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ISBN 978-80-260-6721-4

Copyright© 2014 Czechoslovak Microscopy Society

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Table of content

Plenary lectures

Type of presentation: Plenary

IMC-PL-6095 Light Microscopy at the Nanoscale

Cremer C.1,2
1Superresolution Microscopy, Institute of Molecular Biology (IMB), Mainz, Germany, 2Kirchhoff Institute of Physics (KIP), and Institute of Pharmacy&Molecular Biotechnology (IPMB) University Heidelberg, Heidelberg, Germany
c.cremer@imb-mainz.de

Novel developments in optical technology and photophysics made it possible to radically overcome the diffraction limit (ca. 200 nm laterally, 600 nm along the optical axis) of conventional far-field fluorescence microscopy. Presently, three principal “nanoscopy” families have been established: “Nanoscopy” based on focused laser beams, like 4Pi-, STED- (STimulated Emission Depletion)-, and RESOLFT- (Reversible Saturable OpticaL Fluorescence depletion Transitions) microscopy; nanoscopy based on Structured Illumination Excitation (SIE), like SMI (Structured Modulated Illumination) microscopy, SIM (Structured Illumination Microscopy) and PEM (Patterned Excitation Microscopy); and nanoscopy based on various modes of Localization Microscopy, like PALM (PhotoActivated Localization Microscopy) and FPALM (Fluorescence Photoactivable Localization Microscopy), GSDIM (Ground State Depletion Imaging Microscopy), SPDM Spectral Precision Distance/Spatial Position Determination Microscopy), STORM (STochastic Optical Reconstruction Microscopy) and dSTORM (direct STORM). These and related far-field light microscopy methods have opened an avenue to image nanostructures down to single molecule resolution; they made possible to measure the size of molecule aggregates of few tens of nm diameter and to analyze the spatial distribution of individual molecules with a light optical resolution down to the few nanometer range, corresponding to ca. 1/100 of the exciting wavelength. Application examples obtained by focused, structured, and localization techniques cover a variety of biostructures, such as membrane complexes, neuronal synapses, cellular protein distribution, nuclear nanostructures, as well as the “nanoimaging” of individual viruses and lithographically generated nanostructures. Each of the nanoscopy methods described has its peculiar advantages; as a whole, they provide a tool set of light microscopy approaches to the nanoscale and open a wide range of perspectives in Biology, Medicine and the material sciences. Further improvements are expected to make possible a three-dimensional lightoptical resolution down to the 1 nm scale. The combination with Electron- and X-ray microscopy techniques is anticipated to provide further nanostructural insights.

C. Cremer, Optics far Beyond the Diffraction Limit: From Focused Nanoscopy to Spectrally Assigned Localization Microscopy (2012). In: Springer Handbook of Lasers and Optics, 2nd edition (F. Träger, Edit.), pp. 1351 – 1389.
C. Cremer, B.R. Masters (2013) Resolution enhancement techniques in microscopy. Eur. Phys. J. H 38: 281–344.


Type of presentation: Plenary

IMC-PL-6096 Bioimaging at the nanoscale -- Single-molecule and super-resolution fluorescence microscopy

Zhuang X.1
1Department of Chemistry and Chemical Biology, Department of Physics, Howard Hughes Medical Institute, Harvard University, Cambridge
zhuang@chemistry.harvard.edu

Dissecting the inner workings of a cell requires imaging methods with molecular specificity, single-molecule sensitivity, molecular-scale resolution, and dynamic imaging capability such that molecular interactions inside the cell can be directly visualized. Fluorescence microscopy is a powerful imaging modality for investigating cells largely owning to its molecular specificity and dynamic imaging capability. However, the spatial resolution of light microscopy, classically limited by the diffraction of light to a few hundred nanometers, is substantially larger than typical molecular length scales in cells. Hence many subcellular structures and dynamics cannot be resolved by conventional fluorescence microscopy. We developed a super-resolution fluorescence microscopy method, stochastic optical reconstruction microscopy (STORM), which breaks the diffraction limit. STORM uses single-molecule imaging and photo-switchable fluorescent probes to temporally separate the spatially overlapping images of individual molecules. This approach has allowed multicolor and three-dimensional imaging of living cells with nanometer-scale resolution and enabled discoveries of novel sub-cellular structures. In this talk, I will discuss the general concept, recent technological advances and biological applications of STORM.                                                                                                                                                                                                                                                                                                                                                                               


Type of presentation: Plenary

IMC-PL-6097 Imaging and Spectroscopy of Individual Atoms in Nanostructured Materials

Suenaga K.1
1National Institute of Advanced Industrial Science and Technology (AIST), Tsukuba, Japan
suenaga-kazu@aist.go.jp

It has remained a challenge for scientists to image and discriminate individual atoms since Dalton first proposed distinct properties of atoms in his atomic theory. The requirements to analyze the atomic structures of matter with elemental information are nowadays increasing in importance of cutting-edge research. An elemental analysis down to the single atom limit was first demonstrated with the successful detection of a Gd dopant atom in carbon nano-peapods using a STEM-EELS technique at 100kV [1]. Specimen damage due to the high dose of the incident electron beam, which is required to isolate the signals from individual atoms, is an intrinsic problem for such a highly delicate analysis. Furthermore it is important to prevent the atoms from being kicked out during the observations. In order to reduce the atomic movements and also to enhance the EELS contrast, a lower accelerating voltage is preferred for single atom detection by STEM. Sawada et al. designed a new type of aberration corrector with triple dodecapole elements (the delta system) to reduce the higher-order geometric astigmatism [2, 3, 4], which is critical for the STEM performance operated at low accelerating voltages, i.e., 15 to 60 kV. Here, I demonstrate successful single-atom imaging and spectroscopy in nanostructured materials using STEM together with EELS and/or EDX.

Fig. 1 shows an example for chemical analysis of individual molecules. A carbon nanotube encapsulating two different metallofullerenes (La@C82 and Ce@C82) is examined at 30 kV operating voltage [5]. The annular dark-field (ADF) image clearly shows the molecular structures encapsulated inside the SWNTs (Fig. 1a). Each molecule carries one metal atom, appearing in brighter contrast, inside the cage. We can identify these atoms by simultaneous EELS. Fig. 1b shows two EELS spectra recorded from two atoms. The EELS spectrum shown in green corresponds to the atom indicated by the green arrow. This spectrum is the sum of four spectra, each of which had an acquisition time of 0.05 s. The resulting signal-to-noise ratio is high enough to isolate the La N-edge. On the other hand, the atom indicated by the blue arrow is assigned as Ce. Moreover, its peak position (≈122 eV) fits very well with that for Ce3+ [6]. Though the two edges of La N and Ce N overlap severely, we could identify the elements (La: Z = 57 and Ce: Z = 58) comprising the two encaged atoms. Fig. 1c shows the ADF image, and the elemental mappings for La, Ce, and carbon are shown in Figs. 1d, e, and f, respectively. A further comparison of simultaneous EELS and EDX measurement allows us to directly estimate the fluorescent yield of single atoms [7, 8].
The interrupted periodicities of 2D materials such as graphene, h-BN, and MX2 (dichalcogenides) are of great interest because they govern the physical/chemical properties. Atomic defects, such as a vacancy or impurity/dopant in single-layered materials are investigated with atomic precision. A single-layer of MoS2 exhibits interesting physical properties. The electrical conductivity of MoS2 can be further modulated by doping, such as Re (n-type) and Au (p-type). Typical ADF images of single-layered Re-doped and Au-doped MoS2 are presented in Fig. 2, respectively. The dopants, Re (Z = 75) and Au (Z = 79), appear in brighter contrast in the ADF images than both Mo (Z = 42) and S (Z = 16). Chemical analysis by means of EDX was also done to confirm the doping elements [9]. ADF image in the inset of Fig. 2(left) clearly shows that Re atoms sit at the Mo sites. The Re dopants are well dispersed in MoS2 layers and seldom form clusters on the host material. On the other hand, the Au dopants at similar concentration tend to aggregate on the MoS2 surface (Fig. 2 right). The Au atoms are indeed mobile under the electron beam [9].
A monovacancy in h-BN can be also examined by STEM-EELS (Fig. 3). Core-level spectroscopy on the nitrogen atoms in the vicinity of the boron vacancy was carried out [10]. As shown in Fig. 3a, a monovacancy is induced at the boron site by the knock-on effect, which can be proved by the fact that the darkest contrast appears in the middle of three nitrogen atoms showing brighter contrast. A line spectrum is recorded across the VB (boron monovacancy) along the yellow arrow. From the line spectrum, three typical spectra for the nitrogen K-edge were extracted, with probe positions corresponding to the yellow circles in Fig. 3b. While the first and third spectra are quite similar to the one for the sp2-bonded nitrogen atoms in h-BN with the known * peak at 401 eV, the second spectrum recorded near the VB indeed shows a sharp pre-peak around 392 eV. Although the spectra are rather noisy because of the minimized acquisition time, this pre-peak appears at the same energy level in many different experiments, and arises reproducibly at other VB sites and represents the lowered LUMO state [10].
Identification of individual atoms and examination of their electronic properties in materials are the ultimate goals of all microscopy-based analytical techniques. It is clear that the bonding/electronic states are now accessible from single atoms through EELS fine-structure analysis. For example the radical carbon atoms at the graphene edge have been successfully identified [11, 12, 13]. Moreover the active point defects in 2D materials can now be caught red-handed [14, 15, 16]. I will also show some of the atomic level observations of alloying behavior and phase transition phenomenon of 2D materials, that used to be investigated only by the macroscopic viewpoint [17, 18].                                                                                                                                                                              

References:
[1] K. Suenaga et al., Science, 290 (2000) 2280-2282
[2] H. Sawada, et al., J. Electron Microscopy, 58 (2009) 341-347
[3] H. Sawada, et al., Ultramicroscopy, 110 (2010) 958-961
[4] T. Sasaki, et al., J. Electron Microscopy, 59 (2010) S7-S13
[5] K. Suenaga, Y. Iizumi and T. Okazaki Eur. Phys. J. Appl. Phys., 54, 33508 (2011).
[6] K. Suenaga et al., Nature Chem., 1 (2010). 415-418
[7] K. Suenaga, et al., Nature Photonics, 6 (2012) 545-548
[8] L. Tizei et al., (in this conference)
[9] Y. C. Lin et al., Adv. Mater., (2014). DOI:10.1002/adma.201304985
[10] K. Suenaga, H. Kobayashi, and M. Koshino, Phys. Rev. Lett., 108 (2012). 075501
[11] K. Suenaga and M. Koshino, Nature 468 (2010) 1088-1090
[12] J. Warner et al., Nano Lett., 13 (2013) 4820-4826
[13] J. H. Warner et al., (unpublished)
[14] K. Suenaga et al., Nature Nanotech., 2, 358-360 (2007).
[15] Z. Liu et al., Nature Commun., 2, 213 (2011).
[16] O. Cretu, Y. C. Lin and K. Suenaga, Nano Lett., 14 (2014) 1064-1068
[17] D. O. Dumcenco et al., Nature Commun. 4 (2013) 1351 (5 pages)
[18] Y. C. Lin et al., Nature Nanotech., in press, (2014).


The present research is supported by a JST-CREST and Research Acceleration Programs. All my colleagues in AIST, Y.C. Lin, O. Cretu, L. Tizei, Z. Liu, M. Koshino, Y. Sato, and R. Senga, are gratefully acknowledged. Drs. H. Sawada, T. Sasaki, M. Mukai, Y. Kohno, M. Morishita and K. Kimoto are also acknowledged for the development of dedicated microscopes.

Fig. 1: Single molecular spectroscopy of mixed peapods (La@C82 and Ce@C82) at 30 kV[5]. (a) An ADF image with a rectangle showing where the spectrum image was taken. (b) Two EELS spectra recorded from two metal atoms. The atom indicated by the green arrow is assigned as La, and the other, indicated by the blue arrow, as Ce. (c) ADF image and (d, e, f) chemical maps for La, Ce, and carbon, respectively. Scale bar = 1 nm.

Fig. 2: Detection of single dopant atoms in single-layered MoS2[7]. (Left) An ADF image of Re-doped MoS2. The Re substitution at Mo site (Re@Mo) is pointed by a green arrow. (Right) An ADF images of Au-doped MoS2, where an Au adatom (indicated by a white arrow) located at the hollow-center (Au-HC). Scale bar = 0.3nm.

Fig. 3: Core-level spectroscopy of monovacancy in h-BN layer [8]. (a) ADF image shows a monovacancy in single-layer h-BN. Line spectrum was recorded along the yellow line. (b) Schematic presentation (red: nitrogen, blue: boron) of boron monovacancy. (c) Nitrogen K-edge fine structures extracted from the line-spectrum. Each of the three corresponds approximately to the probe positions marked in (b). A prominent pre-peak in the nitrogen K-edge can be found at 392 eV in the spectrum recorded at position 2, i.e., near the boron vacancy site.

Type of presentation: Plenary

IMC-PL-6098 Electron Tomography for Nanoscale Materials Science

Midgley P.1
1Department of Materials Science and Metallurgy, University of Cambridge, Cambridge, UK
pam33@hermes.cam.ac.uk

The nanoscale complexity of modern materials and devices, be they structural or functional in design, requires high spatial resolution characterisation in all 3 dimensions. The remarkable power and flexibility of a modern TEM makes it the ideal tool for such 3D nanoscale imaging and analysis. Over the past 15 years or so, electron tomography (3D imaging) has grown from a niche technique to one which is now firmly established as an almost routine tool for the 3D study of materials. Early electron tomography used many of the ideas and practices established first in the life sciences. Here, a tilt series of bright-field (BF) images are acquired by rotating the sample about a single axis and recording images every 1-2°. Typically, in the electron microscope, the range of sample tilt is limited either by the sample itself (becoming too thick) or by the objective lens pole piece gap. As such, it is therefore likely that the full tilt range is not accessible, this leads to a ‘missing wedge’ of information and the reconstructions suffer from artefacts, especially an elongation parallel to the optic axis. Dual axis tomography can help in this regard, reducing the missing information through a second tilt series about an axis mutually perpendicular to the first.

For many materials problems, however, BF images may not be ideal and the introduction of STEM HAADF tomography offered materials microscopists an imaging mode that, in many cases, is much more suited for tomography, providing images with greatly reduced diffraction contrast, with a signal that in most cases varies monotonically with thickness (satisfying the projection requirement) and providing compositional contrast through the atomic number (Z) dependence of the high angle (Rutherford-like) scattering [1]. STEM tomography has now become for many the technique of choice for 3D nanoscale imaging in materials science. Fig. 1 shows two examples of STEM HAADF tomography [2,3]. In Fig. 1(a) we see Ge precipitates within an Al-rich matrix revealing a wide variety of morphologies and clear orientation relationships and in (b) the 3D distribution of Ru-Pt catalyst nanoparticles (1-2nm in size) decorating the surface of a mesoporous silica support – here we see only the external surface. The colour of the support indicates the surface curvature with a strong preference of the nanoparticles to be anchored at the ‘saddle points’. STEM tomography (both BF and ADF) has also been developed for the study of defects (especially dislocations) where the reconstruction (or 3D representation) of the dislocation resembles a ‘string’ running through space.

Although determination of the 3D morphology of materials at the nanoscale is now essentially routine, to achieve a high fidelity reconstruction typically ca. 100 images are needed across the tilt range. For many specimens long acquisition times, and thus extended exposure, can lead to damage. However, the number of images required in the tilt series can be reduced if there is prior knowledge about the specimen being reconstructed. Such prior knowledge can be used within a discrete tomography reconstruction (using the physical discreteness of the sample) or, perhaps more generally, within a compressed sensing framework where the primary requirement is that the sample may be described as being ‘sparse’ in some transform domain [4,5]. This sparsity constraint turns out to be very powerful and high fidelity reconstructions can be achieved with remarkably few images (in some cases an order of magnitude reduction compared to conventional reconstructions), see Fig. 2.

Coupling tomography acquisition with analytical techniques, such as EDX and EELS, allows a more detailed exploration of the sample’s chemistry as well as its morphology. Early efforts in this direction included the use of EFTEM, especially using the low loss regime (where loss probability is relatively high), EDX and core-loss EELS. Inevitably, although the speed and efficiency of spectrometers has improved greatly over the past few years, the acquisition time needed for multi-dimensional ‘spectrum-images’ is considerably higher than a conventional image. To keep the total exposure to a reasonable level, fewer images are recorded in the tilt series – ideally perhaps only every 10 or 20°. The reduction in data must be compensated by an increase in prior knowledge to achieve a high fidelity reconstruction; for such ‘multi-dimensional microscopy’ [6], compressed sensing offers an important framework to achieve this. As an example, Fig. 3(a) shows a composite figure illustrating the localised surface plasmon resonances from a silver nanocube. The reconstruction was undertaken on a series of spectrum-images recorded about a single tilt axis every 15°. The 4mm symmetry of the cube-substrate system was imposed at the reconstruction stage as well as a constraint that the reconstruction could be considered as being sparse in a wavelet domain. That constraint provided a reconstruction relatively free of artefact even when using few images [7]. Interpretation of the reconstruction seen in Fig.3(a) can be made within a quasi-static approximation and related back to the potential induced by the electron beam acting back on the electron. Mapping electro-magnetic potentials is also possible using electron holography and coupled with tomography was able to yield 3D reconstructions of the built-in potential near a p-n junction in a silicon device, see Fig. 3(b) [8]. 3D magnetic fields require an enhanced approach using dual axis geometry to determine all the components of the magnetic potential A (or induction B). Here, physical constraints (e.g. in the form of Maxwell’s equations), perhaps again within a compressed sensing framework, could be used to improve a reconstruction of the electro-magnetic potential.

So, what of the future? The electron tomography community is pushing in many directions. Atomic resolution tomography has been demonstrated in some cases: by assuming periodicity within a nanocrystal, the position of an isolated atom in a matrix can be determined and even the location of atoms around a dislocation core. Synergisitc studies with atom probe tomography have been demonstrated already and this may, in the future, develop into an important correlative approach. Mapping physical properties in 3D at the nanoscale continues to be an exciting prospect. Whilst early work showed this to be feasible, further development is needed to improve reconstruction quality. Given the almost ubiquitous use now, in materials-based tomography at least, of iterative techniques (e.g. SIRT, ART, etc) the conventional projection / back-projection approach could evolve into a more model-based one incorporating a detailed description of the beam’s interaction with the sample along its trajectory (e.g. dynamical effects). By iteratively refining an initial model, increased detail about the sample may be obtainable (e.g. strain, fields, induced charges). Lastly, industry requires a robust nanoscale metrology technique that provides reliable 3D measurements of length, porosity, distributions etc. We are still some way in many cases of being able to provide such data with statistical confidence (i.e. error bars!) on our 3D measurements. Improved reconstructions, with fewer artefacts, incorporating prior knowledge, should allow improved and unbiassed segmentation and thus will go a long way to providing a true 3D nanometrology technique.                                                                                                                                                    

References:
[1] P.A. Midgley et al., Chem. Commun. (2001) 907
[2] K. Kaneko et al., Ultramicroscopy 108 (2008) 210
[3] E.P.W. Ward et al., J. Phys. Chem. C 111 (2007) 11501
[4] Z. Saghi et al., Nano Letters 11 (2011) 4666
[5] R. Leary et al., Ultramicroscopy 2013 131 70-91
[6] P.A. Midgley and J.M. Thomas, Angewandte Chemie (2014) DOI: 10.1002/anie.201400625
[7] O. Nicoletti et al., Nature 502 (2013) 80
[8] A. Twitchett-Harrison et al., Nano Letters 7 (2007) 2020


The author thanks his many colleagues, past and present, who have contributed to the work presented here including most recently J.M. Thomas, R. Leary, Z. Saghi, D. Holland, K. Kaneko, S. Hata, O. Nicoletti, F. de la Peña, C. Ducati. PAM acknowledges funding from the European Research Council under FP7/2007-2013 / ERC grant agreement 291522-3DIMAGE.

Fig. 1: (a) 3D reconstruction of Ge precipitates in an Al-rich matrix showing colour-coded to highlight theor different morpholgy [2]; (b) Ru-Pt catalyst particles (red) shown on a colour-coded silica surface where the blue regions indicate positive Gaussian curvature (saddle points) [3].

Fig. 2: (a) Comparison of reconstructions using SIRT and compressed sensing (CS) codes for an iron oxide nanoparticle with a concavity; (b) the apparent concavity volume as a function of projection number [4].

Fig. 3: (a) Colour composite figure indicating five surface plasmon modes on a silver nanoparticle, 100nm in size [7]; (b) Reconstructed electrostatic potential near a p-n junction in a silicon device showing sub-surface carrier depletion [8].

IFSM symposium

Type of presentation: Plenary

IFSM-PL-1670 From Atomic Structure to Properties of Oxides: Applications of Aberration-corrected Transmission Electron Microscopy

Jia C. L.1,2
1Ernst Ruska-Centre for Microscopy and Spectroscopy with Electrons and Peter Grünberg Institute, Forschungszentrum Jülich GmbH, Jülich, Germany, 2International Centre for Dielectric Research, Xi'an Jiaotong University, Xi'an, China
c.jia@fz-juelich.de

Functional oxides provide an important part of the material basis for multifunctional devices as a result of their exceptional range of physical properties. These properties, in turn, depend strongly on the crystal structures, chemical compositions and defect configurations of the materials, which can be characterized on the atomic scale.

In a high-resolution transmission electron microscope equipped with an aberration corrector, the spherical aberration coefficient CS of the objective lens can be tuned to either a positive or a negative value. The use of a negative value of CS combined with an overfocus setting of the objective lens is used in the negative CS imaging (NCSI) technique [1]. Images obtained using the NCSI technique show superior atomic column contrast and intensity than corresponding positive CS images [2], especially for weakly scattering oxygen columns that are in close proximity to strongly scattering cation columns.

Using the NCSI technique, we have investigated the atomic details near 180° domain walls in thin films of PbZr0.2Ti0.8O3 [3,4]. The relative displacements of ions have been measured and on this basis the local polarization across the wall has been calculated. Using this technique we have studied the atomic structure of LaO-TiO2-type interfaces in epitactic LaAlO3/SrTiO3 heterostructures [5]. The prominent result is the oxygen octahedron rotation and the TiO6 octahedra distortion induced by LaAlO3 in SrTiO3 at the interface. The cation-oxygen octahedra represent the prominent structural element of perovskites, which can be modified by distortions, rotations, and particular atomic shifts. Small atomic rearrangements as they are expected to occur at the interfaces between perovskites of different structure can change dramatically the electronic system.

We have recently used the NCSI technique to perform quantitative comparisons between experimental and simulated images on an absolute intensity basis after taking into account the effects of the modulation transfer function of the camera and additional image spread [6]. This absolute intensity matching approach not only allows atomic column positions and defect structures to be determined with picometer precision, but also allows the local chemistry and the three-dimensional morphology of a crystal to be determined on the atomic scale.

Figure 1 shows results obtained from a study of the relationship between the atomic structure and properties of BiFeO3, a room temperature multiferroic material. In the rhombohedrally-distorted perovskite unit cell of BiFeO3 (shown in Fig. 1a), characteristic structural features include relative shifts between the cations and the oxygen anions along the [111] axis and rotations of oxygen octahedra about the [111] axis, which are related to the ferroelectric polarization and the antiferromagnetic properties of the material, respectively. Both the atomic shifts and the rotations of the octahedra can be quantified using the NCSI and ACM techniques and used to understand the electrical and magnetic properties of the material. Figure 1b shows an atomic-resolution image of a 109° domain boundary (thick arrow) between two domains. The use of NCSI conditions and a particular specimen thickness result in the atomic columns appearing bright on a dark background. The domains in the material can then be distinguished by measuring the positions of the atomic columns inside individual unit cells.

In Fig. 1(b), the domain above the boundary is oriented along the [110] direction. The O column positions are shifted upward and downward (Fig. 1c), corresponding to alternating rotations of octahedra. A corresponding off-centre displacement of Fe with respect to the middle point of the line connecting two neighbouring (left and right) O positions is visible and oriented in a downward direction. In this orientation, the [001] component (red arrow) of the [111] polarization vector can be measured and the octahedron rotation can be revealed. Below the boundary (Fig. 1d), the domain is viewed along the [1 ̅10] direction. The octahedron rotation is now not visible due to the overlap of atoms (Fig. 1d). However, the full vector (red arrow) of the atomic column displacement is now revealed. In this way, the polarization of the domain can be determined unambiguously.                                                                                                                       

References:

C.L. Jia, M. Lentzen, K. Urban, Atomic-Resolution Imaging of Oxygen in Perovskite Ceramics. Science 299, 870 (2003).
C.L. Jia L. Houben, A. Thust,and J. Barthel, On the benefit of the negative-spherical-aberration imaging technique for quantitative HRTEM. Ultramicroscopy 110, 500 (2010).
C.L. Jia et al., Atomic-scale study of electric dipoles near charged and uncharged domain walls in ferroelectric films. Nature Mater. 7, 57 (2008).
C.L. Jia et al., Direct observation of continuous electric dipole rotation in flux-closure domains in ferroelectric Pb(Zr,Ti)O3, Science 331, 1421 (2011).
C.L. Jia et al., Oxygen octahedron reconstruction in the SrTiO3/LaAlO3 heterointerfaces, Phys. Rev. B 79, 081405(R) (2009).
C.L. Jia et al., Atomic-scale measurement of structure and chemistry of a single-unit-cell layer of LaAlO3 embedded in SrTiO3. Microsc. Microanal. 19, 310 (2013).


This work was carried out in collaboration with L. Jin, S.B. Mi, K. Urban, A. Thust, J. Barthel, L. Houben, M. Lentzen, D. Hesse and M. Alexe.

Fig. 1: (a) Schematic diagram of the pseudocubic unit cell of BiFeO3. (b) Atomic-resolution image of a 109° domain wall (thick arrow) separating two domains: the domain above the wall and the magnification in (c) correspond to a [110] viewing direction, while the domain below the wall and the magnification in (d) correspond to a [1 ̅10] viewing direction. The red arrows denote the polarization.

Type of presentation: Plenary

IFSM-PL-1780 Some surprises in electron diffraction physics and imaging.

Spence J. C.1
1Physics Department, ASU, Tempe , Az. USA. 85282, and LBNL USA.
spence@asu.edu

The multiple scattering theory on which modern electron microscopy (EM) is based had been fairly well worked out by about 1960, following work by Bethe, Sturkey, Heidenreich, Hirsch, Howie, Whelan, Cowley and Moodie and others. Nevertheless many surprises remained in the ensuing 50 years. For me the most important of these have been i) The finding that multiple energy-loss effects can be removed from EELS spectra, using earlier work on cosmic ray showers. ii) The richness of the "point-projection" geometry, championed by Gabor in 1949. In turn this has produced Ptychography, the theory of STEM lattice imaging for crystals and low-voltage field-emission point-projection imaging. It is remarkable that coherent overlapping convergent beam orders provide a solution to the phase problem, an atomic-resolution "shadow image", Talbot self-imaging, and in-line holography. iii) The discovery of "forbidden" termination reflections and their value for imaging surfaces and sub-surface dislocations and kinks. Their monolayer sensitivity is remarkable. iv) The detection of coherent bremstrahlung tunable X-ray emission lines in STEM EDX. It is remarkable that these lines can be indexed, and are absent when reflections are forbidden by symmetry. v) The explanation for dynamically forbidden reflections, which cancel due to symmetry-related paths for all thickness. vi) The usefulness of electron channelling effects (Alchemi) on EDX for locating foreign atoms in several fields (turbine blades, mineralogy), previously an academic curiosity. vii) The achievement of aberration correction. viii) The success of our TEM CCD camera, whose impact on cryo-em tomography we never anticipated. ix) The surprising sensitivity of low-angle scattering to atomic bonding, with the zero-order scattering the most sensitive quantity known. x) The finding that sufficiently short pulses of radiation can outrun radiation damage, thus breaking the nexus between damage, resolution and particle size if a large number of particles can be packed into a near delta-function pulse. xi) The information extracted from ELS spectra, with its unrivalled spatial resolution.

  The changing agenda of EM over this half-century, from the study of bulk defects such as dislocations, and atomic resolution imaging of interfaces, to nanoscience, cryo-electron and in-situ microscopy (liquid cells, catalysis) has been fascinating to watch. Recent developments - atomic resolution imaging with characteristic X-rays, direct injection detectors, sub-Angstrom resolution, high-resolution imaging in 3D, fast diffraction and imaging - continue to surprise. References in: High Resolution Electron Microscopy (Spence, 4th ed. 2014) and Electron Microdiffraction (Spence & Zuo, 1992).


To many colleagues and friends over half a century in many countries, and to the US funding agencies and Arizona State University.

Type of presentation: Plenary

IFSM-PL-6099 Macromolecules in Motion: Visualization by 4-Dimensional Cryo-Em

Steven A. C.1
1Laboratory of Structural Biology Research, National Institute of Arthritis and Musculoskeletal and Skin Diseases, National Institutes of Health, Bethesda, USA
stevena@mail.nih.gov

Cryo-electron microscopy offers a unique capability to determine the 3-D structures of macromolecular complexes. However, insight into biological activity requires understanding the structural transitions that the complex of interest undergoes. It is not possible, even in principle, to visualize the same molecule in successive states as this would involve the prohibitively difficult task of thawing the specimen, inducing the conformational change, re-vitrifying and re-locating the molecule. However, dynamics may be addressed by a statistical approach in which classification techniques are applied to data sets imaging conformationally mixed populations. Then, provided that there is a basis for ordering the various conformers in a temporal sequence, the reaction dynamics of the complex may be described and movies made.

The process of virus maturation is amenable this approach. With many viruses, the precursor particle undergoes radical structural changes as it matures into an infectious virion. We have investigated the maturation of bacteriophage HK97 capsid, an icosahedrally symmetric shell composed of 420 protein subunits, which expands from 45 nm to 55 nm and angularizes as it matures. These changes in morphology reflect large rotations of the protein subunits and local remodeling (1), and the pathway proceeds via three metastable intermediates (2). The capsid of herpes simplex virus, an animal virus, follows a similar pathway, which may be traced to a capsid protein domain similar in structure to that of HK97, but it passes through many more intermediates (3). Recently, we have found that bacteriophage phi6, which has a RNA genome rather than a DNA genome and an entirely different capsid protein fold from HK97, also undergoes massive subunit rotations and matures via two intermediates (4, 5) – Figure 1.

For this approach to visualization of conformational dynamics, several conditions must be met (6). The differences between states must be large if differences in the images that arise from viewing geometry are to be separated from real structural differences. The number of distinct conformational states must be relatively small. To establish a time-line, one must be able to induce the reaction of interest on a time-scale of seconds to hours. Notwithstanding, recent technical advances in automated collection of large data sets, the improved resolution and signal-to-noise ratio of direct detection cameras, and sophisticated classification techniques promise to expand the range of applicability.                                                                                                                                            

References

1. J.F. Conway et al. Science 292: 744-748 (2001)
2. R. Lata et al., Cell 100: 253-263 (2000)
3. J.B. Heymann et al., Nature Struct. Biol. 10, 334-341 (2003)
4. D. Nemecek et al., J. Mol. Biol. 414, 260-271 (2011)
5. D. Nemecek et al., Structure 21, 1384-83 (2013)
6. J.B. Heymann et al., J. Struct. Biol. 147, 291-301 (2004)


I thank many colleagues, particularly Drs N. Cheng, J.F. Conway, J.B. Heymann and D. Nemecek. This research has been supported by the intramural research program of NIAMS/NIH.

Fig. 1: Maturation dynamics of bacteriophage phi6 capsid visualized by 4-dimensional cryo-EM. The pathway progresses through four states: the initially assembled procapsid with its deeply indented facets; two expansion intermediates with near-planar facets; and the spherical mature nucleocapsid. The capsid is built from 60 P1A/P1B dimers, where P1A and P1B are chemically identical but conformationally distinct (non-equivalent) protein subunits. The top row shows renderings of the outer surfaces viewed along a 5-fold symmetry axis. The middle row shows models of the respective capsids consisting of a pentamer of P1A subunits (blue, green) and surrounding P1B subunits (red, yellow), viewed from above. The bottom row shows slabs through a portion of the structures, passing through one vertex (P1A’s blue, P1B’s blue). The procapsid-to-intermediate 1 transition is achieved by rotations of P1B subunits about the line connecting two 3-fold icosahedral axes. Further expansion to intermediate 2 is achieved by outward movement of the P1A subunits. The final step to nucleocapsid involves primarily local changes affecting the P1A subunits. Adapted from reference (5), where movies of the transition are available, courtesy of J.B. Heymann. Scale bars: 100 Å.

Type of presentation: Plenary

IFSM-PL-6100 From the Prague Spring to a Spring in Electron Microscopy

Křivánek O. L.1
1Nion Co., Kirkland and Dept of Physics, ASU, Tempe, USA
krivanek@nion.com

Prague is my native city: I was born in Praha-Bubeneč, on the plateau behind the Prague Castle. I grew up in the era of the Czechoslovak Socialist Republic (ČSSR), when the Soviet Union and its satellites prided themselves on their space exploits and their education systems. Among the special efforts they made were competitions for talented youngsters in mathematics and physics, and I used to enjoy those. In my senior high school year, I qualified for the national round in both math and physics, and in physics I was invited onto the national team of three that represented Czechoslovakia at the 2nd International Physics Olympiad, held in Budapest in June 1968. Back then only the Soviet Union and its satellites participated – Western Europe, USA and other countries joined the Physics Olympiad later. Our team did well: we got a joint second place with the Hungarians and the East Germans, with the Soviets winning the first place. I have since then had the pleasure of working with one other former International Physics Olympian – Niklas Dellby, my partner at Nion.

That same summer I took the entrance exam to Charles University in Prague, to study physics. I passed and promptly took off on a trip I had planned: a vacation in the south of France, followed by a stay in London where I was planning to work in a summer job while improving my English. 1968 was the year of the famous Prague Spring, when “socialism with a human face,” which included many democratic measures, was introduced by a group of reformers led by Alexander Dubček (Fig. 1a), much to the displeasure of the old guard in the Kremlin. As I was boarding the train to France, my father told me: “If the Soviets invade, stay in the West.” I had not been following the political situation very closely, so this came as a surprise instruction to me. The Soviets invaded 4 weeks later (Fig. 1b), while my whole family happened to be in the West: my parents on vacation in Austria, my sister working in a summer job in France and me working as an office helper in London. We got together on the phone, and decided that none of us would go back to Prague, at least not for the time being. (See [1] for an especially lucid account of the Prague Spring.)

People were very sympathetic to citizens of a small country invaded by Soviet tanks, and the British National Union of Students had a special place in its London office for notices of available openings for prospective Czech and Slovak students. I was checking it daily while working in a new job, as a carpenter. Around the end of September, a small notice appeared, announcing that the University of Leeds was going to offer up to 5 scholarships to qualified Czechoslovak students. I called them up and caught the train to Leeds soon thereafter. There was an entrance interview during which it became clear that I knew my physics all right, and also that 3 years of high school English and a vocabulary of perhaps 3000 English words were not nearly enough for me to slot painlessly into the British university system.

Leeds took a chance on me, and at first they must have wondered how it would turn out. In my first year I got a First in math – understanding equations did not require much English – but only a Pass in physics, in which there were long textbook passages that I studied laboriously, with a dictionary in hand. I did better in later years, graduating with a First, at the top of my class. I was then accepted to do a physics Ph.D. in Cambridge, with Archie Howie as my inspiring supervisor. In my first year, our lab was not far from Ellis Cosslett’s, after whom the award I received is named, and who has been one of my heroes in electron microscopy, especially after I came to appreciate the pioneering nature of much of the work of his group.
I greatly enjoyed my time in Cambridge, both inside and outside the Physics Department. I learned a lot, made many friends, and made good use of Cambridge’s excellent extra-curricular facilities. I raced for Cambridge against Oxford in skiing and won the special and parallel slaloms at the 1975 Oxford-Cambridge ski race, in the Italian Dolomites. The 8-man boat crew I joined the previous spring (Fig. 1c) did three bumps and an overbump in the Cambridge May Races, and by Cambridge tradition, we got to keep our oars as souvenirs.

After Cambridge, I worked at Kyoto University for 3 months, and did post-docs at Bell Labs and UC Berkeley, where I joined the group of Gareth Thomas in the Materials Science Department. Being in Materials Science made me feel that I had to make a choice: I could concentrate on the materials we were studying and become a materials scientist, or on the instruments and techniques we were using and remain a physicist. I had done a little instrument design work and liked it, so the second option seemed more attractive. The technique I thought was especially interesting was a new one (to me) called Electron Energy Loss Spectroscopy (EELS). I got my first taste of it at the 1978 Cornell workshop, where I met people who became lifelong friends, such as Phil Batson, Christian Colliex, Ray Egerton, and Mike Isaacson. One was expected to build one’s own spectrometer in those days – there were no commercial models. When I got back to Berkeley, I climbed the stairs to Professor Thomas’s office and said: “I think I should build an energy loss spectrometer. It will allow us to study oxygen concentrations at grain boundaries in nitrogen ceramics.” – a subject the group was focusing on. Gareth asked just one question: “How much will it cost? ”, I replied “about $10k”, and I had my first OK to build a major instrument.

The spectrometer came together quickly and produced good results (Fig. 2a). In the summer and autumn of 1979, I was showing the results at various conferences. At one of them, at NBS in Washington, Nancy Tighe came up to me and said: “I think your spectrometer would interest Peter Swann of Gatan. You should give him a call.” This started my fruitful collaboration with Peter, from whom I learned on many fronts. Peter passed away in the summer of 2013, and many of us miss him very much.

Over the next year, Peter Swann, Joe Lebiedzik and I, with input from Mike Scheinfein, designed and built a second-generation serial EEL spectrometer. I also started in a new job, as Associate Director of the NSF-funded HREM facility at Arizona State University. With my collaborators at ASU, we applied the spectrometer to many interesting problems, and put together the EELS Atlas [2] that is used to this day. ASU was a great place to work. There were many good instruments, several leading researchers in electron microscopy, and stimulating annual schools and workshops (Fig. 1d), whose organization was my responsibility.
The pull of Gatan, however, proved irresistible when Peter moved its R&D facility from Pittsburgh to California, and in 1985 I became Director of Research at Gatan. A very productive period followed, during which I had the privilege of working with many talented researchers and designers: Dan Bui, Niklas Dellby, Garry Fan, Stuart Friedmann, Sander Gubbens (the current President of Gatan), Robert Keeney, Bernd Kraus, Mike Kundmann, Mike Leber, Chris Meyer, Paul Mooney, Ming Pan, Nils Swann, Peter Swann, Marcel Tence and Jacob Wilbrink, among others. We introduced a number of innovative products, including parallel EELS, imaging filters, CCD cameras, scanned image acquisition systems and DigitalMicrograph software. Gatan grew nearly 10x in size during this time, and I learned that developing instruments commercially can be a great way to fund instrumentation research, especially when working with like-minded researchers and lean and understanding administrations.
The next big change in my scientific life came when Peter decided to retire in 1992, and “professional managers” took over at Gatan. My freedom to do interesting projects was greatly restricted, and I started to look around. It had been clear to me since about 1990 that having managed to correct the second order aberrations of the quadrupole optics of imaging filters, I had a good chance of correcting third order aberrations – a classic problem in electron optics since Scherzer’s work on the subject in the 1930s and 40s. It seemed too speculative a project for Gatan, however, and so I explored doing it elsewhere. My first try for corrector funding was a chat with Uli Dahmen, the Berkeley NCEM director, who consulted with Bob Gottschall, his manager at DOE. Bob’s answer was apparently “over my dead body.” He had gotten burned funding Crewe’s corrector attempts, which never led to a working instrument.

I was more successful persuading Mick Brown of my Alma Mater, Cambridge University, who had a spare VG cold field emission (CFE) scanning transmission electron microscope (STEM), that we should jointly build a corrector for it. We applied for funding to the British Royal Society and secured the maximum allowed amount from the Paul Instrument Fund: £80k. I then moved to Cambridge with my family for two wonderful years. Niklas Dellby and others joined the project, and we had a working proof-of-principle STEM corrector about 2 years later [3], the same summer (1997) as the Heidelberg-Julich CTEM corrector started working.
The 100 kV VG STEM we built our corrector for was older than a research student who joined the project (Andy Lupini), and it had poor aberration coefficients (Cs~Cc~3.5 mm). We improved its resolution, but we did not beat any resolution records relative to the best uncorrected instruments. (The same was true for the Heidelberg effort – 1 MV microscopes were then giving higher resolution than their corrected 200 kV CTEM.) However, a corrector of an improved design we built for Phil Batson’s extensively modified VG at IBM Yorktown Heights achieved a double distinction: it led to the first STEM able to focus an electron probe to <1 Å diameter [4], and it was, as far as I know, the first commercial corrector (delivered in June 2000).

Aberration correction soon became a “hit”, with CEOS GmbH supplying correctors to all the regular manufacturers of electron microscopes, and the company Niklas Dellby and I started near Seattle, Nion, concentrating on correctors for CFE STEM and going it alone. Our idea was a somewhat crazy one: that we could extend our prowess in correctors by designing a whole new electron microscope, and that we would do it better than the regular manufacturers. Not many thought that we would succeed. But there were early believers to whom we owe a great deal, such as John Silcox, Andrew Bleloch, Steve Pennycook and Christian Colliex. Benchmarks established subsequently by Nion for resolution, stability, probe current, ultra-high vacuum, freedom from contamination and powerful software [5,6] have persuaded many others. 

Nion’s very capable team - Niklas Dellby, Neil Bacon, George Corbin, Peggy Cramer, Zeno Dellby, Russ Hayner, Petr Hrncirik, Tracy Lovejoy, Chris Meyer, Savath Phoungphidok, Michael Sarahan, Gwyn Skone, Zoltan Szilagyi, Janet Willis, Tad Yoo and myself for now, and growing, has done some amazing things. We first delivered 10 aberration correctors for VGs, then moved onto making whole electron microscopes. Currently we’re manufacturing Nion microscopes #10-13, and the interest in our instruments is on the rise. Building the instruments has been made easier by the close collaboration we enjoy with Czech Republic’s Delong Instruments, especially Vladimír Kolařík and Petr Homolka. Nion’s progress has also been helped by two simple facts: ordering an electron microscope from a small company is a gutsy thing to do, and gutsy scientists tend to be first-rate. (Figs 2b-3a,b) and references [7-13] show some of the revolutionary results they and their collaborators have obtained with Nion microscopes.

Aberration correction has ushered in an era of electron microscopy in which we can see the structure, composition and bonding of materials better than ever before. It amounts to a new spring in electron microscopy, best captured by the words of David Cockayne: “it is as though a veil of fog has lifted from our samples.” It is about to get better still, because of an exciting new development: studying energy losses with sub-20 meV energy resolution and sub-nm spatial resolution. This has been made possible by Nion’s new monochromator [14], which has been the subject of two separate talks at this congress [15], and which promises to make vibrational excitations in materials (phonons) readily observable (Fig. 4), at a high spatial resolution. It will probably also allow hydrogen to be mapped in the electron microscope, using energy losses that accompany high-angle scattering of fast electrons by hydrogen nuclei. 

My scientific instrumentation journey began with EELS and progressed onto aberration correction and high resolution STEM. It has now come back to EELS, with an energy resolution about 100x better than on my first try. My life’s journey began in Prague, and Prague is where this congress has been held. Both journeys are reminiscent of the famous lines by T.S. Eliot [16]:                                                                                

We shall not cease from exploration,
And the end of all our exploring
Will be to arrive where we started
And know the place for the first time.                                                                                                       

So let us celebrate exploration (also known as research) and knowing where we came from. And also congresses such as IMC, which enrich our knowledge of our field, and of ourselves.                                                                                                                                                        

[1] A. Levy, Rowboat to Prague (ISBN 0-670-60920-X), reprinted as So Many Heroes (ISBN 978-0933256125). See also http://en.wikipedia.org/wiki/Alan_Levy
[2] C.C. Ahn and O.L. Krivanek, EELS Atlas (1983) Gatan and the ASU HREM facility.
[3] O.L. Krivanek et al., Proc. EMAG 1997, IOP Conf. Ser. No 153 (J. Rodenburg, ed.) 35-40.
[4] P.E. Batson, N. Dellby and O.L. Krivanek, Nature 418 (2002) 617-620.
[5] O.L. Krivanek, et al., Ultramicroscopy 108 (2008) 179-195.
[6] N. Dellby et al., The European Physical Journal Applied Physics 54 (2011) 33505 (11 pages).
[7] D.A. Muller et al., Science 319 (2008) 1073-1076.
[8] O.L. Krivanek et al., Nature 464 (2010) 571-574.
[9] T.C. Lovejoy et al., Appl. Phys. Letts 100 (2012) 154101 to 154101-4.
[10] P.Y. Huang et al., Nano Letters 12 (2012) 1081-1086.
[11] W. Zhou et al., Microscopy and Microanalysis 18 (2012) 1342-1354.
[12] Q.M. Ramasse et al., Nano Letters 13 (2013), 4989–4995.
[13] J. Lee at al., Nature communications 4 (2013) 1650.
[14] O.L. Krivanek et al., Microscopy 62 (2013) 3-21.
[15] N. Dellby et al., these proceedings and O.L. Krivanek et al., these proceedings.
[16] T.S. Elliot, Four Quartets (1943) ISBN 978-0156332255.


Fig. 1: a) Alexander Dubček, who led the Prague Spring. b) Soviets tanks rolling through Prague’s Wenceslas square. c) Rowing on the river Cam. I am in seat #5 (bow = 1), holding the oar that does not quite match the others. d) 1981 ASU meeting. Front row: Mike Isaacson, Alan Craven, John Spence, John Venables, Albert Crewe, John Cowley, Bernard Jouffrey, Ian Wardell, Ondrej Krivanek, Colin Humphreys. Spot Ray Carpenter, Mark Disko, Murray Gibson, Sumio Iijima, Kazuo Ishizuka, Masashi Iwatsuki, Charlie Lyman, Peggy Mochel, Steve Pennycook, Jing Zhu and others in the photo.

Fig. 2: a) EEL spectrum of BaTiO3 recorded with the serial EEL spectrometer I built at Berkeley, at about 2 eV resolution. b) Reversible atomic motion in monolayer graphene: one of the 6 substitutional Si atoms moves right, left, right. Nion UltraSTEM100, 60 kV, 6 s between frames. (Ref. [13])

Fig. 3: a) EEL L2,3 spectrum from a single Si atom replacing a C atom in graphene (line) and theoretical fits (solid spectra). The right fit allowed the Si atom to “pop out” 0.65 Å from the graphene plane (inset) and gave better agreement. (Ref. [12]) b) Results from Nion microscopes. Nature vol. 464 (2010) issue 7288 cover: image of BN monolayer with impurities by Matt Chisholm, processing by the author and Tim Pennycook. Angewandte Chemie vol. 50 (2011) issue 43 cover: image of MoS2 by Quentin Ramasse. Nature Materials vol. 11 (2012) issue 10 cover: EELS elemental map by Julia Mundy.

Fig. 4: 60 kV results from the Nion High Energy Resolution Monochromated EELS-STEM (HERMES). a) Spectra obtained with the slit out and in, slit-in acquisition time 0.25 s, courtesy Niklas Dellby (Nion) and Philip Batson (Rutgers U.). b) spectrum from titanium hydride (acq. time 10 s), courtesy Peter Crozier and Jiangtao Zhu (ASU), and Tracy Lovejoy (Nion).

IT-1. Electron optics and optical elements

Type of presentation: Invited

IT-1-IN-2012 Advances in electron vortex experiments in the TEM

Verbeeck J.1, Béché A.1, Clark L.1, Guzzinati G.1, Juchtmans R.1, Van Boxem R.1, Van Tendeloo G.1
1EMAT, University of Antwerp, Groenenborgerlaan 171, 2020 Antwerp, Belgium
jo.verbeeck@ua.ac.be

Electrons in a transmission electron microscope are successfully described by linear combinations of plane waves. The sample and the magnetic lenses deform the wavefronts of these waves in a way that transfers information from the sample onto a detection plane. Alternatives to this plane wave basis are however possible and especially cylindrical harmonics are an interesting option. The plane waves are replaced by waves which have a typical azimuthal phase factor exp(i m φ) with φ the angle in the plane perpendicular to the optical axis and m the so-called topological charge. Such waves are orbital angular momentum (OAM) eigenstates in the sense that a normalized cylindrical wave carries exactly mħ angular momentum around the cylinder axis. These waves are often referred to as vortex waves and they attract considerable attention in many different fields of physics including optics, acoustics, radio communication and quantum information [1].
Electron vortices were theoretically predicted to possess also a quantised magnetic moment mµB on top of the common OAM of mħ due to their electrostatic charge [2]. It took untill 2010 before pure electron vortex modes were demonstrated in a transmission electron microscope [3,4]. Since then, many different ways of producing these waves have followed (see e.g. fig.1), each with different advantages and disadvantages. The latest addition, sketched in fig.2, is the production of electron vortex waves making use of a thin single domain magnetic needle approximating a magnetic monopole [5]. This method holds great promise as it offers pure vortex modes at full beam current. Apart from producing single vortex modes, we also focused on the detection of the OAM in an arbitrary wave. Several methods are possible and will be discussed. In terms of the interaction with a sample we observed magnetic dependence in EELS spectra of ferromagnetic samples relating to electron magnetic chiral dichroism and its X-ray counterpart X-ray magnetic chiral dichroism. On top of this, we will discuss the use of vortex beams in elastic diffraction and the transfer of angular momentum to rotate nanoparticles.

References

[1] J. F. Nye and M. V. Berry., Proc. of the R. Soc. of London. A. 336/1605 (1974) 165.

[2] K. Bliokh et al., Phys. Rev. Lett. 99 (2007) 190404.

[3] M. Uchida and A. Tonomura., Nature, 464/7289 (2010) 737.

[4] J. Verbeeck et al., Nature 467/7313 (2010) 301.

[5] A. Béché et al., Nat. Phys.10/1 (2013) 26.


This work was financially supported by the European Union: ERC grant 246791 COUNTATOMS, ERC Starting Grant 278510 VORTEX, Integrated Infrastructure Initiative grant 312483-ESTEEM2.

Fig. 1: Using a probe aberration corrector and an annular condensor aperture, an approximation to an azimuthally varying phase plate can be obtained. This is one of the alternative ways to create electron vortex beams.

Fig. 2: Sketch of the creation of an electron vortex beam by impinging a plane wave to the end of a long bar magnet approximating a magnetic monopole.

Type of presentation: Invited

IT-1-IN-2597 Electron Optics for High-brightness High-beam-current Column Design --- extracting micro-amperes from point cathodes ---

Fujita S.1, Takebe M.1, Wells T.2, El-Gomati M. M.3, Shimoyama H.4
1SHIMADZU Corporation, Kyoto, Japan, 2York Probe Sources Ltd., York, United Kingdom, 3University of York, York, United Kingdom, 4Meijo University, Nagoya, Japan
fujita@shimadzu.co.jp

A principal goal of designing electron probe forming system is to focus desired beam current into as small a spot on the target as possible. Increasing demand for analytical measurements is making desired beam current higher than ever (Ib>10nA). This article describes strategies to design high-brightness high-beam-current electron optical columns.

Figure 1 shows Probe Property that relates the beam current Ib to the probe size d. The dotted curve assumes as electron source a thermionic gun while the dashed curve is for conventional ZrO/W (100) Schottky emitter (SE) gun system. A higher brightness of the latter makes the probe size substantially smaller in the middle beam current regime. However, the probe blurs fast once the current exceeds a certain threshold.

Probe Property is limited by three different mechanisms with increasing beam current order:

Beam Current Regime   “low”                “middle”          “high”
Limiting Mechanisms     wavelength       brightness       angular intensity
                                        chromatic(OL)   spherical(OL)   spherical(Gun)

Attempts were made to improve “high” beam current performance by increasing the source angular intensity and suppressing the gun spherical aberration. Extended Paraxial Trajectory Method is used to analyze electron rays starting from cathode surface with large slopes [1]. The emission characteristic of SE gun is then given by optical parameters familiar in lens designs.

The first strategy is to adopt an emitter whose tip radius is significantly larger [2]. Figure 2 compares a scaled-up emitter (giant SE = GSE) with a conventional SE. The tip size effect is reflected in “electron gun focal length,” f. The angular intensity is given by JΩ = f2*js where js is the cathode current density. Since the focal length is roughly proportional to the tip size, a large tip leads to an improved angular intensity.

The second strategy is to immerse the emitter in the condenser lens field, which is known to result in a suppressed spherical aberration.

Figure 3 compares the source emittance diagrams of conventional SE and GSE. GSE’s wide and less-distorted diagram demonstrates its high-beam-current capability. It is expected GSE’s improved emittance extends the “middle” beam current regime to Ib ~ 1μA (see Fig.1).

A test column was constructed by combining the GSE gun with an objective lens designed for efficient X-ray detection. SEM image observations at Vacc = 10kV over beam current range 100pA <Ib< 3μA confirmed semi-quantitatively the predicted probe property given in Fig.1.

[1] S.Fujita, M.Takebe, W.Ushio and H.Shimoyama, J.Electron Microsc. 59, 3 (2010).
[2] S.Fujita, T.R.C.Wells, W.Ushio, H.Sato, and M.M.El-Gomati, J.Microsc. 239, 215 (2010).


The authors thank Shimadzu Corporation for the support of this work as well as for the permission of the publication.

Fig. 1: Probe Properties with three different electron sources. Expected probe size is plotted against the beam current.

Fig. 2: Comparison of tip geometries of conventional SE and scaled-up emitter (GSE). Angular intensity can be increased by a larger tip radius.

Fig. 3: Source emittance diagrams of conventional SE and GSE. GSE’s wide and less-distorted emittance demonstrates an improved high-beam-current capability of the emitter.

Type of presentation: Oral

IT-1-O-1664 Magnetic-Field-Superimposed Cold Field Emission Gun for 1.2-MV Transmission Electron Microscope

Kasuya K.1, Kawasaki T.1, Moriya N.1, Arai M.1, Furutsu T.1
1Central Research Laboratory, Hitachi, Ltd., Akanuma 2520, Hatoyama, Saitama, 350-0395, Japan
keigo.kasuya.bp@hitachi.com

     A magnetic-field-superimposed cold field emission gun (M-FEG) was developed for a 1.2-MV transmission electron microscope (TEM)[1]. This microscope is intended to have a point resolution of 40 pm and to take atomic-scale three-dimensional images by electron holography.

     Figure 1 shows the cross section of the developed M-FEG. The gun is designed to have a high brightness and stable emission current. The gun is equipped with a pre-accelerator magnetic lens placed close to the emitter [2]. The superimposed magnetic field causes the emitted electrons to converge so that the aberration-caused blurring with subsequent electrostatic lenses is minimized. As a result, the inherent high brightness of the cold field emitter can be obtained. The chambers of the gun are differentially evacuated with three non-evaporative getter (NEG) pumps and four ion pumps. The pressure of the first chamber, where the emitter is placed, was 3×10-10 Pa. This small pressure stabilizes time variations of the emission current [3].

     Figure 2 shows the measured time variations of the probe and total currents. After performing flashing of the emitter, the initial probe current of 1 nA was obtained at the total current of 1 µA. The probe current stayed almost constant for more than 10 hours during the initial period of the measurement. The 90% decrease time, at which the current falls to 90% of the initial value, was prolonged to 900 min in comparison with 3 min in a previous gun at 5×10-8 Pa [4]. The variation in the probe current over the course of the initial 8 hours was 5.2%.

     Another advantage of the pressure reduction is the increase in probe current. It increased two times higher than that of the conventional field emission gun operating at 10-8 Pa. This reason can be explained by the fact that the clean emitter surface has higher probe current density than the adsorbed surface. The gun provided large probe currents ranging from 1 to 170 nA for total currents ranging from 1 to 300 µA.

     The resulting current characteristics ensure that the 1.2-MV TEM will have fine resolution with a high S/N ratio. The illumination system of the microscope is discussed by Kawasaki in this conference.

[1] K. Kasuya et al., submitted to J. Vac. Sci. Technol. B.

[2] M. Troyon, Optik 57, 401 (1980).
[3] K. Kasuya et al., J. Vac. Sci. Technol. B 28, L55 (2010).
[4] T. Kawasaki et al., J. Elec. Microsc. 49, 711 (2000).

 


This research was supported by the Japan Society for the Promotion of Science through the FIRST Program, initiated by the Council for Science and Technology Policy.

Fig. 1: Cross-section of the developed magnetic-field-superimposed cold field emission gun (M-FEG). The pressure of the first chamber was 3 ×10-10 Pa.

Fig. 2: Measured time variations of probe and total currents. The 90% decrease time of the probe current was 900 min. The variation in the probe current over the initial 8 hours was 5.2%.

Type of presentation: Oral

IT-1-O-1982 Challenges in Phase Plate Development and Applications

Sader K.1, Buijsse B.1, van Duinen G.1, Danev R.2
1FEI, Achtseweg Noord 5, 5651 GG Eindhoven, The Netherlands, 2Max Planck Institute of Biochemistry, Am Klopferspitz 18, 82152 Martinsried, Germany
kasim.sader@fei.com

While there have been attempts to implement phase plates in transmission electron microscopes (TEMs) over a long period of time, a publication by Danev and Nagayama [1] renewed interest that functional phase plates could be produced. In particular in life sciences, the development of thin film vitrification techniques has enabled the examination of unstained macromolecules and thin cells in the electron microscope, but also created the need for phase contrast. Conventionally, contrast at low spatial resolutions has been generated by using a strong defocus, but with the added consequence of introducing oscillations in the contrast transfer function. A phase plate allows one to work in-focus, with a large increase in the contrast at low spatial resolutions.

Many types of phase plates have been proposed, but the most widespread implementation has been the original thin-film Zernike phase plate. This type of phase plate has shown practical performance, especially in life science applications. The most widely tested film type is amorphous carbon, but these suffer from aging problems, making frequent exchanges of the phase plate necessary. Alternatives to conventional amorphous carbon have been investigated and silicon-based films show promise in terms of longevity.

In close collaboration with the Max Planck Institute of Biochemistry in Martinsried, FEI have developed a new type of phase plate with properties that make it very suitable for implementing it as a user friendly device in our TEMs. It produces high-contrast images, providing excellent contrast transfer in the low resolution range which is particularly relevant for cryo-electron tomography and may provide benefits for single particle analysis in the case of small and heterogeneous particles. No fringing effects around high-contrast features are observed and CTF oscillations can be avoided up to better than 10Å while maintaining contrast transfer at low spatial frequencies. Transmission losses by the phase plate are very modest. Moreover, the phase plate shows consistent performance for at least half a year of usage.

To facilitate routine phase plate usage we have added extra alignments and control panels to the microscope software. In particular, accurate adjustment of beam deflection pivot points is included to ensure a stable beam position at the plane of the phase plate. Also, software has been developed to easily navigate the phase plate in the back focal plane. We are developing detailed phase plate workflows for our applications software that will provide a seamless integration of the phase plate in the (automated) applications. In this talk a selection of results will be shown from cryo electron tomography.

[1] R. Danev, K. Nagayama, Ultramicroscopy 88, 243-252 (2001)


Fig. 1: Tecnai F20 results from cryo electron tomography on doxorubicin using conventional TEM at 4 μm defocus. Experimental conditions: total dose of 85 e-/Å2, tilt range +/-60º. 

Fig. 2: Tecnai F20 results from cryo electron tomography on doxorubicin using phase plate TEM at 0.5 μm defocus. Experimental conditions: total dose of 85 e-/Å2, tilt range +/-60º.

Type of presentation: Oral

IT-1-O-2071 Sculpturing the electron wave function using nanoscale phase masks

Shiloh R.1, Lereah Y.1, Lilach Y.1, Arie A.1
1Department of Physical Electronics, Fleischman Faculty of Engineering, Tel Aviv University, Tel Aviv, Israel
royshilo@post.tau.ac.il

Electron beams are extensively used in lithography, microscopy, material studies and electronic chip inspection. Today, beams are mainly shaped using magnetic or electric forces, enabling only simple shaping tasks such as focusing or scanning. Recently, binary amplitude gratings achieved complex shapes. These, however, generate multiple diffraction orders, hence the desired shape, appearing only in one order, retains little of the beam energy. Here we demonstrate a method in electron-optics for arbitrarily shaping electron beams into a single desired shape, by precise patterning of a thin-membrane. It is conceptually similar to shaping light beams using refractive or diffractive glass elements such as lenses or holograms - rather than applying electromagnetic forces, the beam is controlled by spatially modulating its wavefront. Our method allows for nearly-maximal energy transference to the designed shape, and may avoid physical damage and charging effects that are the scorn of commonly-used (e.g. Zernike and Hilbert) phase-plates. The experimental demonstrations presented here – two solutions to the free-space wave equation: on-axis Hermite-Gauss and Laguerre-Gauss (vortex) beams, and computer-generated holograms – are a first example of nearly-arbitrary manipulation of electron beams. Our results herald exciting prospects for microscopic material studies, research in electron-matter interaction, enables electron lithography with fixed sample and beam and high resolution electronic chip inspection by structured electron illumination.


The work was supported by the Israel Science Foundation, grant no. 1310/13 and the German-Israeli Project cooperation.

Fig. 1: On-axis generation free-space modes: images taken at different effective distances near the diffraction plane. (A) Unmodulated beam passing through the membrane, (B) Hermite-Gauss11-like, (C) Laguerre-Gauss01-like (vortex), (D) Bragg diffraction pattern used as metric, (E) Bragg grating, (F) HG11-generating mask, (G) vortex-generating mask.

Fig. 2: On-axis holograms: (A) “TAU” hologram produced by the mask in (B); inset: magnification showing ~60nm holes composing the pixels. (C) Electrons orbiting a nucleus hologram produced by the mask in (D); inset: magnification showing the centre of the mask. Note: contrast and brightness levels in (C) were altered for visibility.

Type of presentation: Oral

IT-1-O-2914 Tuning and Operation of a sub-20 meV Monochromator

Dellby N.1, Lovejoy T. C.1, Křivánek L. O.1
1Nion Co., 1102 Eighth St., Kirkland, WA 98033, USA
dellby@nion.com

When aiming for simultaneous high energy resolution and high spatial resolution in a monochromated scanning transmission electron microscope (STEM), three locations in the microscope are critical:
1) the monochromator’s (MC’s) energy-selecting slit, where the pass-band of energies admitted into the rest of the column is determined,
2) the sample, where the tuning determines the spatial resolution, and
3) the detector of the electron energy loss spectrometer (EELS).
To optimize the performance of the entire system, aberrations in all three locations must be accurately and repeatably tuned, so as to produce the smallest possible beam crossover at each place. In typical operation, all three crossovers are images of the field emission source, and upstream crossovers are re-imaged in subsequent stages. A mistuned monochromator can be largely compensated by a pre-sample aberration corrector that is mistuned in the opposite direction, or by a mistuned EELS.
The ideal method for monochromator tuning should therefore measure the actual aberrations at the plane of the energy selecting slit and not be affected by post-monochromator optics. We use a variation of the method developed by Foucault[1]: we image the far-field shadow of the energy-selecting slit near which the beam crossover is formed.
With a monochromatic beam coming into the monochromator, the aberrations would be tuned when the far-field image of the slit fades out uniformly as the slit is closed up. Non-zero focus and astigmatism would produce a stripe across the image of the beam-defining aperture, and one would focus and stigmate to make the stripe wider until it fills the aperture.
In practice, however, the slit is illuminated with an energy-dispersed beam some 300 meV wide, i.e. about 20 times larger than the energy width of our usual monochromated beam. This means that electrons with different incoming energies fill different parts of the aperture with stripes of different energies (Fig. 1), and the total beam after the slit is an incoherent superposition of a distribution of slit positions.
Fortunately, the tuning information is imprinted on the coherence properties of the beam exiting the MC slit, and we use this to determine the tuning at the slit to first and higher orders (Fig. 2). The end result is a repeatable tuning of the MC to 15 meV and better in Nion’s High Energy Resolution Monochromated EELS STEM (HERMESTM) (Fig. 3), as well as an ability to refocus the beam at the sample to sub-nm dimensions [2].
[1] L. Foucault, Comptes Rendus Academie des Sciences 47 (1858) 958-959.
[2] OL Krivanek et al, Microscopy 62 (2013) 3-21


Fig. 1:  Idealized far-field image (Ronchigram) for three electrons beams with slightly different energies, with astigmatism present at the MC slit.

Fig. 2:  Fourier transform of a Ronchigram obtained with the beam tuned at the MC slit.

Fig. 3:  Zero loss peaks before and after MC tuning

Type of presentation: Poster

IT-1-P-1486 Magnetic monopole like fields and electron vortices

Béché A.1, Van Boxem R.1, Van Tendeloo G.1, Verbeeck J.1
1EMAT, University of Antwerp, Groenenborgerlaan 171, 2020 Antwerp, Belgium
armand.beche@uantwerpen.be

The search for magnetic monopole particles has been in vain so far. However, an approximation to a magnetic monopole field can be obtained at the tip of a long, thin, nanoscopic magnetic needle [1,2]. We demonstrate that the interaction of an electron beam with such a field produces an electron vortex beam just like was predicted for a true magnetic monopole [3]. The total orbital angular momentum (OAM) produced by the magnetic needle can be precisely tuned by carefully selecting the amount of magnetic flux via the needle cross section.

The magnetic needle is extracted from a 60 nm thick nickel film using focused ion beam (FIB) milling. It is then deposited on top of a gold plated silicon-nitride grid with one end suspended over a pre-cut aperture hole (Fig.1 A). This aperture allows the impinging electron beam to interact with only one end of the needle. The magnetic field at the tip causes the fast electrons to obtain a spiral phase shift via the Aharanov-Bohm effect as revealed by holography in field free conditions in a transmission electron microscope (TEM) (Fig. 1B). The width of the needle is reduced in the FIB until the flux approaches one fluxon (total phase shift of to 2pi). Comparing the experimental results with simulations (Fig. 1C), an OAM of 0.8 was estimated.

In order to confirm the existence of a vortex after letting an electron beam interact with the magnetic needle aperture, a focal series was acquired in the far field plane of the needle (Fig.2 A). The presence of a dark center which does not disappear upon focusing is typical for a vortex beam, as demonstrated in simulated images (Fig. 2B). A second confirmation of the vortex character was made by cutting the slightly defocused far field images with the sharp edge of an objective aperture and noting the configuration of the Fresnel fringes [4]. Close to the vortex core, the phase dislocation pattern appears in the Fresnel fringes (Fig. 3A). The number of non-connected lines gives an approximation of the total OAM, close to 1 in the present case, confirming the holography result (Fig. 1B). The Fresnel fringes agree remarkably well with simulations (Fig. 3B).

An aperture containing such a monopole-like field provides a unique way of creating electron vortex beams with a pure OAM value, independent of the electron energy. As almost all the incoming electrons transforms into a specific OAM state, a high intensity vortex beam is created, greatly improving the potential for atomic scale magnetic measurements at much improved signal to noise ratios.

1. Béché A. et al., Nature Physics (2014), 10, p. 26-29.
2. Kasama T. et al., MRS Proc. (2004), 839, p. 107-118.
3. Aharonov Y. and Bohm D., Phys. Rev. (1959), 115, p. 485-491.
4. Verbeeck J. et al., Nature (2010), 467, p. 301-304.


This work was financially supported by the European Union: ERC grant 246791 COUNTATOMS, ERC Starting Grant 278510 VORTEX, Integrated Infrastructure Initiative grant 312483-ESTEEM2.

Fig. 1: A: Overview of needle surrounded by an aperture. B: Experimental phase map at the tip of the needle, figured by the dash square in B. The phase rosacea is scaled from 0 to 2pi. C: Simulated phase map for a total phase shift of 0.8x2pi over the full aperture.

Fig. 2: A: Experimental focus series of the aperture in far field conditions. The destructive interference center is typical of a vortex beam. B: Simulation of the focal series using the phase profile displayed in Fig. 1C.

Fig. 3: A: Cut of the defocused far field image of the needle aperture by a sharp edge, revealing a dislocation like feature in the Fresnel fringes close to the vortex center. As only one branch cannot connect, the total OAM is close to 1. B: Simulation of the cut aperture in the far field using the phase profile displayed in Fig 1C.

Type of presentation: Poster

IT-1-P-1500 Measuring the Orbital Angular Momentum of Electron Vortex Beams in the TEM

Guzzinati G.1, Clark L.1, Béché A.1, Verbeeck J.1
1EMAT, University of Antwerp, Groenenborgerlaan 171, 2020 Antwerp, Belgium
giulio.guzzinati@uantwerpen.be

The exchange of orbital angular momentum (OAM) in the interaction between an electron beam and a sample is determined by the properties of the sample and the beam [1,2]. Studying this interaction could enable a new class of OAM based microscopy techniques if convenient measurement of OAM exchange would exist. These techniques could then be used to study in the TEM, among others, the magnetic state of atoms and the transfer of OAM nanoparticles.

Electron beams possessing intrinsic orbital angular momentum have recently risen to attention after the prediction and demonstration of electron vortex beams[3-5]. This discovery has led to the rapid development in the field of singular electron optics [1-7].
In order to employ electron vortices as a probe to study the OAM exchange between a beam and a sample, methods to manipulate or measure the OAM of the beams are fundamentally important. While several methods have been designed to produce vortex beams, there has not been an equal progress in the detection and measurement of intrinsic OAM in the electron microscope.

Aiming to bridge this gap, we have implemented several diffraction based OAM measurement methods: using a forked grating hologram, a triangular geometrical aperture, a knife-edge and an astigmatic phase plate. Fig.1 shows an overview of the experimental results of the different methods when different incoming vortex beams are used as input.
In particular the triangular aperture and the astigmatic phase allow to recognize high order vortex beams easily , but they require to record and analyze a full 2D diffraction pattern. Intentional astigmatic aberration is easier to implement but the OAM is revealed by observing the beam waist rather than the far field pattern which may be a disadvantage in scanned electron probe setups.
On the other hand the hologram and the knife-edge are only appropriate for the measurement of lower values of OAM, but they allow the measurement to be reduced to a simple electron counting process which makes them ideally suited for automated OAM measurement [7].

[1] P. Schattschneider et al., Phys. Rev. B 85, 134422 (2012).
[2] A. Béché et al., Nat. Phys.10/1 26 (2013).
[3] K. Bliokh et al., Phys. Rev. Lett. 99 190404 (2007).
[4] M. Uchida and A. Tonomura., Nature, 464/7289 737 (2010).
[5] J. Verbeeck et al. Nature 494, 331–335 (2013).
[6] V. Grillo et al., Phys. Rev. X 4, 011013 (2014).
[7] G. Guzzinati et al., Phys. Rev. A 89, 025803 (2014).


We acknowledge funding from the European Union under the FP7 program, ERC Starting Grant No. 278510 VORTEX and Integrated Infrastructure Initiative No. 312483 ESTEEM2.

Fig. 1: (a) Schematic representation of the experiment, depending on the OAM of the input beam and type of aperture used different patterns are produced. Experimental data are show for (b) forked hologram, (c) triangular aperture (d) knife-edge (e) astigmatic aberration.

Type of presentation: Poster

IT-1-P-1541 Investigation of physical and chemical method to produce Möllenstedt electrostatic biprism for off-axis electron holography experiment

Cours R.1, Houdellier F.1
1CEMES-CNRS, Université de Toulouse, 29 Rue Jeanne Marvig, 31055 TOULOUSE FRANCE EU
robin.cours@cemes.fr

Denis Gabor has developed electron holography in 1948, as a method used to quantitatively retrieve the phase of the electron wave. D. Gabor proposed a configuration where the perturbed wave (object wave) and the reference unperturbed wave are observed in a common optical plane below the sample. In this plane a superimposition of the two waves can occur. This superimposition will induce an interference phenomenon and create the so-called in-line electron hologram, used to retrieve the phase difference between the two waves. In this configuration the sample is then out of focus. In 1955 G. Möllenstedt and H.Düker invented the biprism for electrons, a metallic wire biased relatively to the earth. The biprism effectively splits the electron beam into an object wave and a reference wave, which by electrostatic fields are brought to overlap onto one another. An interference pattern will be observed below the wire plane while the sample can still be in focus. This configuration, known as off-axis electron holography, is the one commonly used in all the major holography studies from dopant profiling to strain mapping through studies of nanomaterials magnetic configurations. Biprisms in common use today are constructed by coating ultrasmall quartz fibers with noble metals. The resulting biprisms, although they are quite small by most fabrication standards (approximately 700 nm in diameter), can have various mechanical, electrical, structural … properties. Depending on the quality of the biprism, the properties of the off axis hologram can be strongly modified. As an example, to avoid vibration, which drastically decrease the interference fringes contrast, the wire should be very taut; to minimize charge effect, which induce Fresnel fringes phenomena, the wire should be extremely clean; to increase the phase coherence of the beam across the biprism the wire should be the smaller possible, …

Regarding all these drastic requests that the wire should fulfilled to be a suitable biprism, the question of reproducibility become deeply problematic using standard biprism fabrication method. This question become even more crucial regarding our new microscope, the In situ interferometry TEM (I2TEM), a HF3300 TEM that fits with 4 biprisms wire used for various electron holography developments. In order to choose the most reproducible way which will give the best wire properties (size, vibrations, cleanliness, …), we have investigated several methods to produce them from chemical method to FIB (Focused Ion Beam) approach. The combination of these methods allowed us to make numbers of high quality biprism wire with a higher reproducibility rate.


This work has been supported by the French National Research Agency under the "Investissement d'Avenir" program reference No. ANR-10-EQPX-38-01.

Fig. 1: A: SEM image of a Wollaston wire thinned using a FEI Helios FIB B: The same wire installed inside a Hitachi HF2000 TEM

Fig. 2: C: Special and ultrafast chemical etching method using nanowetting of hot HNO3 onto a Wollaston wire

Type of presentation: Poster

IT-1-P-1672 Development of illumination system of a 1.2 MV Field Emission Transmission Electron Microscope

Kawasaki T.1, Kasuya K.1, Furutsu T.1, Ono S.1, Arai M.1, Moriya N.1
1Central Research Laboratory, Hitachi, Ltd., Hatoyama 2520,Saitama 350-03, Japan
takeshi.kawasaki.qb@hitachi.com

       In the FIRST Tonomura project, we have been developing a 1.2 MV field-emission transmission electron microscope (FE-TEM) for the atomic resolution three-dimensional reconstruction of electro-magnetic fields by electron holography. Here FIRST stands for funding program for world-leading innovative R&D on science and technology. In this paper we report its illumination system with the following requirements:
       (1) high brightness beam for electron holography
       (2) current fluctuation less than 10 % over 8 hours for stable observation
The requirement (1) is discussed in this presentation and the requirement (2) is discussed by Kasuya in this conference.
       Figure 1 shows schematic view of the illumination system and three ray paths. Separate valves are placed between the FE gun and the accelerator tube so that conditioning of emission and high-voltage can be performed separately. The pre-accelerating magnetic lens focuses the beam near the first electrode of the accelerator tube where the Butler lens is formed (Case A), and then the spherical aberration of the accelerator tube can be suppressed. When the magnetic lens excitation becomes stronger, the electron trajectory focuses twice in the accelerator tube (Case C). Between Case A and Case C, beams focus near the condenser lens and cannot focus on the specimen position (Case B). To obtain high brightness beam, total aberration of the illumination system has to be minimized. The optimum condition of the pre-accelerating magnetic lens was obtained by calculating mean brightness and probe current of the spot focused on the specimen position as a function of the lens excitation using WR5 software (MEBS Ltd.). The FE-cathode source diameter, the angular current density, and the energy spread are assumed to be 5 nm, 30 μA/sr, and 0.3 eV, respectively. Figure 2 shows the results. Two peaks of the brightness exist: The left peak corresponds to Case A, the right peak corresponds to Case C, and the bottom region D between two peaks corresponds to the Case B. Preliminary experimental results using the 1 MV FE-TEM showed the following:
       (1) existence of two brightness peaks
       (2) the maximum brightness of 1.8×1010 A/cm2sr [1]
This brightness value is almost the same as that calculated for the 1 MV FE-TEM. The calculated maximum brightness is 3.3 ×1010 A/cm2sr for the 1.2 MV FE-TEM. We expect it to reach 5×1010 A/cm2sr by increasing the angular current density of the cleaner FE-tip under ultra high vacuum condition (3.0×10-10 Pa) [2].

References
[1] T. Kawasaki et al.  J. Electron Microsc. 49 (2000) 711-718.
[2] K. Kasuya et al. submitted to  J. Vac. Sci. Technol. B.

 


This research was supported by the Japan Society for the Promotion of Science through the FIRST Program initiated by the Council for Science and Technology Policy.

Fig. 1: Schematic view of the illumination system and different ray paths A, B, and C.

Fig. 2: Calculated brightness and probe current as a function of the excitation of the magnetic lens in terms of IN(V1)-1/2 , where I is the lens current, N is the number of turns of the coil (1700), and V1 is the FE extraction voltage    (5 kV).

Type of presentation: Poster

IT-1-P-1693 Measurement of current density distribution in shaped e-beam writers

Horáček M.1, Bok J.1, Kolařík V.1, Urbánek M.1, Matějka M.1, Krátký S.1
1Institute of Scientific Instruments AS CR, v.v.i., Brno, Czech Republic
mih@isibrno.cz

The ZrO W(100) Schottky cathode is used in our e-beam writing system working with a rectangular-shaped electron beam. The homogeneous angular current density distribution is crucial for quality of exposures of the shaped beam lithography systems. Two basic types of the angular emission distribution can be observed in dependence on the microscopic final end form shape of the emitter tip, with bright centre and more common dark centre [1]. The stable operation of the cathode thus stable end form shape requires a delicate balance of parameters inside the gun which however can slightly change during cathode life time. This implies the necessity of analysing and periodical monitoring the current density distribution in e-beam. Four methods enabling this measurement are presented.
First we implemented a method based on the modified knife-edge approach [2], when a part of the scanned element of the beam is blanked out and the current within the remaining "open" part is measured. The 2D information of the current distribution is obtained by stepwise opening of selected segments. The measurement error analysis was made and necessary measurement averaging in each segment were used in order to reduce the random error of the current [3]. The size of the scanned element was 6 × 6 µm2, a maximum usable segment for one shot in our lithography system (Fig. 1).
The current distribution obtained by the knife-edge method was compared with a method using a luminescent screen. The YAG:Ce single-crystal screen was irradiated by the e-beam stamp of the 6 × 6 µm2 and the areal light emission was recorded by a magnifying optical system with a CCD camera. The emitted light intensity is directly proportional to the e-beam current, thus the current density distribution can be compared with other measurements methods. However, the absolute measurement is hardly possible (Fig. 2).
Next the same e-beam stamp of the 6 × 6 µm2 was scanned over Faraday cup opening. The advantage of this method is uniform distribution of the measurement error instead of the modified knife-edge method. The absolute value of the current density is affected by the demagnification of the electron optics during measurement (Fig. 3).
Another method is based on evaluation of developed electron resist exposed by the 6 × 6 µm2 separate shaped e-beam stamp using atomic force microscope. The depth of the developed resist depends on the spread of the energy in the electron resist. The real current density distribution was obtained by the deconvolution of the developed resist with electron scattering model (Fig. 4).

References

[1] K-Liu et al., J. Vac. Sci. Technol. B 28, C6C26 (2010).
[2] M. Sakakibara et al., Jpn. J. Appl. Phys. 46, 6616 (2007).
[3] J. Bok et al., J. Vac. Sci. Technol. B 31, 31603-1 (2013).


The authors acknowledge the support from MEYS CR (LO1212) together with EC (ALISI No. CZ.1.05/2.1.00/01.0017), the TACR project No. TE01020118 and institutional support RVO:68081731.

Fig. 1: Modified knife-edge method.

Fig. 2: Luminescent screen method.

Fig. 3: Faraday cup method.

Fig. 4: Electron resist exposure method.

Type of presentation: Poster

IT-1-P-1958 Quantitative measurement of the OAM spectra of electron vortex beams.

Clark L.1, Béché A.1, Guzzinati G.1, Verbeeck J.1
1EMAT, University of Antwerp, Antwerp, Belgium
laura.clark@uantwerp.be

Electron vortex beams have been subject to a great level of interest since their first demonstration only a few years ago [1]. Much of the interest in the field stems from their potential to measure magnetic transitions within a sample, at a previously unreachable scale. While much progress has been made, in producing electron vortices of high purity, high intensity and atomic scale, research into the required counterpart towards full experimental application, of orbital angular momentum (OAM) measurement, has not yet matured to its full potential [2-4].

In the last 12 months, the first methods to measure the OAM make-up of an electron vortex beam have been demonstrated [5-7]. However, the methods presented thus far, are limited to only those cases where the input beam is in a single vortex state, and do not allow measurement of the relative weightings of vortex states in a beam . Indeed, a generic electron wave can be seen as a superposition of multiple vortex modes and the weight of each of these modes can in principle be measured.

We introduce here an experimental technique able to measure the relative weightings of 5 or more OAM modes within an input beam, through the use of a multi-pinhole interferometer (MPI). This is a technique which has recently been used to measure the strength and location of optical vortices, but which is easily adaptable to practical implementation in a TEM, placing an MPI aperture in the SA plane, below the sample.

Experimental results are shown, having measured the OAM spectrum of pure l={-1,0,+1,+2} centred vortex beams, enabling the first quantitative discussion of their experimental purity. We further demonstrate the so-called mode broadening effect, by measuring the changes in OAM composition as a vortex beam is shifted away from the central axis of measurement.

This application of an MPI within a TEM has enabled measurement of an approximate OAM spectrum in the SA plane. We give experimental evidence alongside theoretical models, enabling rapid discrimination of different orders of vortex beams even if the electron beam consists of a superposition of different OAM modes. This capability serves as a promising tool to measure OAM exchanges in the interaction of electrons with a sample.

[1] Bliokh, KY, et al. PRL 99.19 (2007): 190404
[2] Verbeeck, J., et al. Nature 467.7313 (2010): 301-304
[3] Clark, L., et al. PRL 111.6 (2013): 064801
[4] Béché, A, et al, Nature Physics 10.1 (2014): 26-29
[5] Guzzinati, Giulio, et al. arXiv: 1401.7211 (2014)
[6] Saitoh, K, et al. PRL 111.7 (2013): 074801
[7] Shiloh, Roy, et al. arXiv: 1402.3133 (2014)


We acknowledge funding from the European Union under the FP7 program: ERC Starting Grant No. 278510-VORTEX and Integrated Infrastructure Initiative Reference No. 312483-ESTEEM2.

Fig. 1: Plot of idealised electron vortex beam – brightness represents intensity, and hue represents phase

Fig. 2: A five-pinhole multi-pinhole interferometer, enabling measurement of OAM modes in the set l={-2:+2}

Fig. 3: Experimental diffraction pattern from an l=+1 vortex centred on the MPI

Fig. 4: Autocorrelation function produced from the experimental diffraction pattern.

Type of presentation: Poster

IT-1-P-1961 Development of Phase Contrast Scanning Transmission Electron Microscopy

Iijima H.1, Minoda M.2, Tamai T.2, Kondo Y.1, Hosokawa F.1
1EM Business Unit, JEOL Ltd., 3-1-2 Musashino, Akishima, Tokyo 196-8558, Japan, 2Department of Applied Physics, Tokyo University of Agriculture and Technology, 2-24-16 Nakacho, Koganei, Tokyo 184-8588, Japan
hiiijima@jeol.co.jp

Phase contrast transmission electron microscopy (P-TEM) is a powerful tool to enhance the image contrast of transparent materials such as ice-embedded biological specimens and polymer materials. In P-TEM, a phase plate is placed at the back-focal plane (BFP) of the objective lens (OL). It gives a phase shift for scattered electron waves, resulting in a change of phase contrast transfer function (PCTF) from sine to cosine type. Eventually, phase variation of specimens is converted into intensity variation. Among various types of phase plates, a carbon film phase plate with a small central hole is the most practical1. However, there is a serious issue that high-density electron beam (cross-over) on the phase plate causes the charging and/or the alteration of the phase plate, resulting in decreasing the life time of the phase plate.

To overcome this issue, we are developing phase contrast scanning transmission electron microscopy (P-STEM). Figure 1 shows the schematics of P-TEM and P-STEM. According to the reciprocity theorem, the same contrast appears in the P-TEM and the P-STEM if a phase plate is placed at a front-focal plane (FFP) of an OL in P-STEM. In P-STEM, a cross-over is not formed on the phase plate, so that improvement of the phase plate life time is expected. In our experiments, we used a field emission electron microscope (JEM-2100F) equipped with a Schottky electron source, to obtain a coherent small probe on a specimen. Phase plate is placed on a condenser lens aperture plane conjugate to the FFP of the OL.

On the other hand, it is well known that the small detection angle is needed to obtain good phase contrast in STEM imaging. Figure 2 compares a conventional bright-field STEM and a P-STEM images of amorphous carbon film with different detection angle shown in Fig. 1. And Fourier transforms of the conventional bright-field STEM image and the P-STEM image with β = 4 mrad show the sine shape. By contrast, that of the P-STEM image at β = 0.3 mrad shows the cosine shape, which proves that the P-STEM can be achieved with small detection angle.

[1] R. Danev and K. Nagayama, J. Phys. Sci. Jpn. 70 (2001) 696.


This development was supported by the program for "Development of Systems and Technologies for Advanced Measurement and Analysis" under JST.

Fig. 1: Schematic of P-TEM (left) and P-STEM (right). The phase plate is placed at the BFP of the objective lens in P-TEM and the FFP of the objective lens in P-STEM.

Fig. 2: Conventional bright-field STEM and P-STEM images of amorphous carbon film. All images are taken close to focus. (a) Conventional bright-field STEM image. (b) P-STEM image with β = 4 mrad. (c) P-STEM image with β = 0.3 mrad. (d)-(f) Fourier transforms for images shown above. Scale bars; 10 nm in (a)-(c), 4 nm-1 in (d)-(f).

Type of presentation: Poster

IT-1-P-2090 Contrast enhancement of phase objects by using Phase Contrast Scanning Transmission Electron Microscopy

Minoda H.1, Tamai T.1, Iijima H.2, Hosokawa F.2, Kondo Y.2
1Department of Applied Physics, Tokyo University of Agriculture and Technology, 2EM Business Unit, JEOL Ltd
hminoda@cc.tuat.ac.jp

It is well known that an interaction between electron waves and molecules composed of light elements such as biological molecules is very weak. Therefore, it is very difficult to obtain their high contrast image in transmission electron microscopy (TEM). Contrast enhancement of the phase objects by using a phase plate was proposed at the middle of the 20th century [1], but it was realized at the beginning of 21st century [2]. In the pioneering work by Nagayama, a carbon thin film with a hole in its center is used as a phase plate (PP) and it was placed at a back focal plane (BFP) of the objective lens (OL). A role of the PP is giving a phase shift to scattered wave by means of the mean inner potential of the PP material. Electron waves having a phase shift interfere with electron waves without phase shift. Accordingly, phase image would be able to be visualized.

Applying the principle of reciprocity to scanning transmission electron microscopy (STEM), imaging optics of the STEM is equivalent to that of a conventional TEM. Therefore, a phase contrast scanning transmission electron microscopy (P-STEM) can be used to enhance phase contrast of the phase objects. In the present study, a PP can be set on the condenser lens aperture (CLA) plane that is optically equivalent to a front focal plane (FFP) of an OL. The P-STEM image which enhances image contrast could be obtained by getting an appropriate optical condition. Figure 1 show an example of the comparison of (a) the conventional STEM bright field image and (b) the P-STEM image. Ferritin molecules were used as a specimen. This comparison clearly shows contrast enhancement in P-STEM. In this paper, the results obtained by sung phase contrast microscopy to the STEM mode are introduced.

[1] F. Zernike, Physica 9 (1942) 686.

[2] R. Danev and K. Nagayama, J. Phys. Sci. Jpn. 70 (2001), 696.


This development was supported by SENTAN, JST.

Fig. 1: A comparison of (a) C-STEM and (b) P-STEM images of ferritin molecules.  The contrast enhancement in P-STEM is evident.

Type of presentation: Poster

IT-1-P-2325 AC- Voltage Operated Schottky Electron Source

Yada K.1, Saito Y.1
1Daiwa Techno Systems Co.,Ltd
yada@daiwatechno.co.jp

Introduction: Zr-O/W100 Schottky electron source has been widely used in electron beam instruments because of its favorable properties such as 1)vacuum technological tolerance at its operational condition, 2) rather long life time and high brightness. It is still required, however, that a vacuum must be better than 10 -8 Pa and stability of high voltage must be better than 10 -5 when DC high tension and electro-magnetic lens systems are used in the instrument. We tried to find promising materials. Among them, we selected BaZrO3 and SrZrO3 and tested their thermal field emission properties with both DC and AC high tension powers.

Results: Field emission tips of 110- and 100-oriented W wire were made by electrolytic method and powder of BrZrO3 or SrZrO3 was pasted near the apex as usual. Thermal field emission patterns obtained by DC and AC voltage are very similar and crystal facets are indexed very easily. Fig.1 and Fig2 show emission patterns of BaZrO3/W(110) and BaZr3(W100) obtained by AC and DC operation ,where optimal working temperature is 800 degree C. Similar results are obtained with SrZrO3(W)100 cathode but its optimal temperature is little higher than the case of BrZrO3.

As advantages of AC operation of present schottky electron source, followings are concluded:

1)Schottky shield is not necessary because of low working temperature of the emission materials.

2)Emission beams can be focused, deflected and stigmated by using electrostatic lens,deflector and stigmtor, respectively.

3)Work function of newly adopted materials here is so small that working temperature is fairly low (800-850 degreeC). Consequently ,energy spread of the beam will be narrow.

4) As commercial AC electric supply can be used without any rectifire or stabilizer, factor cost will be fairly reduced.

5)We think that the present Schottky electron source is the best selection for a generation of strong and small X-Ray source of projection X-ray microscope.


Fig. 1: Fig.1

Fig. 2: Fig.2

Type of presentation: Poster

IT-1-P-2198 The effect of Detector Thickness on Direct Detector Performance

Clough R. N.1, Moldovan G.2, Kim J. S.1, Kirkland A. I.1
1Department of Materials, University of Oxford, UK, 2Oxford Instruments NanoAnalysis, High Wycombe, UK
robert.clough@materials.ox.ac.uk

Direct detection refers to a detection system where signal is generated in the sensor chip directly by the imaging electrons; indirect systems generate photons in a scintillator from the imaging electron and it is these photon which are coupled to the sensor chip that generate signal. One of the key advantages of a direct detection system is the possibility of producing thin detectors; these are desirable as a thin detector has improved detection performance in terms of Modulation Transfer Function (MTF) and Detective Quantum Efficiency (DQE) [1]. This improvement arises from the fact that many electrons will pass all the way through the sensor and escape the detector system generating signal along the way, before large lateral scattering has occurred.

We have taken a prototype CMOS based direct detector featuring full frame resolution of 1024 by 1024 pixels, with a pixel size of 20µm and readout of 30fps [2]. Two different versions of the detection chip were produced. The first is a 20µm thick p- active layer on a p+ substrate mechanically thinned to 50µm. The second was made from silicon on insulator (SOI) wafer with a 20µm device layer with the handle wafer removed using a chemical etch. For each of these detectors the MTF and DQE were measured using standard techniques [3] at 80 and 200kV. Here we shall present the characterisation data along with images of gold particles on an amorphous substrate to show how thinner detectors lead to improved detector performance, allowing images taken at lower magnification to have improved resolution.

[1] G. McMullan, et al, Experimental observation of the improvement in MTF from backthinning a CMOS direct electron detector, Ultramicroscopy, 109 (2009).
[2] A.J. Wilkinson, et al, Direct Detection of Electron Backscatter Diffraction Patterns, Phys. Rev. Lett. 111 (2013).
[3] R. R. Meyer, et al, Experimental characterisation of CCD cameras for HREM at 300kV, Ultramicroscopy, 85 (2000).


We would like to acknowledge Dr T. Anaxagoras and Prof. N. Allinson from the University of Sheffield for provision of CMOS wafers, and C Wilburn of Micron Semiconductor Ltd. for chip packaging.

Fig. 1: MTF of a 20µm thick detector at 80 and 200kV.

Fig. 2: Au on amorphous Carbon at 80kV and 120,000x magnification taken with a 20µm thick sensor chip.

Fig. 3: Au on amorphous Ge at 200kV and 120,000x magnification taken with a 20µm thick sensor chip.

Type of presentation: Poster

IT-1-P-2263 Maximising Phase Contrast in Aberration-corrected STEM using Pixelated Detectors

Yang H.1, Pennycook T. J.1,2, Nellist P. D.1,2
1University of Oxford, Department of Materials. Parks Rd, Oxford, OX1 3PH, UK, 2EPSRC SuperSTEM Facility, Daresbury Laboratory, WA4 4AD, UK
hao.yang@materials.ox.ac.uk

For imaging weak phase biological specimens, phase contrast imaging using elastically scattered electrons provides the most information for a given amount of radiation damage as compared to electron inelastic scattering as well as X-ray and neutron scattering [1]. In scanning transmission electron microscopy (STEM), most phase information from weak scattering objects lies inside the bright field disc of the convergent beam electron diffraction pattern, which can be reconstructed using the method described by Rodenburg et al [2]. In this work we show that, compared to alternative modes including annular bright field (ABF) and differential phase contrast (DPC), phase contrast using a pixelated detector generates higher contrast in reconstructing the phase and therefore enjoys a higher dose efficiency in imaging weak phase objects.

With zero aberrations, any centrally symmetric detector will give no contrast for a weak phase object, as the two sides of disc overlapping regions in the convergence beam electron diffraction pattern are pi out of phase under weak phase approximation, and cancel each other when integrated using a central symmetrical detector geometry. Therefore, asymmetric detector geometries like DPC are expected to have higher phase contrast than ABF. In DPC, the quadrant detector can be divided into more segments with different collection angles, and the contrast transfer function is found to depend on the collection angles used, therefore the detector geometry of DPC can be further optimized to collect the maximum phase information per detected electrons. A pixelated detector provides even greater flexibility over where the information in the bright-field disc is retrieved from for each spatial frequency in the image.

Simulations have been done using an arbitrary weak phase specimen whose maximum atomic potential equals to that of a carbon atom, and has a Gaussian shape with a full width half maximum (FWHM) of 1nm. The artificially high width of the object is designed to test the lower spatial frequency transfer. The reconstructed phase with a dose as low as 50 electrons/Å2 and Nyquist resolution of 4.6Å still shows an interpretable feature (Figure 1). This dose is close to the critical dose of 5-50 electrons/Å2 for imaging biological specimen. In contrast to using a pixelated detector, neither ADF, ABF (Figure 2) nor DPC (Figure 3) show any recognizable structure feature under the same dose of 50 electrons/Å2. The formation of image contrast in ABF relies the presence of aberrations for a weak phase object, and here we are assuming an aberration-corrected microscope with zero residual aberrations.

[1] Henderson, R. Quarterly Reviews of Biophysics 28, 171-193 (1995).

[2] Rodenburg, J. M. et al. Ultramicroscopy 48, 304-314 (1993).


The authors would like to acknowledge financial support from the EPSRC (grant number EP/K032518/1) and the EU Seventh Framework Programme: ESTEEM2.

Fig. 1: Figure 1: Phase retrieval using a pixelated detector. (a) Schematic of a high speed pixelated detector. (b) A weak phase object with a maximum phase change of 0.15 radian. Reconstructed phase (c) assuming no noise, (d) with shot noise and a dose of 50 electrons/Å2. The scale bar is 5nm.

Fig. 2: Figure 2: Simulated (a,b) ADF and (c,d) ABF images of the weak phase object. The intensity is normalized to the number of incident electrons. (a)(c) assume no noise in image, and (b)(d) consider shot noise with the electron dose being 50 electrons/Å2. The scale bar is 5nm.

Fig. 3: Figure 3: Differential Phase Contrast (DPC) imaging using a 4-quadrants detector in (a). The simulated STEM DPC images using both (b,c) A-C quadrants, and (d,e) B-D quadrants, where (b)(d) assume noise free, and (c) (e) considers shot noise with the electron dose being 50 electrons/Å2. The scale bar is 5nm.

Type of presentation: Poster

IT-1-P-2346 Effects of dielectric substrate on localized surface plasmon in a silver nano-particle

Fujiyoshi Y.1, Nemoto T.1, Kurata H.1
1Institute for Chemical Research, Kyoto University, Kyoto, Japan
fujiyoshi@eels.kuicr.kyoto-u.ac.jp

Recently localized surface plasmons (LSPs) which are collective oscillation of conduction electrons of metallic nano-particles (NPs) attract researchers in nano-optics because of strong optical confinement and electric field enhancement, leading to many applications including biochemical sensors and surface-enhanced Raman spectroscopy (SERS) etc. Since the dielectric environment around the NP affects the property of LSPs, it is important to elucidate the effects of dielectric materials supporting NPs on LSPs.
In the present work, we examined special distributions of LSP excited on a silver NP supported by MgO substrate using electron energy loss spectroscopy (EELS) combined with scanning transmission electron microscopy (STEM). Spectral imaging (SI) data were acquired along the direction parallel to the MgO surface supporting a silver NP, which enabled us to observe the intensity distribution of LSP excitation as a function of the distance from the silver NP/MgO interface. The experiment was performed by an aberration corrected STEM (JEM-9980TKP1) equipped with a cold-FEG.
Figure 1 and 2 show a HAADF image of silver NP on MgO substrate and its LSP map extracted from SI data, respectively. From the HAADF image the NP can be regarded as a sphere. When a spherical metal particle is isolated in vacuum, the excitation probability of LSP should distribute isotropically around the particle. However, the LSP map in Fig. 2 shows anisotropic distribution, that is, the intensity at the top surface of silver NP is strong compared to that at other positions, which means that the effect of dielectric substrate is remarkable. In order to interpret such anisotropic distribution, we simulated the electromagnetic field induced in the silver NP on MgO substrate using finite-difference time-domain (FDTD) method.
Figure 3 shows the spatial distribution of field calculated by assuming the incident plane waves polarized perpendicular (a) and parallel (b) to the substrate surface. When the polarization of incident wave is perpendicular to the substrate, the field strength in the NP on MgO is enhanced compared to that in the isolated NP as shown in Fig. 3(a), which corresponds to the observed strong excitation at position A in Fig. 2. In case of the parallel polarization the field strength in the NP on MgO is weakened (Fig. 3(b)), corresponding to the observed intensity at position B in Fig. 2. Therefore, the anisotropic distribution of the LSP excitation in silver NP on MgO surface can be attributed to the direction of electric polarization induced in the NP depending on the electron positions.


Fig. 1: HAADF image of a silver NP supported on MgO surface.

Fig. 2: LSP map extracted from the energy range from 3.2 to 3.6 eV in the SI data.

Fig. 3: Spatial distribution of electromagnetic field calculated by FDTD simulations. Incident plane waves were assumed to be polarized parallel (a) and perpendicular (b) to the substrate surface. Solid and broken lines correspond to the intensity profiles for an isolated silver NP and the silver NP supported on MgO surface, respectively.

Type of presentation: Poster

IT-1-P-2410 Thermal Emission Properties of GdB6 Cathode

Saito Y.1, Yamagishi K.1, Yada K.1
1Daiwa Techno Systems Co.,Ltd.
saito@daiwatechno.co.jp

Introduction: Hexa boride of lanthanum (LaB6) has been widely used in electron beam
instruments because of its higher brightness than that of tungsten hairpin cathode. But it might
be that there are better materials than LaB6. Among many borides of lanthanide, hexa-borides
of Ce and Gd are promising from the existing data[1] based on Richardson-Dushman equation as
shown in Table 1. So we tested electron emission properties of GdB6.


Table 1
              A      φ(eV)       AT2         I(A/cm2)
LaB6      29      2.66     93960000   3.316363
GdB6    0.84     2.06     2721600     4.607963
GdB6     9.3      2.55     30132000   2.16234
GdB6      10      2.58     32400000   1.915988


R-D Eq. I=AT2EXP(-φ/kT), A:R-D constant, k:Boltsman cont, T:temperature(1800k), I:electron
density


Results: Fig.1 shows photograph of Ta wire covered with GdB6 powder where the central part is
slightly protruded. Fig.2 is beam pattern of GdB6 cathode at working temperature when the
cathode is installed in a scanning electron microscope. Fig.3(a) shows SEM image of ZnO
particles obtained with tungsten hairpin cathode and Fig.3(b) shows that obtained with GdB6
cathode. It is seen that image quality of (b) is superior to that of (a). It is also clear that emission
performance of present GdB6 powder cathode is nearly equal or little better than that of LaB6
single crystal. Sintered GdB6 cathode is now under examination to compare with the single
crystal LaB6 cathode.


Reference:[1]Japan-Soviet Communication “Emission Characteristics of Materials” pp.96-81 by
V.S.Fomenko, Published by Naukova Dunka, Kiev 1970


Fig. 1: Fig.1

Fig. 2: Fig.2

Fig. 3: Fig.3(a)

Fig. 4: Fig.3(b)

Type of presentation: Poster

IT-1-P-2469 Electron differential phase microscopy with an A-B effect phase plate

Tanji T.1, Ikeda U.2, Niimi H.2, Usukura J.1
1EcoTopia Science Institute, Nagoya University, Nagoya, Japan 1, 2Graduate School of Nagoya University, Nagoya, Japan 2
tanji@esi.nagoya-u.ac.jp

    Observations of week phase objects, such as thin films of light elements, thin polymer films, biological sections etc., are available by electron phase microscopy[1]. Many of phase plates utilized are thin film types. Some electrostatic types have been developed, but they are not so general, because the fabrication of the filter with fine structures is very difficult. The mainstream of todays phase plate is the thin film type. This type of the phase plate, however, has some disadvantages, i.e. control of the film thickness, charging up, contamination and so on. We adopted the phase plate with a magnetic thin filament which generates the vector potential around itself by an Aharonov-Bohm (A-B) effect. The filament type phase plate with the A-B effect was proposed and constructed firstly by Nagayama. This type of the phase plate generates the differential phase contrast in the image, and has a longer life time than the thin film type. Any clear differential effect, however, has scarcely reported so far.
    We will report that the effect of a phase plate consisting of a Wollaston platinum filament of 1 µm in diameter covered with ferromagnetic material, Nd-Fe-B of 5 nm thick, deposited by Pules Laser Deposition. The filament with a clean surface selected by SEM is mounted on a single hole Cu grid. The phase difference in the both side spaces of the filament measured by electron holography shows 1.5 rad as shown in Fig.1. Being set on the aperture holder, the phase plate is inserted in the back focal plane of the objective. Figure 2 shows images of a colon bacillus stained with Pb. Fine structures can be observed clearer in the image using the phase plate than in the image taken ordinarily. The direction of the differentiation is shown by the arrowhead.

Refernce

[1]K. Nagayama, Another 60 years in electron microscopy: development of phase-plate electron microscopy and biological applications, Journal of Electron Microscopy, 60(2011) S43-S62.


Fig. 1: (a) Electron phase map reconstructed by electron holography. (b) A line profile along the arrow head in (a) which is averaged along the long side of the rectangle. The phase difference is about 1.5 rad. between both sides of the filament.

Fig. 2: mages of a colon bacillus stained with Pb taken at under-focus condition without the filament(a) , and in-focus with the filament(b).

Type of presentation: Poster

IT-1-P-2476 Effect of a phase plate on TEM imaging

Edgcombe C. J.1
1TFM Group, Dept of Physics, University of Cambridge
cje1@cam.ac.uk

The type of phase plate that has been most widely reported (eg [1]) consists of a plain disc of material such as carbon, of controlled thickness, with a central hole to pass the direct beam. Images made with this type of plate show bright outlines or halos around certain features [2, 3].  Analysis of geometrical imaging has shown how these halos occur.  It is necessary to consider the response to all spatial frequencies that are present in a typical object. In principle this can be done straightforwardly by Fourier transforming the object phase to find its spatial frequency distribution at the back focal plane (BFP), multiplying by the response of the phase plate and further transforming to find the image distribution.

The response has been found [4] for a weak phase object consisting of a circular disc of radius b, centred on the microscope axis. The phase plate is assumed to advance the phase of components with angular frequencies greater than a value q0 , defined as
q0 = 2 (pi) r2 ⁄ λf
where r2 is the radius of the central hole in the plate for the direct beam, λ is the electron wavelength and f is the focal length of the lens. The resulting image intensity is shown in figure 1 for a phase advance of (pi)/2 and a range of values of B = q0b. The object is imaged with little overshoot when B is less than about 1. Reported results [5] agree with this transition value for B.

The step changes at radius b are always imaged fully but as B increases, the low-frequency components are progressively lost from the image and for B > 1, the mean intensity across a step falls to the background value.  The full range of the step is maintained, so the intensity changes from +(half the range) at radii just less than b, to –(half the range) just outside the step. Thus a bright halo or outline is produced just outside the boundary r = b, for objects with B > ~1.  The darker central patch for B = 8 agrees with observation [3]. The maximum object diameter that corresponds to Bmax , the maximum B for accurate imaging, is
2bmax = 2Bmax ⁄ q0 = λ Bmax f ⁄ (pi) s2
where s2 is the radius of the hole needed to pass the direct beam.  To increase the size of object that can be imaged accurately, it will be necessary to reduce s2 or increase the focal length of the objective lens.

References
Danev R and Nagayama K 2001 Ultramicroscopy 88 243-52
Fukuda Y, Fukazawa Y, Danev R, Shigemoto R and Nagayama K 2009 J Struct Biol 168 476-84
Danev R and Nagayama K 2011 Ultramicroscopy 111 1305-15
Edgcombe C J 2014 Ultramicroscopy 136 154-9, http://dx.doi.org/10.1016/j.ultramic.2013.09.004
Hall R J, Nogales E and Glaeser R M 2011 J Struct Biol 174 468-75


Fig. 1: Figure 1. Image intensities produced by a (pi)/2 phase plate with fixed q0 for uniform disk objects with a range of diameters 2b.  Responses are shown for values of B = q0b (increasing from top left) of 0.2, 0.5, 1, 2, 5 and 8.

Type of presentation: Poster

IT-1-P-2510 Toward electron polarizators

Grillo V.1,2, Karimi E.3, Balboni R.4, Gazzadi G. C.1, Frabboni S.1,5, Mafakheri E.1,5, Tang W. X.6,7, Boyd R. W.3,8
1CNR-Istituto Nanoscienze, Centro S3, Via G Campi 213/a, I-41125 Modena, Italy, 2CNR-IMEM, Parco delle Scienze 37a, I-43100 Parma, Italy. , 3Department of Physics, University of Ottawa, 150 Louis Pasteur, Ottawa, Ontario K1N 6N5, Canada, 4CNR-IMM Bologna, Via P. Gobetti 101, 40129 Bologna, Italy, 5Dipartimento FIM, Universitá di Modena e Reggio Emilia, Via G. Campi 213/a, 41125 Modena, Italy, 6College of Materials Science and Engineering, Chongqing, 400044, China, 7School of Physics, Monash University, Clayton, VIC, 3800, Australia, 8Institute of Optics, University of Rochester, Rochester, New York 14627, USA
vincenzo.grillo@cnr.it

We describe the experimental and theoretical improvements toward the realization of an efficient electron spin polarizator. The initial proposed polarizator [1] was based on the spin-orbit conversion of a vortex beam [2] to a beam with a defined polarization. The conversion occurred within a compensated quadrupolar Wien Filter (WF).

The theoretical improvements are supported by simulations of the beam-field interaction through a new multislice for propagation including spin [3]. The experimental steps are based on the introduction of phase holograms to produce e-beams close to ideal Bessel beams [4]. To improve the flexibility and feasibility of the polarizer we have considered different possible alternative design: e.g. when the pitch fork hologram is positioned below the WF it is possible to obtain simultaneously the 2 polarized beams and switch between them [3]. Alternative design permit also to remove the electric fields. We have also studied the higher order corrections of the WF by magnetic multipoles of higher order and calculated the possible effects of the fringing fields: the efficiency in the selection of the polarized states increases with the order of the vortex and consequently of the multipoles in the WF.

Fig 1 is an example of simulation of the wavefunction after a WF for a beam at 15 KeV (e.g. for SPLEEM and low voltage TEM applications ) for 2 initial spin state. The brightness is proportional to the wave intensity, the phases encoded in the color. Due to the spin orbit coupling different spin are transformed, inside the WF, in different phase factors and orbital momentum. Only the center of the state |ℓ=0,↑> has stationary phase and therefore contributes to the intensity at the center of a pupil in far field diffraction.

For this simulation we corrected the asymmetric aberrations by multipolar elements but still obtained a strong phase oscillation beyond a radius dependent of the size of the field that must be further corrected to obtain maximal efficiency.

Fig 2 a,b is an example of phase hologram described in its thickness map and overall pattern. This pattern reaches an efficiency of 40%. In fig 2c an example of Bessel beam with ℓ=2 is shown. These beams, in the diffraction plane (see fig d), transform to narrow rings. This strongly reduce the demand of lateral stability of the fields and the problems of phase oscillations described in fig. 1

[1] E. Karimi et al. Phys Rev. Lett 108, 044801 (2012)
[2] J. Verbeeck et al Nature 467, 301 (2010).
[3] E. Karimi et al Ultramicroscopy 138, 22 (2014)
[4] V. Grillo et al. Phys. Rev. X 4, 011013 (2014)


Fig. 1: Wavefunction after a Wien filter for a beam at 15KeV. The initial beam had ℓ=1 and 2 spin states were considered. The final spin state are also separately plotted. The external phase oscillation are due to residual aberrations.

Fig. 2: Example of phase hologram: the thickness profile in Si3N4 a) and the full pattern b)are shown. Example of Bessel beam with ℓ=2 in the Fresnel c) and Fraunhofer d) regime

Type of presentation: Poster

IT-1-P-2578 Design of a monochromator for aberration-corrected low-voltage (S)TEM

Mukai M.1, Omoto K.1, Sasaki T.1, Kohno Y.1, Morishita S.1, Kimura A.1, Ikeda A.1, Somehara K.1, Sawada H.1, Kimoto K.2, Suenaga K.3
1JEOL Ltd., Akishima, Tokyo, Japan, 2National Institute for Material Science (NIMS), Tsukuba, Ibaraki, Japan, 3National Institute of Advanced Industrial Science and Technology (AIST), Tsukuba, Ibaraki, Japan
mmukai@jeol.co.jp

Low-voltage analytical electron microscope equipped with delta-type aberration correctors for image- and probe-forming lens system [1] was developed under a project “Triple-C phase-1” to study the atomic structures of carbon materials sensitive to the damage by irradiation of electrons. It enabled us to reveal the characters of graphenes by EELS [2] and to visualize and specify an encapsulated single metal atom in a fullerene [3]. However this microscope was equipped with a cold field emission gun to obtain high brightness therefore its energy resolution remains at approximately 0.3 eV.
For the next challenges, we have started to develop a new type of low-voltage aberration-corrected analytical electron microscope equipped with a monochromator working at 15-60 kV under a project “Triple-C phase-2”, whose targeted energy resolution is better than 25 meV. Fig. 1(a) and 1(b) show an appearance of the microscope and a configuration of components inside the cover.
The developed monochromator employs a double Wien-filter system, arranged between the extraction anode of Schottky source and the accelerator, which is similar configuration to previous design [4]. The electron trajectories from the electron source to the plane of the exit crossover of the monochromator are calculated as shown in Fig. 2. Electron trajectories are set to be symmetric to the plane of energy-selection slit so that the energy-dispersion formed by the first Wien-filter at a slit plane is cancelled by the second Wien-filter at an exit plane as a consequence of the double Wien-filter system. Thus, after the monochromator, the electron probe is achromatic and the energy spread is controllable by choosing the width of the slit, independently on the probe size. In addition, the setting of the monochromator and the electron trajectories inside the monochromator are independent of the change of the accelerating voltage since the accelerator of the electron gun is located after the monochromator and the potential along the optical axis inside the monochromator is kept constant.
We intend to evaluate the performances of the developed low-voltage monochromated electron optical system and the enhancement of spatial resolution arising from a small chromatic aberration in TEM at low accelerating voltage with large scattering cross-section and small specimen damage by reducing a primary electron energy.

References
[1] H. Sawada, et al.: J. Electron. Microsc. 58 (2009) 341.
[2] K. Suenaga and M. Koshino, Nature 468 (2010) 1088.
[3] K. Suenaga, et al.: Nature chemistry 1 (2009) 415.
[4] M. Mukai, et al.: Ultramicroscopy (2014) accepted.


This work is supported by Japan Science and Technology agency, Research Acceleration Program.

Fig. 1: Computer graphics of a low-voltage aberration-corrected analytical electron microscope equipped with the developed monochromator, (a) An appearance with the cover and (b) a column of the microscope inside the cover.

Fig. 2: (a) Calculated trajectories along optical axis from source to exit of monochromator, (b) beam shapes at slit plane with an energy-dispersed 1st focus and (c) beam shapes at exit plane with an achromatic 2nd focus. The red lines and the green lines show the trajectories having different energies of 1 eV inside of the monochromator.

Type of presentation: Poster

IT-1-P-2604 Holographic generation of Electron quasi-Bessel beams

Frabboni S.1,2, Grillo V.2,3, Karimi E.4, Balboni R.5, Gazzadi G. C.2, Mafakheri E.1,2, Boyd R. W.4,6
1Dipartimento FIM, Università di Modena e Reggio Emilia, Via G. Campi 213/A, 41125 Modena, Italy , 2CNR-Istituto Nanoscienze, Centro S3, Via G Campi 213/a, I-41125 Modena, Italy, 3CNR-IMEM, Parco delle Scienze 37a, I-43100 Parma, Italy. , 4Department of Physics, University of Ottawa, 150 Louis Pasteur, Ottawa, Ontario K1N 6N5, Canada, 5CNR-IMM Bologna, Via P. Gobetti 101, 40129 Bologna, Italy, 6Institute of Optics, University of Rochester, Rochester, New York 14627, USA
stefano.frabboni@unimore.it

Recently the attention of electron microscopy community has been attracted by the generation of electron beams by means of holographic element that allows to shape the electron wavefront through a modulation of the phase or amplitude transmittance. This new degree of freedom has already demonstrated huge potentialities in application with electron vortex beams [1]. In this contribution we discuss the case of the quasi-Bessel beams obtained as a coherent superposition of conical plane waves along a closed ring of finite angular aperture [2].
Fig 1a shows the simulated transverse distribution of the electron Bessel beam at the first order of diffraction propagating, in the Fresnel region, from the hologram shown in b). In Fig 1c is reported the scanning electron microscope image of the nanofabricated phase hologram with a zoom-in image of the central region shown in the upper inset. The hologram is obtained from of a FIB-milled silicon nitride membrane, which is almost transparent to the 200keV electron beam [3]. Different depths modify the local projected potential; thus, electrons see different effective paths at grooves. In Fig 1d the distribution of the diffracted electrons in the Fraunhofer region of propagation, is reported. In the first order of diffraction, the Bessel beam forms a ring in the far-field. Due to the limited number of grooves of the hologram, the ring, typical of the Bessel beam, is convoluted with the Airy function of the hologram aperture, thus forming a quasi-Bessel beam. In Fig 1e is shown the measured transverse intensity distribution of the quasi-Bessel beam of the zeroth order generated by the hologram shown in Fig 1c, in Fresnel regime. In Fig 1f the experimental radial intensity distribution of the Bessel beam, blue solid curve, is compared with simulations by varying the convergence of the beam incident on the hologram plane, thus showing the effect of the partial coherence on the Fresnel ring contrast.
Bessel beams have many interesting properties, namely resistance to diffraction and the smallest spot diameter compared to other ordinary type of beams that could be exploited in STEM tomography. In Fig 2 is reported the diffraction free range of the quasi Bessel beam shown in Fig 1c.


[1] J. Verbeeck, H. Tian, and P. Schattschneider, Nature 467, 301 (2010).
[2] V. Grillo, E. Karimi et al. Phys. Rev. X 4, 011013 (2014)
[3]V. Grillo, G.C. Gazzadi, E. Karimi et al. Appl.Phys. Lett. 104, 043109 (2014)


M.E. acknowledges  the support  of SPINNER 2013.

Fig. 1: Computer generated hologram and electron Bessel beams of the zeroth order.

Fig. 2: Propagation of Bessel beams of the zeroth order in the Fresnel regime.

Type of presentation: Poster

IT-1-P-2685 Low Voltage Mini TEM

Coufalová E.1, Mynář M.1, Štěpán P.1, Drštička M.1, Sintorn I. M.2, 3
1DELONG INSTRUMENTS a.s., Brno, Czech Republic, 2Centre for Image Analysis, Uppsala University, Sweden, 3Vironova AB, Stockholm, Sweden
michal.drsticka@dicomps.com

On the basis of experience with the low voltage transmission electron microscopy at 5 kV, which is intended for the study of samples with low contrast (organic matters), we tried to design a TEM optimized in many aspects:
1) Maintaining relatively low voltage to keep up high contrast.
2) The use of such energy, which would open the possibility to increase the resolution of the system to the area of atomic (molecular) resolution using the monochromatization of the primary beam and Cs correction in future.
3) Practical standpoints – reasonable dimensions, resistance to external influences.
4) Energy sufficient for the transmissivity of electrons through samples of "standard" thickness.
It turned out to be suitable to base such electron-optical system on the use of magnetostatic (the objective lens) and electrostatic (projection system) elements. For the above reasons, we have chosen a range of energy of 10-25 keV. This choice enables to maintain the concept of combination of electron-optical and light-optical magnification, which leads to a significant reduction of the dimensions of the unit and solving simultaneously the problem of TEM image digitalization. It emerged that the working energy of 25 keV is the highest possible energy, at which there is no degradation of the applicable high light-optical magnification due to scattering in the single crystal fluorescent screen.
Using light lenses with large numerical aperture (up to 0.95), we achieve a high collection efficiency of the light from the screen. Also, the level of the light signal is high enough at 25keV energy. We have verified that the electron-optical system can be operated in several modes:
1) TEM at 25 keV
2) STEM at 15, 10 keV
3) DIFF at 25keV
The first experimental results confirm the assumptions obtained by electron-optical simulations, in particular the expected resolution in various modes.
It is further confirmed that the contrast inevitably decreases at the energy of 25 keV compared to the lower energies, however, it is still significantly higher than in the energy area of above 50 keV. Even thin sections for which there is no significant increase of chromatic aberration provide sufficient contrast in the image at this working energy. This brings the opportunity to study both stained and unstained samples at low radiation damage.
This version has been optimized for identification of viruses – samples prepared with negative staining and fixation. It allows mobility of the device, and is equipped with user friendly control system with a simple concept that provides remote control resources to allow to be controlled by upper level image analysis software for automatic virus recognition (Kylberg and Sintorn EURASIP J. on Image and Video Processing 2013, 2013:17).


The work has been supported by Eurostars Programme of EUREKA and European Community.

Fig. 1: The body on MiniTEM on a standard desk

Fig. 2: Section of the column

Fig. 3: ATCC and rota viruses stained with 2%Uac in TEM mode at 25 keV

Fig. 4: ATCC and rota viruses stained with 2%Uac in STEM mode at 10 keV

Type of presentation: Poster

IT-1-P-2688 Energy analyzer for point electron sources

Kolařík V.1, Coufalová E.1, Mynář M.1, Drštička M.1
1DELONG INSTRUMENTS a.s., Brno, Czech Republic
michal.drsticka@dicomps.com

We have built an energy analyzer for characterization of parameters of various types of point emitters, electron guns, and illumination blocks of electron columns. It can be also used for characterization of electron monochromators, and for studying the influence of electron – electron interaction on the beam energy spread.
The concept of the analyzer is very simple and physically straight, based on dispersion characteristics of magnetic prism: It is configured for measuring energetic spread of emitters with the virtual source size between 1 nm and 50 nm independently of the electron source distance, it means any design or type of electron gun can be measured.
The theoretic resolution of the analyzer is:
• < 15 mV for the virtual source size of 50 nm,
• < 3 mV for the virtual source size of 15 nm.
The image of virtual source is focused only in the dispersion direction (see Fig. 2). The dispersion of the magnetic prism in this plane is about 3 µm/V at the output edge of the prism. The optical set guarantees the resolution of electron spectrometer on the level of 10 mV or better, the use of slit aperture provides the capability of statistical evaluation of 2048 spectra (pixel columns).
Although the dispersion itself is relatively small (units of µm/V), the analysis is possible at the level of units of mV, because the source image size in the spectral plane is in units of nm. The dispersion plane can be enlarged electron-optically so that it is projected onto a screen with the size accessible for imaging by high-quality light optics (the dispersion and source image are magnified in the same proportion).
The significant input parameters that determine the resulting energy resolution are the virtual source size and used aperture angle. We illustrate on the chart that the effect of the virtual source size for cold field emission and Schottky cathode is in a significant range of aperture size under the resolution of a light objective lens with NA as high as 0.95 in this arrangement.
The high energy resolution of the electron-optical part can be used for very effective monochromatization of en electron beam.
Reference: V. Kolařík, M. Maňkoš, L. H. Veneklasen, Close packed prism arrays for electron microscopy, Optik 87, No.1(1991)


The work was supported by ”Electron Microscopy“ Competence Centre of Technology Agency of the Czech Republic

Fig. 1: Section along the optical axes

Fig. 2: Optical scheme

Fig. 3: The influence of the aperture on the energy resolution for CFE and Schottky emitter at dispersion of 2.8 µm/V in relation to the optical limit (at NA = 0.95, M=40×, pixel size = 7.4 µm)

Fig. 4: Examples: calibration of measurement, profiles - CFE energy spread – 0.34 eV, Schottky emitter energy spread – 0.58 eV

Type of presentation: Poster

IT-1-P-2710 Spiral phase plates for electron vortices

Béché A.1, Winkler R.2, Planck H.3, Hofer F.3, Verbeeck J.1
1EMAT, University of Antwerp, Groenenborgerlaan 171, 2020 Antwerp, Belgium, 2Center for Electron Microscopy, Steyrergasse 17, 8010 Graz, Austria, 3Institute for Electron Microscopy and Nanoanalysis, Graz University of Technology, Steyrergasse 17, A-8010 Graz, Austria
armand.beche@uantwerpen.be

Vortex beams have been recently developed in electron optics and generate a lot of interest due to their potential ability of retrieving magnetic information down to the atomic scale [1, 2]. Several techniques are now available to produce such beams like the holographic mask [2] or the more recent magnetic needle [3]. In this work we propose to extend the idea of Uchida & Tonomura [1] by creating a spiral phase plate with smoothly increasing thickness.
The phase plate should be composed of a light material to prevent too much absorption from the plate itself and be ideally thicker than 100 nm at its highest point to allow a smooth increase of thickness. Focused electron beam induced deposition (FEBID) is an ideal tool to realize such structures as it can deposit functional materials with high spatial resolution. In the present case, ultrathin silicon nitride (SiN) was successfully used as substrate to fabricate SiO2 spiral phase plates as shown in Fig. 1. In order to prevent unwanted scattering from the central hole in the spiral, it was filled with a small amount of platinum via FEBID.
The phase plate was then introduced into the Qu-Ant-TEM, an FEI Titan3 transmission electron microscope, operated in Lorentz mode, to achieve a large field of view with extended spatial coherence conditions. Carefully illuminating the phase plate with a uniform electron beam and looking in the far field, typical features of vortex beams were recorded. Fig. 2 displays a through focus series of the resulting beam which reveals the presence of a doughnut like intensity pattern with the destructive interference centre of the vortex beam.
In order to quantify the orbital angular momentum (OAM) carried by the outgoing beam, electron holography was performed at the edge of the phase plate. By measuring the phase shift between the thickest and thinnest area, the total OAM was estimated to be 0.6 (Fig. 3).
Further tuning of this setup provides another method for creating atomic sized electron vortex beams with the advantage of providing a single vortex beam that is easy to obtain in a standard TEM.

[1] Uchida M. & Tonomura A., Nature Letters (2010), 464, p737-739.
[2] Verbeeck J. et al., Nature Letters (2010), 467, p.301-304.
[3] Béché A. et al., Nature Physics (2014), 10, p. 26-29.


This work was financially supported by the European Research Council under the 7th Framework Program (FP7), ERC grant 246791 COUNTATOMS, ERC Starting Grant 278510 VORTEX and Integrated Infrastructure Initiative No. 312483 ESTEEM2.

Fig. 1: (a) TEM view of the spiral phase plate with the central hole filled with platinum. (b) Atomic force microscope image revealing the thickness profile of the phase plate.

Fig. 2: Far field through focus series of an electron beam evenly illuminating the phase plate. The destructive interference area in the middle of the fully condensed beam (black hole) is typical of a vortex beam.

Fig. 3: (a) Phase image of the edge of the phase plate, between the thickest and thinnest area of the spiral, acquired by holography. (b) Phase profile taken along the dotted area displayed in (a) revealing a total phase shift of 4 rad corresponding to a total OAM of ~0.6.

Type of presentation: Poster

IT-1-P-2835 Projector lens and CCD camera distortions in a Hitachi HF-3300 TEM

Denneulin T.1, Gatel C.1, Houdellier F.1, Hÿtch M. J.1
1CEMES CNRS, 29 rue Jeanne Marvig, 31055 Toulouse, France.
hytch@cemes.fr

In a transmission electron microscope (TEM) the projector lenses are known to introduce large-scale distortions. The magnification and the rotation in the image can vary up to 5% and 2° across the field of view [1]. Therefore the accurate mapping of any physical field (strain, magnetic or electric field) using high resolution TEM or holography requires the calibration of those distortions. The method used here does not add noise to the phase image and alleviates the need for a reference hologram.
We have investigated the projector and the CCD camera distortions on a recently installed aberration-corrected HF-3300 Hitachi TEM (I2TEM-Toulouse). The distortions were measured using off-axis electron holograms acquired in the vacuum. A double biprisms setup was used to remove the Fresnel fringes [2]. The voltages of the biprisms were set so that the interference pattern fills entirely the 4k Gatan CCD camera. Two holograms with a different orientation of the biprisms were acquired in order to reconstruct the 2D strain field using geometrical phase analysis (GPA) [3]. Before GPA calculation, the reconstructed phase images were fitted using a 4th order polynomial to remove the noise.
The influence of the magnification and the values of P1 and P2 was investigated. It was found that the distortions are mainly dependent on the value of P2. Fig. 1 shows the strain field obtained for 4 different values of P2. Increasing P2 is equivalent to “zoom” into the distortion pattern. The variations across the image are then lower for high values of P2. At a nominal magnification of ×1.5M (P2 is equal to 5.3 A) the mean dilatation Δxz varies from 0 to 3% and the rigid body rotation ωxz varies from 0 to 1° from the center to the corner of the image.
According to the theory [1], Δxz and ωxz should be circular shaped. However it can be noted that the rotation image is slightly triangular shaped. After analysing the ronchigram of the camera provided by Gatan [4] we found that this is due to the low frequency distortions of the camera (see Fig. 2 after correction of the camera distortions). We then created an artificial ronchigram for correcting both the projector and the camera distortions. The procedure will be detailed during the presentation. Fig. 3(a) is an example of dark-field hologram acquired on a SiGe layer grown by epitaxy on a Si substrate. Without correction (Fig. 3(b)) the reconstructed phase image exhibits some variations in the substrate and the phase ramps in the layer are slightly distorted. Those artifacts are removed after correction (Fig. 3(c)).

[1] F Hüe et al, J. Electron. Microsc. 54(3) (2005), 181–190
[2] K Harada et al, Appl. Phys. Lett. 84(17) (2004), 3229–3231
[3] MJ Hÿtch et al, Ultramicrosopy 74 (1998), 131–146
[4] P Mooney, private communication


This work received financial support from the European Union under the Seventh Framework Programme under a contract for an Integrated Infrastructure Initiative Reference 312483-ESTEEM2 and the European Metrology Research Programme (EMRP) Project IND54 Nanostrain. The EMRP is jointly funded by the EMRP participating countries within EURAMET and the European Union.

Fig. 1: Distortions of the projector lenses as a function of the displayed value of P2. The strain field was calculated by geometrical phase analysis after fitting the phase images reconstructed from the holograms. From left to right are shown the horizontal εxx, the vertical εzz, the shear εxz strain, the mean dilatation Δxz and the rotation ωxz.

Fig. 2: Distortions (for P2 = 6.0 A) obtained after correcting the phase images with the CCD camera ronchigram.

Fig. 3: (a) (004) dark-field electron hologram of a SiGe layer epitaxially grown on a Si substrate. (b) Phase image reconstructed from the hologram without correction. (c) Phase image reconstructed with correction of the projector and camera distortions.

Type of presentation: Poster

IT-1-P-2975 Design and Characterization of a Single-Atom Electron Column

Lin C. Y.1,2, Chang W. T.1, Hsu W. H.1,3, Lai W. C.1,2, Chen Y. S.1, Hwang I. S.1
1Institute of Physics, Academia Sinica, Taipei, Taiwan, 2Department of Physics, National Taiwan University, Taipei, Taiwan , 3Department of Materials Science and Engineering, National Tsing Hua University, Hsinchu, Taiwan
cylin@phys.sinica.edu.tw

        It has been shown that noble-metal covered W(111) single atom tips (SATs) can be reliably prepared [1,2]. We have demonstrated full spatial coherence of electron beams emitted from the SATs [3]. Thus, single atom electron sources are suitable for phase retrieval imaging methods, such as holography and coherent diffractive imaging. We have proposed a SAT-based low-keV electron microscope that allows different imaging modes, as shown in Fig. 1. For this purpose, we plan to build an electron column with the capability to accelerate electron beams to 1~5 keV and a focused beam spot smaller than 100 nm. The column is composed of two parts: an electron gun and a condenser lens.

        The electron gun consists of a SAT, an extractor/suppressor, and an acceleration electrode. The tip is mounted on a holder that can be translated, tilted, and rotated in nanometer scale by piezo-positioners. Therefore, the tip-lens alignment can be done in vacuum without alignment coils. We have recorded the opening angles of the electron beams. As shown in the inset of Fig. 2, the emitter can be moved to different positions with the piezo-positioners and the corresponding beam profiles are recorded. Fig. 2 shows the half opening angles of the beams at an electron energy of 2.5 keV measured at different extraction voltage and different separations. Clearly the beam opening angle varies with the tip position. When the tip is positioned at about -2.5 mm, the half opening angle can be smaller than 1 mrad. We also find that the suppressor design that is often used in normal field emitters is not effective in reducing the beam divergence for the SAT emitter.

        The condenser lens consists of a limiting aperture, an einzel lens, and an octupole stigmator. We used Simion 8.1 software to simulate the lens parameters and determine the aperture diameter. In our simulations at the electron energy of 2.5 keV and the working distance of 2 mm, a spot size of 140 nm is obtained when the limiting aperture of 100 μm is used; a spot size of 20 nm is obtained when the limiting aperture of 20 μm is used. Fig. 3(a) shows the whole assembly of our instrument. As shown in Fig. 3(b), we have obtained a diffraction pattern on a small region of a suspended CVD graphene, which show two domains with different orientations. We are also designing a microcolumn based on the MEMS technique. Our ultimate goal is to determine the atomic structures of few-layer two-dimensional structures such as graphene and one-dimensional structures such as carbon nanotubes and bio-molecules.

References

[1] H. S. Kuo et al, Nano Lett. 4(12) (2004), p. 2379.

[2] H. S. Kuo et al, J. J. Appl. Phys. 45 (2006), p. 8972.

[3] C. C. Chang et al, Nanotech. 20 (2009), p. 115401.


This work is supported by National Science Council of ROC and Academia Sinica.

Fig. 1: Schematic of a multi-mode low-keV electron microscope

Fig. 2: Beam divergence, measured with the half opening angle at an electron energy of 2.5 keV, versus the extractor/suppressor voltage at different emitter positions. The inset is the schematic for characterization of the beam profile at different emitter positions.

Fig. 3: (a) Illustration and photo of a low-keV electron microscope (b) The diffraction pattern of a CVD graphene sample

Type of presentation: Poster

IT-1-P-2984 Wire corrector for aberration corrected electron optics

Nishi R.1, Ito H.2, Hoque S.2
1Osaka University, Osaka, Japan, 2Hitachi High-Technologies Corporation, Ibaraki, Japan
rnishi@uhvem.osaka-u.ac.jp

The wire corrector on the analogy of multipole correctors was proposed by H. Ito [1]. Two-parallel line current (Fig. 1) makes the magnetic field similar to that of a quadrupole as shown in Fig. 2. When using two-parallel line current, the filed cancels at the rotation symmetric axis and the two-parallel line current can generate quadrupole magnetic field because each magnetic field has opposite rotation direction of magnetic flux.
The wire corrector is only arranged by parallel line currents without using any magnetic materials, so it can be easily and simply fabricated and arranged in comparison to a conventional multipole. Adverse effect of hysteresis of magnetic material does not exist and homogeneity of magnet property is not needed. Magnetic field can be controlled by superimposition of parallel line currents. In actual layout, the wire corrector is configured to a coil shape in addition to the parallel currents with infinite length, but the effect of a coil shape can be reduced by consideration of its shape. Applying constant current to a main coil, fine adjustment of magnetic field can be performed by applying current to a sub coil. The wire corrector is valuable to the aberration corrected electron optics with high precision alignment and reproducibility.
When using the wire corrector of N=2, the magnetic field is similar to quadrupole field but the magnetic field is expanded in a series which also contains octapole field as a higher order term, as shown in Eq.(1) inset of Fig.1. Due to the wire corrector has octapole component, the wire corrector has possibility of simultaneous correction of spherical aberration in addition to chromatic aberration. Symmetric curved ray optical system constituted by combining both components of a deflector and the wire corrector of N=2, is expected that chromatic and spherical aberration is potentially corrected in such configuration.
The combination of the round lenses and the wire correctors of N=3 decreases the spherical aberration [2]. This shows the wire correctors of N=3 worked as a hexapole. The wire corrector has a potential of consisting an easy-to-use aberration corrector.

[1] Hiroyuki Ito et al, USP 7,872,240 B2 (date of patent: Jan. 18, 2011).
[2] H. Rose, Optik, 85 (1990) 19.


A part of this work of calculation was done by Dr. Eric Munro and Dr. John Rouse in Munro's Electron Beam Software Ltd.

Fig. 1: The wire corrector consisting of two parallel line currents (N=2).

Fig. 2: Magnetic flux in the wire corrector (N=2).

Type of presentation: Poster

IT-1-P-3022 Measuring the Orbital Angular Momentum of Electron Vortex Beams by Forked Grating

Saitoh K.1, Hasegawa Y.2, Hirakawa K.2, Tanaka N.1, Uchida M.3
1EcoTopia Science Institute, Nagoya University, 2Department of Crystalline Materials Science, Nagoya University, 3Advanced Science Research Laboratory, Saitama Institute of Technology
saitoh@esi.nagoya-u.ac.jp

After the first report of the production of an electron vortex beam, an electron traveling in free space with orbital angular momentum (OAM) [1], electron vortex beams have been attracting a great attention owing to the unique physical property and application to a new microscopy in materials science [2]. In the present paper, we show the how the electron vortex beams are diffracted by forked gratings and how the OAMs of the electron vortex beams are transferred to each of the diffracted waves (Fig.1(a)). [3].

Figures 1(b) and 1(c) show a schematic diagram of the experimental setup of the present study. The binary masks of the spiral zone plates [Fig. 1(d)] and the forked gratings [Fig. 1(e)], fabricated from 200 nm thick PtPd films using a focused-ion-beam instrument (Hitachi FB-2100). The spiral zone plates and forked gratings were inserted into the condenser lens aperture position and selected-area aperture position, respectively, of a transmission electron microscope (JEOL JEM-2100F), which was operated at an acceleration voltage of 200 keV.

Figures 2(a) and 2(b) show electron vortex beams with OAMs of 10h and -10h, respectively, produced by the spiral zone plate. Each of the electron beams show a ring composed of 10 peaks at the center [4]. Figure 2(c) shows an electron diffraction pattern for an incident electron vortex beam with m = 10h. The diffraction pattern shows a series of diffracted rings, as indicated by the arrows. The central ring, composed of 10 peaks, is the transmitted beam with m = 10h. The 1st- and -1st-order diffracted electron beams show similar ringlike features, but have 11 and 9 peaks, respectively. This indicates that the electron OAMs of the 1st- and -1st-order diffracted beams are 11h and 9h, respectively. Figure 2(d) shows an electron diffraction pattern for an incident electron vortex beam with m = -10h. The pattern shows a series of diffracted rings as in Fig. 2(c), but is horizontally inverted from that shown in Fig. 2(c). The transmitted (0th-), 1st-, and -1st-order diffracted rings show 10, 9, and 11 peaks, respectively, indicating that the electron OAM of the 1st- and -1st-order beams are -9h and -11h, respectively. Our results indicate that the forked grating with a Burgers vector of b = 1 transfers not only linear momentum but also OAM, where the electron OAM transfer of the nth-order diffracted electron beam is nh. This diffraction property could be used as an electron OAM analyzer, as the nth-order diffracted beam shows a normal peak.

References

[1] M. Uchida and A. Tonomura, Nature 464, 737 (2010).

[2] J. Verbeeck, H. Tian, and P. Schattschneider, Nature 467, 301 (2010).

[3] K. Saitoh et al, Phys. Rev. Lett. 111, 074801 (2013).

[4] K. Saitoh et al., J. Electron Microsc. 61, 171 (2012).


The present work was partly supported by the Grant-in-Aid for Scientific Research (A) (No. 23241036), the Ministry of Education, Culture, Sports, Science and Technology, Japan, and the Mitsubishi Foundation.

Fig. 1: (a) Schematic drawing of the present experiment. (b),(c) Ray-path diagrams of the experimental setups for incident electron vortex beams of m = 10h (b) and m = 10h (c). (d) A spiral zone plate with a 20 μm diameter introduced to the condenser lens system (e) A forked grating with a 30 μm diameter introduced to the image-forming lens system.

Fig. 2: Incident electron vortex beams with m = 10h (a) and m = -10h (b), and diffraction patterns of the vortex beams with m = 10h (c) and m = -10h (d) generated by the forked grating shown in Fig. 1(e). 

Type of presentation: Poster

IT-1-P-3222 Contrast enhancement in TEM imaging by use of a central beam stop

Zandbergen H.1, Xu Q.1
1Kavli Institute of Nanoscience, Delft University of Technology, Delft, The Netherlands
h.w.zandbergen@tudelft.nl

In TEM in life science, beam damage is the most important limitation. This is also the case for materials science samples like graphene, polymers and hybrid materials. On the imaging side an important boost is expected from the introduction of a phase plate. Phase plates have been researched over several decades and no easy to use system has emerged yet, indicating that is not easy task. Given the importance of efficient imaging, it is clearly necessary to explore other routes. We have explored [1] the possibility of dark-field imaging for contrast enhancement in which we have tried to block the central beam [2] and leave as many of the diffraction beams un-blocked.
Central beam block apertures (the abbreviation DF-000 is used in this abstract) in the shape of Mercedes star (see Figure 1) were made with a FIB. In our experiments we have observed no sign of charging, possibly due to the DF-000 shape. In central disk should preferably be smaller than the frequency, g, one wants to observe, which is of course much smaller for biological samples than for most materials samples. Our DF-000 removed frequencies corresponding to d-spacings of 8.7 Å and larger. In the presentation we will report how far we can decrease the size of central disk without charging problems and with still good blocking of the central beam.
For the drilling of holes in exfoliated graphene without contamination build-up we heated the graphene to 600°C. TEM experiments were done at 300keV and post-specimen aberration correction at 600°C. Figure 1c and 1d shows high-resolution and DF-000 images of multilayer graphene (4-5 layers). A hole was made in this sample using an e-beam. This hole can be seen very well in the DF-000 image and only faintly in the BF image. In both cases one can see that the graphene lattice continues up to the edge of the hole. The gradual decrease in thickness is clearly visible in the DF-000 image and not at all in the BF image. Thus we can obtain in the DF-000 image high-resolution information with a similar resolution limit as the BF image.
Figure 2 shows several images of graphene with three holes with varying size imaged at various focus values, showing that the bright field images in the range from -1500 to + 1500 nm show hardly any contrast and none at zero focus. On the contrary, the contrast in the DF-000 taken at 0 focus shows the largest contrast and in particular the smallest delocalization. In this case selecting ~ zero focus is easy by minimizing the blurring in the image.

1. Zhang C, Xu, Q, Peters PJ, and Zandbergen, H, Ultramicroscopy 134, 200 (2013)
2. Cowley, J., Acta Crystallographica Section A 1973, 29, 529-536


Fig. 1: (a) SEM image of the DF-000 objective aperture used to stop the central beam. (b) shows a typical DF-000 aperture in diffraction space. (c) and (d) show HREM images taken without an aperture and with the DF-000 image taken from the exactly the same area.

Fig. 2: Images from the same area of single layer graphene with three holes showing the effect of focus in BF and in DF-000 modes. No DF-000 at defocus values of -1000 and -3000 nm are given because these are too blurred. The two small holes are only barely visible in the BF image taken at -3000 nm and not at all in the other three BF images.

Type of presentation: Poster

IT-1-P-6004 Determinationof geometrical form factor of emitter from Schottky plot

Emura Y.1, Murata H.1, Rokuta E.1, Shimoyama H.1, Yasuda H.2, Haraguchi T.2
1Faculty of Science and Technology, Meijo University, 2PARAM Corporation
hkmurata@meijo-u.ac.jp

In this paper we report preliminary experimental results on a LaB6 Schottky emission electron gun, which also includes our new findings that the electric field strength on the emitter surface can be estimated experimentally from the Schottky plot whose slope depends not on the work function but only on the reciprocal of the emitter temperature. According to the theoretical considerations on the Schottky emission, if the values of log10 j (j: emission current density) are plotted as a function of √F (F: field strength on the emitter surface), then the graph becomes a straight line with the slope of 1.913/T (T: emitter temperature), which is known as “the theoretical Schottky plot”. In experiment, on the other hand, the beam current I is measured as a function of the extraction voltage Va. Thus, the slope of “the experimental Schottky plot” is different from that of “the theoretical Schottky plot”. From I = j × ΔS (ΔS: emission area on emitter surface), the vertical axis of “the experimental Schottky plot” is expressed as log10 j + log10 ΔS, which means the graph is moved parallel to the vertical direction without changing the slope of the graph. We mark a new scale on the horizontal axis of “the experimental Schottky plot” in order that the slope may be equal to 1.913/T. Then, the new horizontal axis should be graduated in √F. This procedure makes it possible to relate the field strength F directly to the extraction voltage Va as F = β Va, where β is the geometrical form factor of the emitter.

The Schottky emission experiment has been done in the ultra-high vacuum chamber, using the experimental circuit shown in Fig. 1. A flat top LaB6 emitter is embedded into a rhenium conical sheath, and is heated by a tungsten hairpin filament, as shown in Fig. 1.

The beam current IF was measured as a function of the extraction voltage Va for a constant emitter temperature T = 1600 K by the Faraday cup placed behind the fluorescent screen. In Fig. 2, the values of log10 IF are plotted as a function of √Va. It can be seen that the plot is almost a straight line, which indicates that the emission is under Schottky emission mode. Figure 2 (a) and (b) also show emission patterns observed on the fluorescent screen at Va = 2 kV and 5 kV, respectively.

Figure 3 shows the above procedures, where the new horizontal axes scaled in √F are placed in addition to the original horizontal axes scaled in √Va. From the relation between F and Va, we have found that β = 99.6 [1/cm] for T = 1600 K. We have also performed the field calculation for the experimental system shown in Fig. 1. According to the calculation, the geometrical form factor β has been found to be β = 95.0 [1/cm], which is in good agreement with that estimated from “the experimental Schottky plot”.


Fig. 1: Experimental circuit for measuring the beam current by Faraday cup.

Fig. 2: Experimental Schottky plot for emitter temperature T = 1600 K. Emission patterns observed on the fluorescentscreen at Va = 2 kV (a) and 5 kV (b).

Fig. 3: Schottky plot for T = 1600 K for determination of geometrical formfactor β. A newhorizontal axis scaled in the square root of F is added so that the slope ofthe Schottky plot may be equal to 1.913/T.

Type of presentation: Poster

IT-1-P-6020 Design and realisation of variable C shaped structured illumination

Mousley M.1, Thirunavukkarasu G.1, Babiker M.1, Yuan J.1
1Department of Physics, University of York, Heslington, York, YO10 5DD, United Kingdom
mgm514@york.ac.uk

Structured illumination is a new development in electron microscopy, with the advantage such as longer column channeling distances in crystals by donut-shaped illumination of atomic scale vortex electron beams [1]. In this paper, we introduce a controlled way to realize C shaped structured illumination. Analytical equations determining the parameters of the C shaped illumination pattern have been derived using phase gradient analysis, allowing independent control of the C-opening angle and radius of the C shape. Experimentally, we have used computer generated hologram (CGH) method to generate C shaped structured illumination in a 200 keV transmission electron microscope. Both amplitude and phase CGH masks have been used and comparisons with simulations show a strong match between the theoretical results and the experimentally recorded electron microscope images. C-shaped illumination has promises in potential applications such as electron beam lithography for production of metamaterials which utilise split ring resonance structures [2]. Physical dimensions of the artificial electromagnetic resonance structures as small as nanometres should now be possible. Furthermore the orientation of the C shape illumination can be readily identified, allowing the easy identification of the Faraday rotational effects of the vortex beams [3].

[1] H. Xin and H Zheng (2012) Microscopy and Microanalysis, Vol. 18, p711-9

[2] D. R. Smith, W. J. Padilla, D. C. Vier, S. C. Nemat-Nasser and S. Schultz (2000) Phys. Rev. Lett. 84, p4184-7.

[3] C. Greenshields, R. Stamps, S. Franke-Arnold (2012) New J Phys. 14, 103040


We wish to thank the UK Engineering and Physical Science Research Council (EPSRC) for financial support to this research by a grant (EP/J022098) and M. Ward of Leeds Electron Microscopy and Spectroscopy Centre, University of Leeds for the help with focused ion beam experiment.

IT-2. High resolution TEM and STEM

Type of presentation: Invited

IT-2-IN-2458 Advanced scanning transmission electron microscopy with segmented annular all field detector

Shibata N.1,2
1The University of Tokyo, 2JST-PRESTO
shibata@sigma.t.u-tokyo.ac.jp

In this talk, I will review our recent and on-going findings from our exploration of new atomic-resolution imaging modes using an area detector which is capable of atomic-resolution STEM imaging [1]. One possibility is atomic-resolution differential phase contrast (DPC) imaging [2]. It has been reported that, to a good approximation, DPC STEM images represent the gradient of the object potential (= fields) taken in the direction of the diagonally opposed detector segments, provided the object scatters weakly [3-5]. Here, we show atomic-resolution DPC STEM images of SrTiO3 observed from the [001] direction [2]. Fig. 1(a) shows the orientation relationship between the SrTiO3 crystal and the detector segments used in this study. The probe-forming aperture angle was 23 mrad and the polar angle range of the detector segments was 15.3 to 30.6 mrad. Fig. 1(b) shows the experimental difference image and its intensity profile projected over the vertical direction in the image. The simultaneous ADF STEM image and its intensity profile are used for reference since the peaks in ADF image are a well-established indicator of the true atomic positions. Fig. 1(c) shows the results of corresponding image simulation. It is clear that the DPC STEM profile has a node (zero crossing) at the atom location. The profile is antisymmetric about this point, reflecting the reversal of the electric field direction across the atom along the direction of diagonal detector segments. Combined with detailed image simulations, atomic-resolution DPC STEM is found to provide information on the local electric field distribution in the vicinity of the atomic columns. Some application results of DPC STEM imaging for ferroelectics and their interfaces will be presented.
Another possibility is annular bright-field (ABF) imaging and its derivatives. Fig. 2(a) shows a schematic of the ABF detector geometry. We form “enhanced” (e)ABF images [6] by simply taking the difference between ABF images and the corresponding BF images using the area detector. As shown in Fig. 2(b), we find that light element imaging can be selectively enhanced by this process. We anticipate that the area detector will offer still further possibilities for new atomic-resolution STEM imaging modes useful for material characterization.

References
[1] N. Shibata et al., J. Electron Microscopy 59, 473 (2010).
[2] N. Shibata et al., Nature Phys., 8, 611 (2012).
[3] N.H. Dekkers and H. de Lang, Optik, 41, 452 (1974).
[4] H. Rose, Ultramicroscopy, 2, 251 (1977).
[5] W.C. Stewart, J. Opt. Soc. Am., 66, 813 (1976).
[6] S.D. Findlay et al., Ultramicroscopy, 136, 31 (2014).


I deeply thank S.D. Findlay and Y. Ikuhara for their collaboration in materials characterization and Y. Kohno, H. Sawada and Y. Kondo for their collaboration in the detector development. This work was supported by the PRESTO, JST. A part of this work was conducted in Research Hub for Advanced Nano Characterization, The University of Tokyo. 

Fig. 1: (a) Schematic illustration showing the relationship between the crystallographic orientation of SrTiO3 and the two detector segments. (b) The DPC STEM image formed by subtracting the signal in detector segment Y from that in detector segment X and its image intensity profile [2]. 

Fig. 2: (a)Schematic illustration of BF and ABF detector geometry. (b)ABF and eABF images of LaTiO3 observed from [001] direction [6].

Type of presentation: Invited

IT-2-IN-2576 Insights into Materials Properties with Quantitative STEM and EELS

Botton G. A.1, Bellido E. P.1, Bugnet M.1, Dudeck K. J.1, Gauquelin N.1, Liu H.1, Prabhudev S.1, Rossouw D.1, Scullion A.1, Stambula S.1, Woo S. Y.1, Zhu G. Z.1
1Department of Materials Science and Engineering, McMaster University, Hamilton, ON, Canada
gbotton@mcmaster.ca

The development of aberration correctors in scanning transmission electron microscopy (STEM) has dramatically improved the analytical “toolkit” of materials scientists. In particular, when combined with electron energy loss spectroscopy (EELS), STEM makes it possible to detect compositional and spectroscopic changes at the atomic level that can be used to understand the structure, and ultimately the performance of materials. Here we present some examples of quantitative STEM and EELS as applied to the study of graphene-based materials, complex nanoparticles used in electrocatalysts, and the defects generated in implanted Si and plasmonic structures.

An FEI Titan microscope was used for this work. With this system, we imaged Pt atoms on multilayer graphene nanosheets (GNS) and demonstrate that single Pt atoms are stabilized during atomic layer deposition on N-doped GNS. Quantitative analyses of images show that the single atoms are located at GNS edge steps and that the doping strongly suppresses the growth of Pt clusters (Figure 1a, b) [1]. Similarly, quantitative images have been used to detect atomic displacements on PtFe intermetallic core-shell nanoparticles that exhibit very high specific activity compared to pure Pt [2,3]. Not only is elemental mapping at the atomic scale possible, but the high beam current and fast spectrometers also allow the acquisition of maps with large sampling of the nanostructure. This is illustrated in the study of PtRu nanocatalysts used in fuel cells where Ru core-Pt shell structures are very clearly mapped (Figure 1c).

Beyond the “simple” deduction of the distribution of elements in nanostructures from maps, quantification is essential to understand the detailed structure of defects and correlate compositional measurement with the optical response of materials. The detailed quantification of the atomic position of a defect, in this case a so-called {311} defect [4] generated by the implantation of ions in Si [4,5] shows that an excellent agreement is obtained between the experimental atomic positions and molecular dynamics simulations (Figure 3) [4] with an accuracy of better than 0.05nm for more than 100 atomic columns. Similarly, quantitative analysis of SiGe alloys has allowed us to deduce compositional fluctuations and interdiffusion in proximity of interfaces [6].

[1] S. Stambula et al., J. Phys. Chem. C, on-line (2014), DOI: 10.1021/jp408979h
[2] S. Prabhudev et al., ACS Nano 7, 6103-6110, (2013)
[3] M.C.Y. Chan et al, Nanoscale 4 (22), 7273-7279, (2012)
[4] K.J. Dudeck et al., Physical Review Letters, 110, 166102 (2013)
[5] K.J. Dudeck et al., Semiconductor Science and Technology, 28, 125012, (2013)
[6] G. Radtke et al., Physical Review B 87, 205309, (2013)


The authors are grateful to NSERC for supporting this research. The microscopy was carried out at the Canadian Centre for Electron Microscopy, a National facility supported by NSERC and McMaster. We are grateful to Paolo Longo (Gatan Inc.) for the help in setting up the Quantum 966 spectrometer.

Fig. 1: HAADF STEM image of single Pt atoms and clusters stabilized on N-doped GNS. Raw signals (a), edges and atoms detected (b). Green arrows point to the few Pt atoms (pink) stabilized on GNS terraces (edges labeled in yellow) [1]. (c) elemental maps of PtRu core-shell nanoparticles

Fig. 2: HAADF STEM image of a {311} defect in Si (a) and the deduced atomic positions (crosses) in (b) with the expected atomic positions deduced by molecular dynamics calculations [4].

Type of presentation: Invited

IT-2-IN-2893 Extending the capabilities of high-resolution STEM:measuring depth dependent strain using optical sectioning and aberration-free phase contrast imaging of low-Z materials

Nellist P. D.1,2, Yang H.1, Lozano J. G.1, Hirsch P. B.1, Pennycook T. J.1,2
1Department of Materials, University of Oxford, UK, 2SuperSTEM Laboratory, Daresbury, UK
peter.nellist@materials.ox.ac.uk

The development of aberration correction in scanning transmission electron microscopy (STEM) has had a major impact on spatial resolution and analytical capability. Unsurprisingly, alongside these developments come further complications but also opportunities. The increased numerical aperture allowed by aberration correction leads to a reduced depth of focus (DOF), which in a modern instrument may be just a few nanometres, and typically less than the sample thickness. The increased numerical aperture of the probe converging optics also leads to a larger bright-field (BF) disc in the detector plane, and as a result much of the scattering by the sample remains in the BF disc. In this presentation we will two explore STEM imaging modes that make use of each of these effects to provide aberration-corrected STEM with new capabilities.


The reduced DOF means that in principle a three-dimensional (3D) data-set can be recorded as a focal series of images. In practice, a confocal configuration is generally required. At atomic resolution, however, nanometre-scale depth resolution is also available in the conventional STEM configuration [1]. For dislocations in GaN viewed end-on we show the detection of depth-dependent Eshelby twist displacements associated with screw dislocations. We also show that ADF STEM optical sectioning can be used to measure the screw displacements parallel to the dislocation line for dislocations lying in the plane of the TEM sample, and we use this effect to measure the dissociation reaction of mixed dislocations in GaN. Despite the channelling of the probe, the depth sensitivity persists, and Fig. 1 shows how a simple weighted potential model is a reasonable approximation to a full channelling simulation.


Use of a pixelated detector to record the entire BF disc in the detector plane as a function of probe position results in a 4D data set. A phase contrast image can be retrieved from this data set using a processing method proposed by Rodenburg et al [2]. Interference between the BF disc and a diffracted disc leads to intensity in the overlap region that oscillates with respect to probe position. Figure 2 shows the magnitude and phase of that oscillation for a bilayer graphene sample. From such data a full phase contrast image can be retrieved and we compare the sensitivity of this imaging mode with alternative techniques such as annular bright-field and differential phase contrast. The data is also an excellent instrument diagnostic, and effects such as aperture charging, residual aberrations and the effect of chromatic aberrations can also be observed.

[1] P.D. Nellist and P. Wang, Annual Review of Materials Research 42 (2012) 125-143.
[2] J.M. Rodenburg, B.C. McCallum and P.D. Nellist, Ultramicroscopy, 48 (1993) 303-314.


This research has received funding from the EPSRC and the EU 7th Framework Programme under Grant Agreement 312483 - ESTEEM2 (Integrated Infrastructure Initiative–I3) and was partly performed at the EPSRC National Facility for Aberration-Corrected STEM.

Fig. 1: Aberration-corrected ADF STEM simulated images along the a lattice direction for a 10 nm thick sample of GaN containing a screw dislocation lying parallel to [0001] in the mid-plane of the foil. For each defocus , the left panel shows a full frozen phonon calculation using the QSTEM code and the right panel a simple weighted potential approach.

Fig. 2: The (a) amplitude and (b) phase of the interference observed in the BF disc for one particular spatial frequency with respect to probe position in the 4D data set. The data was recorded from bilayer graphene at 60 kV using a Nion UltraSTEM 200 with a convergence angle of 30 mrad.

Type of presentation: Oral

IT-2-O-1555 Imaging of light elements by annular dark-field Cs-corrected STEM

Lotnyk A.1, Poppitz D.1, Gerlach J. W.1, Rauschenbach B.1
1Leibniz Institute of Surface Modification (IOM), Leipzig, Germany
andriy.lotnyk@iom-leipzig.de

Nowadays, many crystalline lattices can be imaged directly at atomic resolution in Cs-corrected STEM. Recently, it was shown that light and heavy elements in crystalline lattices can be detected with an ABF method1 or with a double-detector STEM method.2 However, imaging of atomic columns of light elements by ADF method remains challenging. Particularly, the observation of light element columns at the interface between two different materials is still a difficult issue. In this work, we were able to detect directly and simultaneously the N and C atomic columns at the GaN-SiC interface and within the GaN and SiC materials. Additionally, the O atomic columns in a SrTiO3 single crystal were also observed by our method. We have studied the influence of imaging conditions on the appearance of N and C atomic columns in the GaN and SiC materials. The obtained results are discussed and are supported by image simulations.

       The GaN thin film for this study was grown on 6H-SiC(0001) substrate by ion-beam assisted molecular beam epitaxy. STEM experiments were performed on a probe Cs-corrected Titan3 G2 60-300 microscope operated at 300 kV. A probe forming aperture of 20 mrad was used. Cross-sectional samples for STEM work were prepared by FIB technique. To improve the surface quality of the TEM specimens and to reduce the samples thicknesses, a focused low-energy argon ion milling (NanoMill system) was applied.3 Ion energies from 900 eV down to 200 eV were applied to remove implanted Ga ions and amorphous regions caused by the FIB. Image simulations were performed with the xHREM/STEM software package.

       Figures 1 and 2 show the results of our work.4 We found that by adjusting the settings of  HAADF detector and defocus value in STEM, the light element columns at the GaN-SiC interface and within the w-GaN, 6H-SiC and SrTiO3 lattices can be imaged using only a single HAADF detector. We concluded that image simulations for interpretation of atomic-resolution STEM images are only necessary when the probe forming aperture angle overlaps the inner angle of an annular STEM detector or when a complex defect structure is observed in a studied TEM sample. Our method works well using either ADF or HAADF detector, because their angular ranges and defocus values can be easily adjusted on any Cs-corrected STEM. Thus, on TEM systems equipped with only one HAADF detector, the technique can be used without any doubt and upgrades to an ABF detector.

1. S.D. Findlay, N. Shibata, H. Sawada et al., Appl. Phys. Lett. 95, 191913 (2009).
2. Y. Kotaka, Appl. Phys. Lett. 101, 133107 (2012).
3. D. Poppitz, A. Lotnyk, J.W. Gerlach, B. Rauschenbach Acta Mater. 65, 98 (2014).
4. A. Lotnyk, D. Poppitz, J.W. Gerlach, B. Rauschenbach Appl. Phys. Lett. 104, 071908 (2014).


The financial support of the European Union and the Free State of Saxony (LenA project; Project No. 100074065) is greatly acknowledged.

Cs-corrected: aberration-corrected; STEM: scanning transmission electron microscopy; ABF: annular bright-field; ADF: annular dark-filed; HAADF: high-angle ADF; FIB: focused ion beam; w-GaN: wurtzite-type GaN; i: detector inner angle; o: detector outer angle.

Fig. 1: (a) Atomic-resolution STEM image of the GaN-SiC interface taken with a HAADF detector (i20.4-o124.6 mrad) and schematic representation of w-GaN and 6H-SiC lattices along the [2-1-10] zone axis. (b) and (c) Simulated images of w-GaN and 6H-SiC, respectively, at 5 nm underfocus. The TEM sample thickness is measured to be about 16 nm.

Fig. 2: High-resolution STEM images of SrTiO3 acquired with (a) ADF (i19.1-o106.5 mrad) and (b) ABF (i10.1-o19.1 mrad) detectors. The insets in (a) and (b) show the SrTiO3 structure viewed along the [001] zone axis. The TEM sample thickness is measured to be about 60 nm.

Type of presentation: Oral

IT-2-O-1645 Atomically Resolved 3D Shape Determination of an MgO Crystal from a Single HRTEM Image

Jia C. L.1,2,3, Mi S. B.1,4, Barthel J.3,5, Wang D.1, Dunin-Borkowski R. E.2,3, Urban K. W.2,3, Thust A.2,3
1International Center of Dielectric Research, The School of Electronic and Information Engineering, Xi'an Jiaotong University, Xi'an 710049, China, 2Peter Grünberg Institute, Forschungszentrum Jülich GmbH, 52425 Jülich, Germany, 3Ernst Ruska-Centre for Microscopy and Spectroscopy with Electrons, Forschungszentrum Jülich GmbH, 52425 Jülich, Germany, 4Shenyang National Laboratory for Materials Science, Institute of Metal Research, Chinese Academy of Sciences, Shenyang 110016, China, 5Central Facility for Electron Microscopy, RWTH Aachen University, Ahornstr. 55, 52074 Aachen, Germany.
a.thust@fz-juelich.de

High-resolution transmission electron microscopy (HRTEM) allows one to investigate the structure of matter on an atomic level [1]. However, most atomic structure characterizations obtained by HRTEM were so far restricted to the determination of atomic column positions in the image plane perpendicular to the incident electron beam. Due to the fact that the depth resolution of the TEM technique along the beam direction is inferior to its lateral resolution, full 3D structure determinations on an atomic level remain highly challenging. The 3D structure retrieval problem can be solved with tomographic methods, where a multitude of images is acquired from different observation directions. Such multi-image approaches are very demanding at atomic resolution due to instrumental instabilities [2] and due to a possible radiation damage of the object. Alternatively, single-image approaches, where only one exposure is taken along a crystallographic zone axis, have been successfully used to count the number of atoms in crystalline columns running parallel to the beam direction. However, a full 3D determination of the crystal shape would additionally require a highly accurate determination of all column positions along the beam direction, which has not been achieved so far with the single-image approach.

We demonstrate that the full 3D shape of a thin MgO crystal can be determined in a nearly unique way from a single HRTEM image (Fig. 1). Our 3D determination of the crystal shape is based on refining an atomic structure model (Fig. 2) in such a way that a HRTEM image simulated on the basis of this model fits best to the experimental image. In contrast to the usual simplifying assumption of flat lower and upper object surfaces in conjunction with a single global defocus value [3], our structure refinement is executed now locally column-by-column, allowing also for atomically corrugated object surfaces. The comparison between simulation and experiment is made on the basis of absolute image intensity values [4]. A crucial part of our procedure is an extended statistical confidence test which yields detailed quantitative statements on the uniqueness and the reliability of the retrieved 3D crystal shape.

References:

[1] K.W. Urban, Science 321 (2008) 506.
[2] J. Barthel and A. Thust, Ultramicroscopy 134 (2013) 6.
[3] C.-L. Jia et al, Microsc. Microanal. 19 (2013) 310.
[4] A. Thust, Phys. Rev. Lett. 102 (2009) 220801.


Fig. 1: High-resolution image of the edge of an MgO crystal taken along the [001] zone axis with a CS-corrected FEI Titan 80-300 electron microscope at 300 kV accelerating voltage. The 3D shape reconstruction was performed at the area indicated by the dashed box.

Fig. 2: 3D structure model retrieved from the boxed area in Fig. 1. Red balls indicate Mg atoms, yellow balls O atoms, purple balls indicate formally half-occupied Mg positions, and green balls formally half-occupied O positions.

Type of presentation: Oral

IT-2-O-1857 Towards a quantitative exit wave function: the influence of phonon scattering

Liberti E.1, Kim J. S.1, Kirkland A. I.1
1Department of Materials, University of Oxford, Parks Road, Oxford, OX1 3PH, UK
emanuela.liberti@materials.ox.ac.uk

To provide the utmost quantitative information about the atomic structure of the specimen is the ultimate challenge sought by modern high-resolution transmission electron microscopy. At the specimen exit surface, quantitative structural information is embedded in the object complex wave function, which can be recovered, with atomic resolution, from a focal (or tilt) series of aberration corrected HRTEM images [1]. Nonetheless, the quantitative information that is obtained from the exit wave is often in disagreement with imaging simulations. This disagreement is in effect a contrast mismatch, or Stobbs factor, which accounts for a reduction of the experimental image contrast by a factor of three with respect to the calculations [2]. The scattering of phonons following the electron beam-specimen interaction is amongst the possible causes of the Stobbs factor [3].
In this contribution, we discuss the role of phonon scattering in the quantification of the exit wave function of a single layer of graphene. For this idealized object, the contribution of the thermal phonon scattering to the total elastic scattering can be directly investigated by quantifying the exit wave function at different temperatures. For the imaging simulations, the influence of thermal motion upon modeling of the elastic scattering is studied quantitatively, using both the absorptive potential and frozen phonon approaches, addressing the role of the Debye-Waller factor in predicting the thermal displacement of graphene atoms. Experimentally, the exit wave function is recovered in the linear imaging approximation, in both heating and cooling conditions, as well as at room temperature.
To conclude, we present, and discuss, the comparison between the quantitative exit wave functions, obtained in both calculated and experimental approaches.

[1] A.I. Kirkland, S. J. Haigh, Jeol news, 44 (2009) 6 – 11.
[2] M.J. Hÿtch, W.M. Stobbs, Ultramicroscopy 53 (1994) 191 – 203.
[3] A. Howie, Ultramicroscopy 98 (2004) 73 – 79.


The research leading to these results has received funding from the European Union Seventh Framework Programme under Grant Agreement 312483-ESTEEM2 (Integrated Infrastructure Initiative–I3).

Type of presentation: Oral

IT-2-O-1921 Assessment of lower-voltage TEM performance using 3D Fourier transform of through-focus images

Kimoto K.1, Ishizuka K.2
1National Institute for Materials Science, Tsukuba, Japan, 2HREM Research Inc., Higashimatsuyama, Japan
kimoto.koji@nims.go.jp

The performance of an aberration-corrected TEM is determined by the information limit that is often demonstrated using Young's fringe method. However Young's fringe method could show unexpected high frequency information due to the non-linear terms as pointed out by several researchers [1,2]. The three-dimensional (3D) Fourier transform (FT) of through-focus TEM images allows us to discriminate between the linear and the non-linear imaging terms [3,4]. The linear imaging terms are observed on twin Ewald spheres in the 3D FT using an amorphous specimen. Here, we use the 3D FT of through-focus TEM images for the assessment of two low-voltage TEM systems.

Two spherical-aberration-corrected microscopes were assessed and compared. One was a Titan3 (FEI) equipped with a monochromator and a spherical aberration corrector for image forming (CEOS, CETCOR) operated at an acceleration voltage of 80 kV. The energy spread of the electron source was 0.1 eV under monochromated condition. The other microscope, the TripleC microscope, was equipped with a cold field-emission gun (CFEG) and the spherical aberration corrector developed for the TripleC project. This microscope was operated at 60 and 30 kV [5], and the energy spread was 0.3-0.4eV.

Figure 1 schematically shows various 3D data processed in this study [6]. Acquired through-focus TEM images are stacked as a function of the defocus z (Fig. 1a). The 3D Fourier transform Iuvw (Fig. 1c) of through-focus images shows two paraboloids called Ewald spheres, attached at the origin. The information limit can be estimated as an observable range of the Ewald spheres.

The signal of Ewald spheres depends on various factors, such as atomic scattering factors, a specimen structure, thickness, and the modulation transfer function of an imaging device; therefore, the quantitative evaluation of diverse TEM systems is not straightforward. Here we apply the tilted incidence in the 3D Fourier transform method (Fig. 2) to normalize those factors. We evaluate the spatial frequency at which information transfer decreases to 1/e2 (Fig. 3). It was found that the energy spread of the electron source is the major limiting factor even in a monochromated TEM [7].

[1] M. Haider et al., Microsc. Microanal. 16 (2010) 393. [2] J. Barthel, et al., Phys. Rev. Lett. 101 (2008) 200801. [3] Y. Taniguchi, et al., J. Electron Microsc. 40 (1991) 5. [4] M. Op. de Beeck et al., Ultramicrosc. 64 (1996) 167. [5] H. Sawada et al., Ultramicrosc. 110 (2010) 958. [6] K. Kimoto et al., Ultramicrosc. 121 (2012) 31. [7] K. Kimoto et al., Ultramicrosc. 134 (2013) 86.


We thank Drs. Nagai, Freitag, Sawada, Sasaki, Ohwada, Sato and Suenaga for invaluable discussions. This work is supported by Nanotechnology Platform of MEXT and Research Acceleration Program of JSPS.

Fig. 1: Schematics of (a) through-focus TEM images Ixyz, (b) stack of 2D FTs Iuvz, and (c) 3D FT of the through-focus images Iuvw. Since Iuvz and Iuvw are complex, their moduli are shown in gray scale. The cross section Ivz is similar to the Thon diagram. Two Ewald spheres attached at the origin are observed in the 3D Fourier space Iuvw.

Fig. 2: Cross sections of 3D FTs under on-axial and tilted incidence conditions. (a) Titan3 (80kV) and (b) TripleC (60kV).

Fig. 3: Information limit of (a) monochromated Titan3 (80kV), TripleC at 60kV (b) and 30kV (c).

Type of presentation: Oral

IT-2-O-2157 An alternative normalization method for quantitative STEM

Martinez G. T.1, Jones L.2, Béché A.1, Verbeeck J.1, Nellist P. D.2, Van Aert S.1
1Electron Microscopy for Materials Science (EMAT), University of Antwerp, Antwerp, Belgium, 2Department of Materials, University of Oxford, Oxford, United Kingdom
gerardo.martinez@uantwerpen.be

Techniques such as annular bright-field (ABF) or high-angle annular dark-field (HAADF) scanning transmission electron microscopy (STEM) have become widely used in quantitative studies because of the possibility to directly compare experimental and simulated images when the experimental data is expressed in units of ‘fraction of incident probe’ [1]. This is achieved by subtracting by the amplifier’s ‘black-level’ normalizing the experimental image by the mean sensitivity of the annular detector. Since the detector response is spatially inhomogeneous [2], a ‘detector sensitivity’ profile needs to be included in image simulations in order to account for these irregularities. Unfortunately, the quantification procedure now becomes both experiment and instrument specific, with new simulations needing to be carried out for the specific response of each instrument’s detector. This not only impedes the comparison between different instruments but can also be computationally very time consuming.

In this work, we propose an alternative method for normalizing experimental data in order to compare these with simulations that consider a homogenous detector response. To achieve this, we determine the electron flux distribution reaching the detector by means of a camera length series, which is then used to determine the corresponding weighting of the detector response. Figure 1a) shows the detector scan and b) its corresponding active area. The electron flux reaching the active area of the detector is shown in Figure 1 c), which was determined using a camera length series (Figure 2). Next, after normalizing this flux profile to unity, it is multiplied pixel-wise with the experimental detector map, Figure 1d), in which the detector response inhomogeneity is clearly observed. By integrating Figure 1d), we obtain an overall ‘flux-weighted detector sensitivity’ value, which can be used for the experimental data normalization. To validate the proposed methodology, we simulated a [100] oriented Pt crystal using the StemSim software under the frozen lattice approach [3]. The simulations considered homogeneous and inhomogeneous detector sensitivities for 60 – 190 mrad detector acceptance angles. Figure 3 shows that the total intensity for a simulation considering inhomogeneous detector sensitivity followed by electron flux weighting (analogous to experimental conditions) is in perfect agreement with simulations performed with homogeneous detector sensitivity (the ideal case).

[1] J. M. Lebeau and S. Stemmer, Ultramicroscopy 108 (2008), p.1653-8
[2] K. MacArthur, L. Jones, and P. Nellist, Journal of Physics: Conference Series (2013)
[3] A. Rosenauer and M. Schowalter, Springer Proceedings in Physics, vol. 120 (2007), p. 169–172.


Funding from the FWO Flanders, the EU FP7 (312483 - ESTEEM2), and the UK Engineering and Physical Sciences Research Council (EP/K032518/1) is acknowledged.

Fig. 1: Proposed flux-weighted normalization steps: a) experimental detector map, b) detector active area, c) determined flux pattern using camera length series, and d) flux-weighted sensitivity resulting from product of plots a) and c).

Fig. 2: Measured electron flux distribution from simulated camera length series. Using this plot, Figure 1c) is computed for the detector active area.

Fig. 3: Total scattered intensity for homogeneous (blue) and inhomogeneous (black) detector sensitivity. Red circles correspond to the total scattered intensity of inhomogeneous detector sensitivity after electron flux weighted normalization.

Type of presentation: Oral

IT-2-O-2221 Channeling Effects on the Accuracy of HAADF STEM Quantification of Bimetallic Catalyst Nanoparticles

MacArthur K. E.1, Lozano-Perez S.1, Ozkaya D.2, Nellist P. D.1
1Department of Materials Science, University of Oxford, Oxford, UK , 2Johnson-Matthey Technical Centre, Reading, UK
katherine.macarthur@materials.ox.ac.uk

Quantification of high angle annular dark-field scanning transmission electron microscope (HAADF STEM) images uses atomic resolution images as data sets to extract sample composition and thickness information from. A new quantification method based on calculating the scattering cross-section (CS) of each atomic column has been shown to be more robust to microscope image parameters.1 Using an automated code,2 this analysis involves converting images to an absolute scale through detector normalisation3, integrating over each atomic column within an image and multiplying by pixel area.
Whilst mathematically robust to microscope imaging parameters there are many other factors which affect the accuracy of quantification results. Channelling occurs when the columns of atoms in a crystal are aligned parallel to the incident electron beam and they act like miniature lenses providing an extra focusing effect on the probe. The subsequent atoms in the atomic column see a more focused probe than the first atom; resulting in them supplying increased scattering out to the detector. Along the length of the column, oscillations in intensity are seen, much as though the electrons are propagating in a waveguide. The whole column may therefore have a different scattering CS than the sum of the individual CSs of its constituent atoms. The ordering of atom types within an atomic column also affects the overall CS. Comparably another process known as de-channelling provides cross-talk between neighbouring columns of atoms. Cross-talk occurs when part of the probe is scattered and becomes channelled by a neighbouring column of atoms and then scattered out to the detector, thereby contributing information to the signal from neighbouring columns.
Atomic resolution requires viewing a crystal down a low order zone axis; any sample mis-tilt away results in a reduction in the channelling contribution and therefore a loss in CS, Figure 1. By 4̊ of tilt the effects of channeling are almost completely lost, whilst some atomic resolution remains. Top-bottom effects in the bimetallic columns are also diminished by 4̊ mis-tilt. At small tilts, however, there is a plateau region where the CS is independent of tilt, the size of which is dependent on probe convergence angle size, Figure 1. We believe the robustness to tilt when imaging on axis is more beneficial than the potential composition information gain from tilting far off a zone axis. This is particularly the case for nanoparticles which tilt under the beam. Combining with spectroscopy techniques will be necessary for gaining compositional information.

1 H E et al, Ultramicroscopy 133 (2013), p109-19
2 The Absolute Integrator is free for academic use from www.lewyjones.com/software/
3 JM LeBeau et al, Nano Letters 10 (2010), p4405-8


The research leading to these results has received funding from the European Union Seventh Framework Programme under Grant Agreement 312483 - ESTEEM2 (Integrated Infrastructure Initiative–I3), and from the EPSRC (grant number EP/K032518/1).

Fig. 1: The CS of a column of 7 Pt atoms in a crystal plotted against tilt away from a <110> zone axis. (a) Pt, blue, with the top or bottom atom replaced with Co, red and green, show a plateau region before dropping with tilt. With no channelling the CS would be 7x1 Pt atom, dotted red. (b) Tilt plot with different probe convergence angles.

Type of presentation: Oral

IT-2-O-2283 Measuring surface atom structures in Pt and Au nanocatalysts with high precision STEM imaging

Yankovich A. B.1, Berkels B.2,3, Dahmen W.2,4, Binev P.2, Sanchez S. I.5, Bradley S. A.5, Voyles P. M.1
1University of Wisconsin – Madison, USA, 2University of South Carolina, USA, 3Universität Bonn, Germany, 4RWTH Aachen, Germany, 5UOP LLC a Honeywell Company, USA
ayankovich@wisc.edu

TEM and STEM aberration correctors make sub-Ang resolution imaging routine. Once atoms are resolved, the question is how precisely can their positions be measured? TEM and STEM regularly achieve precision smaller than the resolution, but STEM encounters practical limits, such as image distortions from instabilities, before reaching the signal to noise ratio (SNR) fundamental precision limit. Combining multiple frames improves SNR and precision. Rigid registration is a common approach, but it does not correct for all types of instabilities. We have developed a non-rigid (NR) registration scheme for STEM images that accounts for all types of image distortions caused by instabilities during acquisition[1, 2].

Fig. 1 shows the results of the NR registration and averaging of a series of 512 HAADF STEM images of GaN. We show that sub-pm precision is achieved by fitting each Ga column to a 2D Gaussian, calculating the interatomic separations as shown in the histograms in Fig. 1(c) and (d), and using the standard deviation as the precision. The sub-pm precision in the x and y directions (0.74 and 0.85 pm) is reproducible and is 5-7 better than rigid registration.

A multislice simulated HAADF STEM image of a Si [110] dislocation core, shown in Fig. 2(a) was used to create an image series that includes distortions representative of real experiments, including thermal drift, floor vibrations, acoustic noise, electromagnetic fields, and electronic instabilities. Fig. 2(b) shows the NR registered and averaged image of the distorted series, demonstrating that inhomogeneous strain is preserved by NR registration.

NR registering and averaging STEM images allows for pm-scale measurements of surface atom bond length variation in Pt nanoparticles, which are prototypical noble metal catalysts. NP’s surface structure is crucial to their chemical activity but measuring it is extremely challenging. Fig. 3(a) shows that a Pt nanocatalyst exhibits pm-scale contraction of atoms at a (1-11)/(-1-11) corner and expansion of a (1-11) facet, with very little lateral displacement. Standardless atom counting on the same NR registered STEM image shows that the Pt NP is between 1 and 8 atoms thick with <1 atom uncertainty, as shown in Fig. 3(b). High precision in both positions and thickness are enabled by the extremely high SNR after NR registration. In general, STEM imaging with pm-precision will aid in understanding atomic displacement fields important in catalysis, defects, interfaces, and ferroic materials.

[1] Berkels et al, Ultramic. 138, 46 (2014).

[2] Yankovich et al, “Picometer-Precision Analysis of STEM Images of Pt Nanocatalysts” under review (2014).


We acknowledge funding from the Department of Energy, Basic Energy Sciences (DE-FG02-08ER46547), NSF (DMS 1222390), USC’s Special Priority Program SPP 1324, and the Excellence Initiative of the German federal and state governments, and the UW Materials Research Science and Engineering Center (DMR-1121288).

Fig. 1: (a) The first of 512 HAADF STEM images of GaN [11-20]. (b) Average of 512 frames after NR registration. The red dots are the positions of the columns identified by fitting. (c) and (d) Histograms of the X and Y separation measurements from (b).

Fig. 2: (a) Simulated HAADF STEM image of a Si [110] dislocation core model displaying an inhomogeneous strain field. (b) NR registered and averaged HAADF STEM image after the simulated image in (a) was made into an image series with typical distortions representative of experimental series.

Fig. 3: (a) Average of 56 HAADF STEM images of a Pt [011] nanoparticle after NR registration. The red arrows show magnified displacement vectors of the surface atoms and the displacement magnitudes are labeled in white. (b) The number of atoms in each column determined by comparing the experimental absolute intensities to simulations.

Type of presentation: Oral

IT-2-O-2611 The detection of single dopant atoms by high-resolution off-axis electron holography

Cooper D.1, Mayall B.1, McLeod R.2, Haigh S.3, Dunin-Borkowski R. E.4
1CEA-LETI, Minatec, 17 rue des Martyrs, 38054 Grenoble, Cedex 9, France, 2CEA-INAC, Minatec, 17 rue des Martyrs, 38054 Grenoble, Cedex 9, France, 3School of Materials The University of Manchester, Manchester M13 9PL, U.K., 4Ernst Ruska-Centre for Microscopy and Spectroscopy with Electrons, D-52425 Jülich, Germany
david.cooper@cea.fr

In 2003, a paper was published that discussed the need for a technique to detect single dopant atoms. It was stated that “single dopant atom detection is critical to device design, but it can also unravel complex and unexpected phenomena which may also open up new areas of materials exploration”. In 2014, it is still highly challenging to measure the locations, chemical identities and electrostatic potentials of single dopant atoms [1].

The technique of off-axis electron holography in the transmission electron microscope (TEM) involves the use of an electron biprism to interfere an electron wave that has passed through a thin specimen with a reference wave, in order to form an interference pattern that can be used to determine the phase shift of the electrons. Although electron holography has been used for many years to measure dopant potentials in semiconductors, soon devices will become so small that measurements of the electrostatic potentials of individual dopant atoms may be required.

Based on simulations, the expected step in phase shift across a single ionized P atom in Si is ~2π/1000 radians. This level of sensitivity can be reached easily in electron holographic measurements at low spatial resolution if long acquisition times are used. However, it is more of a challenge at atomic resolution. Here, we demonstrate progress towards the detection of single dopant atoms using electron holography. Figure 1(a) shows an electron hologram acquired at 80 kV using an aberration-corrected FEI Titan Ultimate TEM equipped with a high brightness gun, a monochromator and a single biprism. A careful choice of microscope lens settings allows holograms to be acquired with excellent interference fringe contrast and fine fringe spacing. Figure 1(b) shows an intensity profile extracted from the hologram, while Fig. 1(c) shows the experimentally measured phase resolution plotted as a function of interference fringe spacing, demonstrating that the conditions required to detect single dopant atoms are within reach if large numbers of phase images are added together.

Figure 2(a) shows part of an off-axis electron hologram of a thin MoS2 crystal recorded using an interference fringe spacing of 40 pm. The corresponding reconstructed phase image in Fig. 2(b) has a spatial resolution of 0.12 nm, while the line profile in Fig. 2(c) demonstrates that individual atomic columns with a spacing of 0.12 nm can be resolved. We are presently working towards the acquisition of signals from single dopant atoms in graphene and silicon and comparing our results with scanning TEM images. Great care is required to optimize specimen preparation and to minimize radiation damage, electron beam induced charging and contamination.

[1] Castell et al. Nature Materials 2, 129-131 (2003)


DC and RDB thank the ERC for the starting grant “Holoview” and the advanced grant “IMAGINE” respectively.

Fig. 1: (a) An off axis electron hologram acquired with a fringe spacing of 50 pm (b) profile of the fringe intensity showing a contrast of 33 %. (c) Experimentally measured phase resolution as a function of fringe spacing (spatial resolution is 2-3 times the fringe spacing).

Fig. 2: (a) Detail of an off-axis electron hologram of a MoS2 crystal with a fringe spacing of 40 pm (b) reconstructed map of the electrostatic potential profile and (c) profile showing that the 1.2 A spaced atoms have been resolved.

Type of presentation: Oral

IT-2-O-2414 A Method to Analyse the Chemical Composition in (InGa)(NAs) based on Evaluation of HAADF Intensity in STEM

Grieb T.1, Müller K.1, Mahr C.1, Cadel E.2, Beyer A.3, Talbot E.2, Schowalter M.1, Volz K.3, Rosenauer A.1
1Institute of Solid State Physics, University of Bremen, Otto-Hahn-Allee 1, 28359 Bremen, Germany, 2Groupe de Physique des Matériaux (GPM) UMR 6634, Normandie Université, Université et INSA de Rouen–CNRS, Av. de l’Université, BP 12, 76801 Saint Etienne du Rouvray, France, 3Materials Science Center and Faculty of Physics, Philipps University Marburg, Hans-Meerwein-Straße, 35032 Marburg, Germany
rosenauer@ifp.uni-bremen.de

InxGa1-xNyAs1-y is of technological interest for laser diodes in telecommunication and solar cells as both, In and N, lower the semiconductors band gap to emit or absorb in the infra-red spectral range. It was shown for ternary materials that an unknown chemical concentration (eg. of In in InGaN [1]) can be determined by high-angle annular dark field (HAADF) scanning transmission electron microscopy (STEM). For this purpose experimental HAADF intensities are compared with simulated ones. The experimental intensities are normalized to the total beam intensity which allows for determining the thickness in regions with known chemical composition. In this contribution this method is extended to evaluate the quaternary system InxGa1-xNyAs1-y. As a specific HAADF intensity cannot be allocated to a pair of concentrations (x,y) in a unique way, further information is needed. To this end, the local strain state is additionally determined from the high-resolution HAADF-STEM image.

The HAADF intensities were simulated with a frozen-lattice multislice approach implemented in the STEMsim software [2], considering thermal diffuse scattering (TDS). It was shown that for (In)GaNAs, besides TDS, Huang-scattering at static-atomic displacements (SADs) has to be taken into account [3]. SADs are distortions of the atomic lattice due to different covalent radii of In and Ga as well as As and N. The SADs were computed by relaxing the supercells using Keating's valence force field parametrization [4] in the LAMMPS code [5]. Fig. 1. shows the ratio of the simulated HAADF intensity of InGaNAs and GaAs versus specimen thickness for different In and N concentrations. For thicknesses above approx. 50 nm the intensity ratio increases not only with In but also with N concentration, although N has a smaller atomic number than As. This effect reveals the strong influence of additional scattering at SADs. An MOVPE grown InGaNAs/GaAs quantum-well sample is characterized by the outlined method. The mean concentrations of 32 % In and 2 % N (see concentration profiles in Fig. 2) are in good agreement with the results from XRD (marked by arrows). In addition, atom-probe tomography was applied to this sample, and the corresponding In profile is also shown in Fig. 2. Both, profile shape and mean concentration are in good agreement with the HAADF-STEM results.

[1] Rosenauer et al., Ultramicroscopy 111 (2011) 1316.

[2] A. Rosenauer and M. Schowalter, Springer Proc. Phys. 120 (2007), 169.

[3] Grillo et al., Phys. Rev. B 77 (2008), 054103.

[4] P. N. Keating, Phys. Rev. 145 (1966), 637.

[5] S. Plimpton, J. Comput. Phys. 117 (1995), 1.


We thank the DFG under contracts SCHO 1196/3-1, RO 2057/8-1 and GRK1782.

Fig. 1: Ratio of the simulated HAADF intensity for InGaNAs and GaAs (material contrast) as a function of specimen thickness for different indium concentrations (color) and nitrogen concentrations (line style). The HAADF intensity increases for specimen thicknesses above 50 nm with In and N concentration due to Z-contrast and scattering at SADs.

Fig. 2: Determination of the chemical composition of an InGaNAs layer embedded in GaAs. Concentration profiles from averaging concentration maps (HAADF analysis: indium and nitrogen) and from atom probe tomography (only indium). Concentrations derived from HRXRD are marked by arrows.

Type of presentation: Oral

IT-2-O-2538 Optimizing Phase Contrast Imaging in Aberration Corrected TEM

Kahl F.1, Hartel P.1, Linck M.1, Müller H.1, Haider M.1
1Corrected Electron Optical Systems GmbH, Heidelberg, Germany
frank.kahl@gmx.net

Since the realization of the first aberration corrected TEM [1], the number of corrected TEMs is still rapidly growing. Two key benefits are boosting the tremendous success of spherical aberration correction: vanishing delocalization and improved point resolution limit. The latter is achieved by using CS=C3 as additional optimization parameter to increase the aperture radius where the phase shift distribution (phase plate) of the elastically scattered electrons stays close to the optimum of +π/2 or −π/2 for dark or bright atom contrast, respectively. Various efforts have been made to optimize the phase plate, e.g. [2, 3], employing different measures for the distance of real and ideal phase plate as criterion for optimization. However, the different criteria lead to similar results.

Many users still use a π/4-limit for each aberration coefficient separately to assess the corrected state. However, advanced criteria such as minimizing the integrated mean quadratic deviation [3] or minimizing the largest deviation from the ideal phase over the aperture are much more sensible. While designing the SALVE II corrector [5] we used the latter criterion for optimizing the phase plate to assess imaging quality.

Fig. 1 shows the phase shifts generated by all axial aberration coefficients up to fifth order. Compensation schemes can be applied for aberrations of same multiplicity but different orders. In Fig. 1 potential partners are arranged within one column. The highest-order coefficient and the aperture size determine the optimum values for the lower-order coefficients (e.g. multiplicity 2: S5 given by corrector design; S3A1 optimized during alignment). The procedure for optimizing coefficients of non-zero multiplicity is similar to optimize C3 and defocus C1 for a given C5, except that the deviation from zero instead of +π/2 or −π/2 is minimized.

The performance of the SALVE II corrector for 40 kV is shown in Fig. 2. The phase plate (a) corresponds to the output of the CEOS software after aberration correction. The π/4-circle is misleadingly small as it is largely determined by C3. In image (b) only the sum of all non-round contributions is shown. Optimizing C1 for given C3 and C5 yields passband (c). Only with a full compensation scheme (d, e) for all fourth- and fifth-order aberrations using all adjustable lower order aberrations, a passband of up to 50 mrad can be achieved.

References:

[1] M. Haider et al, Nature 392 (1998), 768-769.

[2] O. Scherzer, Ber. Bunsen-Gesellschaft phys. Chemie 74 (1970), 1154-1167.

[3] M. Lentzen, Microsc. Microanal. 14 (2008), 16-26.

[4] M. Born, E. Wolf, Principles of Optics, 6th edition (Cambridge university press, Cambridge), p. 468.

[5] SALVE II project, <http://www.salve-project.de>.


none

Fig. 1: Table of axial aberrations visualized as phase plates. At the aperture edge ±6π is adopted; the phase is wrapped to [-π (black); π (white)[. Aberrations of same multiplicity but different order can partly compensate each other.

Fig. 2: Left: Phase plates for measured aberration coefficients at the SALVE II microscope operated at 40 kV. The passband in (c) demonstrates the reduced contrast transfer due to non-round residual aberrations. Right: With a suitable compensation scheme a passband up to 50 mrad can be achieved. The aperture radius of all phase plates is 75 mrad.

Type of presentation: Oral

IT-2-O-2644 Low-voltage and energy-filtered chromatic aberration-corrected high-resolution TEM on the PICO instrument

Houben L.1,3, Luysberg M.1,3, Barthel J.2,3, Mayer J.2,3, Dunin-Borkowski R. E.1,3
1Peter Grünberg Institut 5, Forschungszentrum Jülich, Germany, 2Gemeinschaftslabor für Elektronenmikroskopie, RWTH Aachen, Germany, 3Ernst Ruska-Centre for Microscopy and Spectroscopy with Electrons, Forschungszentrum Jülich, Germany
l.houben@fz-juelich.de

The advent of chromatic aberration correction in the transmission electron microscope (TEM) offers new prospects for high-resolution imaging at low accelerating voltages and for energy-filtered TEM (EFTEM). Examples of low voltage and energy-filtered images of complex oxides, thin layered materials, and nanoparticles will be presented, demonstrating the unique optical properties of an achroplanatic CEOS CCOR Cc/Cs corrector on Jülich’s chromatic aberration-corrected “PICO” microscope.
Atomic-resolution imaging at low voltages, currently down to 50 kV, allows high-resolution studies of radiation-damage-sensitive nanomaterials, such as CdSe/CdS nanostructures obtained from a cation exchange reaction, graphene and carbon nanotubes. It is also beneficial for the study of organic ligands, ligand-stabilised materials and materials that are functionalized with organo-metallic compounds.
The ability to acquire dose-efficient atomic-scale EFTEM elemental maps using the achroplanatic CCOR corrector on this microscope with a large field of view and large energy windows results from the fact that the chromatic focus spread is negligible after chromatic aberration correction. Figure 1 shows an example of an atomic-resolution elemental map of Ca obtained from a thin TEM foil of a calcium-titanate/strontium-titanate multilayer. Figure 2 shows a structural and compositional modulation in a (CeS)1.2CrS2 misfit-compound nanotube, which comprises alternating hexagonal CrS2 and rock-salt CeS sheets that have a repeat period of 11.2 Ångstrom.
The quantification of EFTEM elemental maps to provide atomic-resolution information about the local chemical composition of a specimen is complicated by the preservation of elastic contrast due to elastic scattering, which gives rise to thickness and defocus dependent contrast with fine details at all energy losses. Optical stability over minutes of collection time and careful image alignment and background subtraction are also required to obtain meaningful and reliable atomic-scale EFTEM elemental maps.


The authors thank J. Schubert (Forschungszentrum Jülich) and L. Penchakarla and R.Tenne (Weizmann Institute of Science) and M. Bar Sadan (Ben Gurion University) for kindly providing the materials used in this study

Fig. 1: High-resolution EFTEM images of a CaTiO3/SrTiO3 [001] multilayer sample taken at 300 kV. Ca L23 (a) pre-edge image, (b) post-edge image and (c) background-subtracted map. (d, e, f) Noise-reduced images obtained by averaging over 5x5 periods in the CaTiO3 layer, revealing Ca on the A sites of the pseudo-cubic perovskite lattice.

Fig. 2: Atomic-resolution micrographs and spectroscopic images of a CeCrS3 nanotube. (a) Schematic view of the CeCrS3 misfit lattice and a tubular structure. (b) HRTEM image and (c) magnified region of (b) with crystallographic projections superimposed. (d) EFTEM maps showing alternating Ce and Cr signals and (e) magnified region of (d).

Type of presentation: Oral

IT-2-O-2658 Atomic resolution secondary electron imaging and simulation of the SrTiO3 (001) c(6x2) surface reconstruction

Ciston J.1, Brown H. G.2, D'Alfonso A. J.2, Koirala P.3, Lin Y.3, Ophus C. L.1, Inada H.4, Zhu Y.5, Allen L. J.2, Marks L. D.3
1National Center for Electron Microscopy, Lawrence Berkeley National Lab., Berkeley, USA, 2School of Physics, University of Melbourne, Parkville, Victoria, Australia, 3Department of Materials Science and Engineering, Northwestern University, Evanston, USA, 4Hitachi High Technologies Corporation, Ibaraki, Japan, 5Condensed Matter Physics and Materials Science, Brookhaven National Laboratory, Upton, USA
jciston@lbl.gov

Aberration-corrected scanning transmission electron microscopes (STEM) enable simultaneous collection of atomically resolved signals relating to coherent scattering (bright field and annular bright field imaging), structural information based on thermal scattering to large angles known as high-angle annular dark-field (HAADF) imaging, bonding information using electron energy-loss spectroscopy (EELS), and element identification using both EELS and energy dispersive x-ray spectroscopy. Atomic resolution imaging based on secondary electron (SE) signals was demonstrated in 2009 [1], but the technique is only slowly growing in use. While these SE signals are highly surface sensitive due to the narrow escape depth of electrons with energies <50eV [2], atomic resolution SE imaging of surface structures that differ from a simple bulk crystal termination has not been previously demonstrated.


We have recently imaged the c(6×2) reconstruction on the (100) surface of single crystal SrTiO3 (Fig 1) through simultaneous atomic resolution SE and HAADF STEM with complementary HREM imaging. The ability to simultaneously record surface sensitive SE and bulk dominated HAADF signals at atomic resolution makes the problem of surface structure registration to the bulk lattice highly tractable, which is a distinct advantage over other scanning probe methods. By inspection it is clear that the registration of the previously reported structure, primarily refined from surface x-ray diffraction and scanning tunneling microscopy (STM) experiments [3], is incorrect. Interpretation of the experimental SE measurements from first principles is now possible using a recently developed quantum mechanical model to simulate SE images. This approach takes into account the probability and angular distribution of electrons that are ejected from atoms in the specimen when ionization of both core and semi-core electrons occurs [4]. Our preliminary simulations of a newly proposed structure of the SrTiO3-<100>-c(6×2) reconstruction are in good agreement with the bulk-subtracted experimental SE data (Fig 2), and consistent with previously reported data from STM, Auger spectroscopy, and x-ray diffraction measurements. The structure solved by SE imaging is also stable in density functional theory simulations, and is on the thermodynamic convex hull of known reconstructions on SrTiO3 <100>.


[1] Y Zhu et al., Nat. Mater. 8 (2009) p. 808
[2] A Howie, J. Microsc. 180 (1995) p.192
[3] CH Lanier, et al., Phys. Rev. B 76 (2007) 045421
[4] HG Brown et al., Phys. Rev. B 87 (2013) 054102


A portion of this work was performed at NCEM, supported by the Office of Science, Basic Energy Sciences of the U.S. Department of Energy under Contract No.: DE-AC02-05CH11231.

Fig. 1: (a) Weak beam dark field image of SrTiO3 001 c(6x2) single crystal with g=(200) showing atomically flat terraces (b) Transmission electron diffraction pattern of the c(6x2) reconstruction with c2mm symmetry acquired off-zone to minimize bulk dynamical diffraction

Fig. 2: (a) Bulk subtracted experimental secondary electron image of a 6×2 unit cell surface reconstruction on a <100> SrTiO3 substrate averaged over ~600 unit cells with c2mm symmetry enforced. The 200-keV probe had a convergence semi-angle of 25 mrad. (b) Prelimenary simulation of the result in (a) using the (projected) surface reconstruction indicated.

Type of presentation: Oral

IT-2-O-2674 Fast imaging with inelastically scattered electrons by off-axis chromatic confocal electron microscopy

Zheng C. L.1, Zhu Y.1, Lazar S.2, Etheridge J.1
1Monash University, Victoria, Australia, 2FEI Electron Optics, Eindhoven, The Netherlands
changlin.zheng@monash.edu

Imaging with inelastically scattered electrons is an important method for studying the composition and electronic properties of materials down to the atomic scale [1]. In this work, we describe an approach for fast mapping of inelastically scattered electrons using a scanning transmission electron microscope in a confocal mode, without using a spectrometer. We develop an off-axis scanning confocal electron microscope configuration using a double spherical-aberration corrected STEM/TEM. The electron probe is focused onto the sample at a significant angle to the optic axis of the imaging lens (Fig 1) and the probe-corrector retuned to form an atomic-scale electron probe in the specimen plane. Under the effect of the chromatic aberration of the imaging lens system, electrons with a chosen energy loss, the confocal energy, Ec, can be focused to a confocal point on the detector plane, while electrons of all other energies, including the zero loss electrons, will be chromatically defocused at that plane. In addition, the tilting of the incident beam laterally shifts the object exit wave in the back focal plane of the imaging lens, introducing an energy-related lateral displacement of the defocused probe. The inelastically scattered electrons are then chromatically dispersed both parallel and perpendicular to the optic axis, effectively separating electrons with different energies. In particular, electrons with the confocal energy can be detected selectively using an integrating detector. Using a synchronized set of scan-descan coils, these confocal electrons can remain focused on the detector as the electron probe is scanned across the specimen (Fig 1).

We illustrate the method with nanoscale core-loss chemical mapping of silver (M4,5) in an aluminium-silver alloy and atomic scale imaging of the low intensity core-loss La (M4,5@ 840eV) signal in LaB6 (Fig 2). The scan rates are up to 2 orders of magnitude faster than conventional STEM spectrum imaging methods recorded by CCD, enabling a corresponding reduction in radiation dose and improvement in the field of view [2]. Moreover, this off-axis chromatic confocal configuration offers the potential for fast nanoscale three-dimensional chemical mapping when coupled with the improved depth and lateral resolution of the incoherent confocal mode [3].

[1] R. F. Egerton, Electron energy-loss spectroscopy in the electron microscope (Plenum Press, New York, 1996), 2nd edn.

[2] C. Zheng, Y. Zhu, S. Lazar, J. Etheridge, Physical review letters, accepted (2014).

[3] T. Wilson and C. Sheppard, Theory and practice of scanning optical microscopy (Academic Press, London ; Orlando, 1984)


Funding is acknowledged from the Australian Research Council Grants DP110104734 and LE0454166.

Fig. 1: (a) Optical diagram of off-axis scanning confocal electron microscopy. (b) Chromatic defocused probe image of amorphous carbon. The refocused chromatic confocal energy is centered at energy loss of 300 eV. Chromatic confocal point is indicated by the arrow.

Fig. 2: Atomic resolution off-axis SCEM map of lanthanum M4,5 (~840 eV) core loss electrons in LaB6. The pixel dwell time is 1.5 ms with image size 256 x 256.

Type of presentation: Oral

IT-2-O-2881 Putting a New Spin on Scanning Transmission Electron Microscopy

LeBeau J. M.1, Sang X.1, Grimley E. D.1
1North Carolina State University, Raleigh, North Carolina, USA
jmlebeau@ncsu.edu

The full text of the abstract is not available. Please contact the presenting author.


Type of presentation: Oral

IT-2-O-2885 Three-dimensional location of a single dopant with atomic precision by aberration-corrected ADF STEM

Ishikawa R.1,2, Lupini A. R.2, Findlay S. D.3, Taniguchi T.4, Pennycook S. J.5
1University of Tokyo, Japan, 2Oak Ridge National Laboratory, USA, 3Monash University, Australia, 4National Institute for Materials Science, Japan, 5The University of Tennessee, USA
ishikawa@sigma.t.u-tokyo.ac.jp

Impurity doping is the key technology for enhancing physical and chemical properties in semiconductors. These functional dopants usually take the form of isolated single atoms, and the materials properties have strong sensitivities to the doping concentration, spatial distribution and three-dimensional location of the dopants. The recent development of aberration-corrected electron microscopy has allowed the determination of the two-dimensional spatial distribution of single dopants with atomic spatial resolution. However, this resolution has been achieved only in the lateral directions, and the last dimension, depth, has not yet achieved atomic resolution.

Here we use quantitative annular-dark field scanning transmission electron microscopy (ADF STEM)[1] to directly visualize isolated single Ce dopants accommodated in bulk w-AlN single crystals[2], exhibiting strong visible-light photo-luminescence. Through combining with frozen phonon image simulations, we determine the three-dimensional location of the Ce dopant with single atomic-layer precision in depth[3].

On the basis of the mean signal value comparison between the experiment and the simulations, we estimate the number of atoms per column, and the atomic-resolution thickness map is shown in Figure 1a. During sequential acquisition, we observed a single Ce dopant jump from X to Y through the interstitial site (Figure 1b-d). For the two columns of X and Y, we performed image simulations of all the possible dopant configurations in depth. In the thicker specimen, it may be difficult to uniquely determine the depth location of a single dopant owing to strong dynamical intensity oscillation. To overcome this issue, we implemented multi-component analysis such as mean signal value, maximum peak intensity and profile fitting. As shown in Fig. 1e, the experimental profile at atom X is well matched with that of the simulation of dopant location to a 9 unit-cell depth. And similarly, atom Y is located to a depth of 8 unit-cells. To develop more general method, we also analyze the same data set with Bayesian statistical model, which does not require a priori knowledge of the number of atoms. And we obtained the same depth locations of Ce dopant. By tracking a single dopant, we could begin to determine the three-dimensional atom diffusion path within bulk materials.

References

[1] R. Ishikawa, A.R. Lupini, S.D. Findlay and S.J. Pennycook, Microsc. Microanal. 20, 99 (2014).

[2] R. Ishikawa, et al., Sci. Rep. 4 3778 (2014).

[3] R. Ishikawa, A.R. Lupini, S.D. Findlay, T. Taniguchi and S.J. Pennycook, Nano Lett., (2014) in press.


R.I. acknowledges support from JSPS Postdoctoral Fellowship. A.R.L. acknowledges support by the U.S. DOE. S.D.F. acknowledges support under the Discovery Projects funding scheme of the Australian Research Council (Project No. DP110101570). T.T. acknowledges support by a Grant-in-Aid for Scientific Research on Innovative Areas "Nano Informatics" (Grant No. 25106006) from JSPS.

Fig. 1: Figure 1. Sequentially acquired Z-contrast images of w-AlN viewed along the [11-20] direction, (a) atomic-resolution thickness map, (b) averaged over frames of 1-19, (c) frame 20, (d) frames of 21-40. (e) Z-contrast profiles obtained from atom X (exp.) and the simulations of Ce locations to 8, 9, 10-unit cells.

Type of presentation: Oral

IT-2-O-3220 3D strain in HAADF – STEM images

Guerrero-Lebrero M. P.1, Bárcena-González G.1, Guerrero E.1, Liu Y.3, Kepaptsoglou D. M.4, Ramasse Q.4, Li L.3, Lazarov V. K.2, Galindo P. L.1
1Department of Computer Science and Engineering, Universidad de Cádiz, 11510 Puerto Real, Spain, 2Department of Physics, University of York, Heslington, York, United Kingdom, 3Department of Physics, University of Wisconsin, Milwaukee, WI 53211, USA, 4SuperSTEM Laboratory, STFC Daresbury Campus, Warrington, WA4 4AD, United Kingdom
maria.guerrero@uca.es

Strain mapping can be used to analyze materials at the atomic-column level, measuring local displacements and strain, and so revealing lattice translations, dislocations and/or rotations. Several methodologies have been developed to determine 2D strain field mapping from HRTEM images, either in real space (peak finding) [1, 2] or in Fourier space (geometrical phase analysis, GPA) [3]. Since 3-dimensional strain is independent of the image plane it might be ideal to gain insight into the behavior, shape and deformations of nanomaterials.

First, a HAADF focal series of 93 images (between 0nm and 14nm at steps of 0.15nm) of epitaxial Bi2Se3 (0001) thin films grown by atomic layer molecular beam epitaxy was taken with a Nion UltrastemTM 100 transmission electron microscope at 100 kV. Then, strain mapping of 23 images (from 19 to 41) was calculated using the Peak Pairs Analysis (PPA, [2]) plug-in for DigitalMicrograph available from HREM Research Inc.

Figure 1 shows one image of this focal series at a depth of 7nm, figure 2 shows the corresponding 2D strain map of this slice. It can observe that the screw dislocation position (red circle in figure 2) does not correspond to the apparent intensity change in HAADF STEM image (red circle in figure 1). The dark area in Figure 1 can be interpreted as a triangular spiral characterized by atomically smooth terraces, the way in which this material grows [4].

Figure 3 shows the 3D strain reconstruction where the screw dislocation movement, shape and tilt can be observed. Eshelby-Stroh twist [5] in the screw dislocation can be recognized, the upper part of the dislocation rotates in clockwise direction and the lower part turns in a counter – clockwise direction. The dislocation tilt has been estimated to be 6.3º with regard to the optical axis.

[1]Kret, S., Ruterana P., Rosenauer A., Gerthsen D. Extracting quantitative information from high resolution electron microscopy. Phys. Status Solidi (b) 227(1):247-295 (2001)

[2]Galindo, P. L, Sławomir, K., Sanchez, A.M., Laval, Y., Yañez A., Pizarro, J., Guerrero, E., Ben, T., Molina, S.I. The Peak Pairs algorithm for strain mapping from HRTEM images. Ultramicroscopy 107:1186-1193 (2007)

[3]Hÿtch, M. J., Snoeck, E., Kilaas, R. Quantitative measurement of displacement and strain fields from HREMicrographs. Ultramicroscopy 74:131–146 (1998)

[4]Liu.Y, Li, Y. Y., Rajput, S., Gilks, D., Lari, L., Galindo, P.L., Weinert, M., Lazarov V. K., Li, L.. Tuning Dirac states by strain in the topological insulator Bi2Se3. Nature Physics. (2014)

[5] Eshelby, J.D., Stroh, A.N. CXL. Dislocations in thin plates, Philosophical Magazines Series 7 42, 1401 (1951)


Fig. 1: HAADF Bi2Se3 slice at 7nm thickness. Apparent screw dislocation position is marked by the red circle.

Fig. 2: exx strain map that corresponds to the image in the Figure 1 calculated using PPA software [2].

Fig. 3: 3D strain reconstruction of a Bi2Se3 screw dislocation. Positive and negative strains are shown in red and blue respectively and the dashed red line represents the optical axis. The reconstruction makes it possible to collect a great deal of information about the dislocation motion.

Type of presentation: Oral

IT-2-O-3292 Position resolved single electron response of the HAADF-STEM detector and improved method for intensity normalisation

Schowalter M.1, Krause F. F.1, Grieb T.1, Mehrtens T.1, Müller K.1, Rosenauer A.1
1Institut für Festkörperphysik, Universität Bremen, Bremen, Germany
schowalter@ifp.uni-bremen.de

Quantification of HAADF-STEM images as demonstrated in [1,2] is based on normalising the image intensity with respect to the incident electron beam and comparison with image simulations. For that the electron beam is scanned over the detector and the intensities Iout outside and Idet on the detector yield the normalized intensity Inorm=(I-Iout)/(Idet-Iout).

Recently, it has been shown that accidental electrons can hit the HAADF detector, although the electron beam is scanned in a specimen free area [3]. From such a “vacuum image” the number of counts caused by a single electron can be inferred and intensity can be scaled in units of electrons per pixels which enables an alternative way for STEM image quantification [3] and error estimation based on electron statistics.

In this contribution we show that accidentally impinging electrons cause artifacts in the normalization of image intensity using the detector scan technique (DST) [1,2]. We introduce an improved DST which is able to avoid such errors. In addition, we demonstrate a method for measuring single electron signals as a function of detector position.

The red line in Fig. 1 depicts a linescan through an HAADF image of an a-C wedge evaluated using the conventional DST. The normalised intensity exhibits a significant shift towards negative intensities in vacuum. This can be attributed to accidentally impinging electrons [3], whose dose is different for a detector scan (image mode) and a vacuum scan (diffraction mode). To account for this difference we suggest to replace Iout in the numerator by the intensity Ivac obtained from a vacuum scan. The result of this is shown by the blue line in Fig. 1, where the intensity in the vacuum region vanishes. Fig. 2 shows histograms of vacuum images for different dwell times, revealing a large peak at 9900 due to the background level of the detector as well as further peaks corresponding to one or more electrons per scan position. Different dwell times yield fundamentally different curves so that Ivac depends on dwell time. Therefore, vacuum scan and image scan must be performed with the same dwell time.

We also measured the spatially resolved response of the detector to a single electron by drastically decreasing the beam current and taking a series of 256 detector scans with 2048 by 2048 pixels. The position of the single-electron peak was measured in bins of 16 by 16 pixels and the position of the zero-electron peak was subtracted. Fig. 3 nicely depicts the position sensitive single-electron response.

[1] J. M. LeBeau and S. Stemmer, Ultramicroscopy, 108, 1653 (2008).
[2] A. Rosenauer et al., Ultramicroscopy, 109, 1171 (2009).
[3] R. Ishikawa, et al., Microscopy and Microanalysis, 20, 99 (2014).


Fig. 1: Normalized intensity along a linescan. Normalization was done using Iout and Idet as derived from a detector scan (red) as well as using the background level from the vacuum image (blue).

Fig. 2: Log. of the freq. of intensities in vacuum images for different dwell times. Each curve shows peaks corresponding to one or more electrons impinging on a certain pixel. The intensity is normalised with respect to the dwell time, so that the distance between e.g. one-electron peak and zero-electron peak is inversely proportional to the dwell time.

Fig. 3: Two-dimensional map showing the one-electron response as a function of position on the detector.

Type of presentation: Oral

IT-2-O-3488 Linking Thickness, Channelling and Secondary X-ray signals in Atomic Resolution Scanning Transmission Electron Microscopy

Weyland M.1, Findlay S. D.2, D'Alfonso A. J.3, Allen L. J.3
1Monash Centre for Electron Microscopy, Monash University, Melbourne 3800, Australia, 2School of Physics, Monash University, Melbourne 3800, Australia, 3School of Physics, The University of Melbourne, Melbourne 3010, Australia
matthew.weyland@monash.edu

Quantification of EDX signals at atomic resolution can be treated by separation into two components; the scattering of electrons prior to ionisation and the subtleties of X-ray generation, emission and collection. Significant progress has recently been made concerning the first of these factors, with Forbes1 showing that consideration of both elastic and thermal scattering is required to explain anomalous contrast variations. However, true quantification requires a similarly detailed approach to the X-ray side of the system, with proportionality between signal and composition dependent on a multitude of factors including scattering cross-section, X-ray Fluorescence yield, Adsorption and detector geometry. Kotula has demonstrated a reference based approach2, scaling signals to averages from areas of known chemistry and Kothleitner recently showed the use of a ‘non-channelling’ (off-axis) approach to scale signals for quantification3. Both of these approaches offer a potential solution, but one of the main limitations is a lack of experimental data linking thickness, channelling and collected signal for a known specimen and well characterised instrument.

Results will be presented of a systematic study between thickness and EDX signal for known crystal structures and compositions. These will be matched with image simulation taking into account elastic and thermal scattering. The results presentenced will be carried out using a dual aberration corrected FEI Titan3, with well-defined probe illumination conditions, fitted with a standard 30mm2 ultra-thin window Si(Li) detector (0.13 sr) and a new 60mm2 windowless SSD detector (0.3 sr). Thickness will measured by position averaged convergent beam electron diffraction (PACBED), with EDX spectra acquired scanning over the same specimen area. Results will be presented from several specimens including Strontium titanate, GaAs/InGaAs radial nanowire heterostructures and Al-Cu alloys. By recording data from multiple areas with different thicknesses, trends between thickness, X-ray signal and channelling condition and its implications for quantitative high resolution EDX will be explored.

1. B. D. Forbes, A. J. D’Alfonso, R. E. A. Williams, R. Srinivasan, H. L. Fraser, D. W. McComb, B. Freitag, D. O. Klenov and L. J. Allen, PRB 86 024108, 2012
2. P. G. Kotula, D. O. Klenov and H. S. von Harrach, M&M 18(4), 2012
3. G. Kothleitner, M. J. Neish, N. R. Lugg, S. D. Findlay, W. Grogger, F. Hofer and L. J. Allen, PRL 112(8) 085501, 2014


The Australian research council is acknowledged for financial support through grants DP130102538 and LE0454166 (FEI Titan3).

Type of presentation: Poster

IT-2-P-1522 Modification of an existing laboratory room to house a Cs corrected microscope.

Papworth A. J.1, Nellist P. D.1
1The Department of Materials, The University of Oxford, Oxford OX1 3PH UK.
adam.papworth@materials.ox.ac.uk

The first Cs corrected microscopes became generally available at the beginning of the 21st century. Cs corrected microscopes require very tight environmental conditions, which often means that when purchasing a Cs corrected microscope you also have to build a new building as well. An example of purpose built high resolution microscope laboratory is SuperSTEM in the UK. However, it is possible to meet the environmental conditions by converting existing rooms, removing the need for a new building, and therefore making ownership of a Cs corrected microscope more affordable.

The required environmental conditions fall into four groups; Electromagnetic force (EMF), Temperature, Acoustic and Vibration, where the biggest cause of instabilities can come from outside interferences, such as trains, power cables and general road traffic. In most cases these outside interferences can be mitigated, for example moving power cables; however trains and traffic cannot be relocated. This paper outlines measures that can be made to minimise the environment factors by careful design and choice of equipment.

The environmental targets set by the design team were as follows; EMF AC and DC <0.5mG, mechanical displacement (vertical and horizontal) <0.3µm, acoustic noise for all frequencies with a flat field response microphone <60db, room temperature and air movements targets were set as; temperature 20oC ±0.2 hr-1 fluctuation, air flow within the room was to be vertical with a minimum air flow of 100mm sec-1. These targets were considered as reasonable to obtain while also meeting the requirements of the Cs corrected microscopes, which were under consideration at that time of planning.

The final design of the rooms, equipment, anti-vibrational block and services gave the following results. The EMF measurement gave an AC X and Y of 0.05mG and Z 0.2mG with no significant DC component. The acoustics were compromised by noise coming from the floor above with a maximum of 55db at 120Hz. Room temperature was measured at 20oC ±0.08 over a five hour period with a 2.5kW load. The room remain within specification, even when the door was left open for two hours. The isolation block showed no external vibration being measured from the roads or surrounding buildings above the normal background. Vibration measurements were also taken during the night as wel


Type of presentation: Poster

IT-2-P-1525 Cs CORRECTED ATOMIC RESOLUTION TEM IMAGES OF THE HUMAN TOOTH ENAMEL CRYSTALS

REYES-GASGA J.1,3, TIZNADO-OROZCO G. E.2, BRÈS E. F.1
1Unité des Matériaux et Transformation (UMET). Université de Lille 1, Sciences et Technologies. Bâtiment C6. 59650 Villeneuve d’Ascq. Lille, France., 2Unidad Académica de Odontología. Universidad Autónoma de Nayarit. Edificio E7, Ciudad de la Cultura “Amado Nervo”, C.P. 63190 Tepic, Nayarit, Mexico., 3Permanet Address: Instituto de Física, UNAM. Circuito de la Investigación s/n. Cd. Universitaria, 04510 Coyoacán, Mexico D. F., México
jreyes@fisica.unam.mx

Aberration-corrected HR-STEM and HAADF images of the human-tooth-enamel crystallites are presented. These spatial and energy resolutions images have allowed to get information on the physical meaning of the central dark line (CDL) defect which leads to the anisotropic dissolution of the crystals.
Human tooth enamel is composed in 95% of hydroxyapatite crystals (HAP, Ca10(PO4)6(OH)2) which are elongated-plate-like of 30 to 60 nm wide and 100 to 200 nm long, approximately [1]. They are organized in microns-sized structures named “rods” or “prisms” that go from the enamel–dentin junction to the enamel surface (figure 1). The chemical analysis in the micron range of enamel by different analytical techniques, mainly EDS spectroscopy, has indicated the existence of trace elements. Thus, carbonated hydroxyapatite (c-HAP) with Na, Mg, Cl, as trace elements, has been stabled for these crystals [2].
When observed with the Transmission Electron Microscope (TEM), the enamel crystallites show a structural defect of 1 to 1.5 nm width in their central region approximately, the Central Dark Line (CDL) (figure 2), whose structure and role in the enamel structure itself is unknown yet [3, 4].
Several studies have shown that the CDL favors their anisotropic dissolution [4, 5]. During the carious process, for example, the enamel crystals are destroyed in a systematic fashion: first a series of hexagonal holes aligned along the [11-20] are observed, then the holes develop anistropically along the [0001] direction and cross the whole crystals [4, 5].
Enamel crystals are electron beam sensitive, other important parameter against the HRTEM observation (figure 3). Therefore, the use of low electron doses is critical during the study of the CDL. Therefore aberration corrected HR-STEM is the appropriated equipment for carrying out the chemical and structural analyses of the enamel crystallites.
Human tooth enamel samples were obtained from permanent non-carious human molar teeth, extracted for orthodontic or periodontal reasons. Samples were prepared in the FIB-FEI QUANTA 200 3D equipment using the two beams system.
JRG thanks to DGAPA-UNAM (contract IN106713), CONACYT and PASPA-DGAPA-UNAM for sabbatical support.

References
1. R.Z Le Geros, Calcium Phosphates in Oral Biology and Medicine ed H M Myers (San Francisco, CA: Karger). 1991.
2. G.E. Tiznado-Orozco et al., J. Phys. D: Appl. Phys. 42 (2009).
3. J. Reyes-Gasga et al., J. Mater. Sci. Mater. Med. 19, 877-882 (2008).
4. E. Brès et al., Journal de Physique, 51, C1-97-102, (1990).
5. E. Brès et al., Ultramicroscopy, 12, 367-372 (1984).


This research has received funding from the European Union Seventh Framework Programme under Grant Agreement 312483-ESTEEM2 (Integrated Infrastructure Initiative–I3).

Fig. 1: TEM image of human tooth enamel crystals inside the micron-sized structure named “prism”.

Fig. 2: Magnification of one of the human tooth enamel crystals shown in figure 1. The row indicates the presence of the “central dark line”.

Fig. 3: HRTEM image of a human tooth enamel crystals aligned along the [0001] direction. The arrow indicates the electron beam damage.

Type of presentation: Poster

IT-2-P-1600 Phase Contrast Transfer Function for Differential Phase Contrast in High Resolution Local Electric Field Measurements

Majert S.1, Kohl H.1
1Physikalisches Institut und Interdisziplinäres Centrum für Elektronenmikroskopie und Mikroanalyse (ICEM), Westfälische Wilhelms-Universität Münster, Wilhelm-Klemm-Straße 10, 48149 Münster, Germany
s_maje03@uni-muenster.de

Differential Phase Contrast (DPC) is a contrast mechanism that can be utilized in the Scanning Transmission Electron Microscope (STEM). Since the advent of DPC, the technique has been used to image magnetic fields within a specimen [1]. To this end, a ring detector is divided into four quadrants and the direct electron beam is placed within the ring, only overlapping a small part of the detector. In a classical interpretation, the direct beam is slightly tilted by the magnetic fields in the specimen, so that subtraction of different detector segement signals yields DPC. Recently, this DPC geometry was also employed to investigate local electric fields with high resolution [2,3].

To determine whether this interpretation of DPC is still valid in high resolution, the wave nature of the electrons has to be taken into account. This can be done by calculating the Phase Contrast Transfer Function (PCTF) [4] to examine the contrast mechanism. For DPC, the PCTF should be proportional to the spatial frequency k=2π/λ whereas a PCTF constant as a function of the spatial frequency k would indicate conventional phase contrast.

Assuming an ideal lens, which is a good approximation for an aberration corrected STEM, the PCTF for a weak phase object can be calculated using elementary geometry. A cut through the two dimensional PCTF, evaluated for the parameters of a local electric field measurement, is shown in fig.1. It is striking that the area in which the PCTF is proportional to k is rather small (up to ca. 0.2 1/Å as seen in fig. 2), indicating that, for high spatial frequencies, DPC would not occur. While this is unproblematic at low resolutions (where the configuration described above leads to an improved signal to noise ratio [5]), it suggests that under these conditions the classical model is not valid for high spatial frequencies and the detector setup is therefore not suited for high resolution DPC applications.

The calculated PCTF shows that, for the given parameters, DPC is limited to spatial frequencies of about 0.2 1/Å. We are currently looking for possibilities to increase the resolution by optimizing the detector geometry.

[1] J. N. Chapman et al., Ultramicroscopy 3 (1978) 203
[2] M. Lohr et al., Ultramicroscopy 117 (2012) 7
[3] N. Shibata et al., Nature Physics 8 (2012) 611
[4] H. Rose, Ultramicroscopy 2 (1977) 251
[5] J. N. Chapman et al., IEEE trans. on magn. 26 (1990) 1506


Fig. 1: PCTF L(k) for an ideal microscope with an acceleration voltage of 300 kV, an aperture angle of 21.6 mrad, an inner detector angle of 21.0 mrad and an outer detector angle of 40.7 mrad, corresponding to the configuration in high resolution local electric fields measurements.

Fig. 2: Low spatial frequency region of the PCTF in fig.1, showing that DPC only occurs for spatial frequencies k smaller than ca. 0.2 1/Å.

Type of presentation: Poster

IT-2-P-1633 Lifetime of the aberration-corrected optical state in HRTEM

Barthel J.1, Thust A.2
1Central Facility for Electron Microscopy, RWTH Aachen, Germany, 2Peter Grünberg Institute, Forschungszentrum Jülich GmbH, Germany
ju.barthel@fz-juelich.de

The technique of high-resolution transmission electron microscopy (HRTEM) experienced an unprecedented progress through the introduction of hardware aberration correctors, and by the improvement of the achievable resolution to the sub-Ångström level. As a consequence, the required precision level to measure and to adjust the optical properties of transmission electron microscopes has become increasingly demanding. A second consequence of this development, which has received little attention so far, is that aberration correction at a given resolution requires additionally a well-defined amount of optical stability. We investigate the qualification of a variety of high-resolution electron microscopes to maintain an aberration-corrected optical state in terms of a lifetime.

A comprehensive statistical framework is introduced for the estimation of the optical lifetime [1]. The temporal evolution of the twofold astigmatism is extracted from a series of images recorded over several minutes (Fig. 1). The twofold astigmatism serves as representative indicator for the optical stability since it is one of the most volatile image aberrations, has a strong influence on the image contrast, and can be measured rapidly in a simple experiment [2]. A model-based evaluation method was developed, which allows us to distinguish between two major components of astigmatism fluctuations, a random walk and a constant drift. A very useful output of the model-based evaluation is a probability curve (Fig. 2), which informs the operator about the chance to still work in an aberration-corrected state after a given timespan.

Optical stability evaluations for different high-resolution microscopes reveal surprisingly short lifetimes on the order of a few seconds up to a few minutes. The observed short lifetimes denote a critical limitation of the timespans between aberration measurement, aberration correction and the actual imaging. Therefore further investigations and technical developments are necessary in order to stabilize electron microscopes with respect to their sub-Ångström qualification. Since the topic of optical stability turns out to be of similar importance as the topic of resolution itself, we recommend to include a routine assessment of the optical stability in acceptance tests for high-resolution microscopes operating in the discussed resolution regime. For this purpose, the lifetime evaluation procedures developed in this work have been implemented in a user friendly and freely downloadable software [3].

References:

[1] J. Barthel, A. Thust, Ultramicroscopy 134 (2013), p. 6.
[2] J. Barthel, A. Thust, Ultramicroscopy 111 (2010), p. 27.
[3] J. Barthel, http://www.er-c.org/barthel/pantarhei/, (Feb 2014).


J.B. gratefully acknowledges funding within the core facilities initiative of the German Science Foundation (DFG) under the grant number MA 1280/40-1.

Fig. 1: Evolution of the twofold astigmatism extracted from images of amorphous carbon. Already after one minute the astigmatism fluctuations violate the π/4 limit for 300 keV electrons and for a microscope resolution of 0.8 Å.

Fig. 2: Decay of the probability to still work in an aberration-corrected state, evaluated from the astigmatism fluctuations shown in Fig. 1.

Type of presentation: Poster

IT-2-P-1677 Preparation of thin film specimen by Cryo Ion Slicer for TEM cross-section (XTEM) observation 

Siddheswaran R.1, Medlín R.1
1New Technologies Research Centre, University of West Bohemia, Plzeň-30614, Czech Republic
rajendra@ntc.zcu.cz

An essential part of research in thin film fabrication is the microstructural analyses like morphology, grain distribution, texture, thickness of the layers and orientation of film structure. For such characterization, cross-sectional transmission electron microscopy (XTEM) is a very essential tool for the study of structure, phase, defects and interfaces. For such analyses, it is necessary to make the film electron transparent in a direction perpendicular to the interfaces. One of the methods is cryo ion slicing (from JEOL) with specific sample preparation procedure different from the PIPS from Gatan. The preparation of cross-sectional specimens with ion slicer are usually done by fabricating a sandwich structure (Thin film/Glue/Cover glass) and subsequently thinning it to transparent for electrons (thickness of the order of <50 nm for TEM and <10nm for HR-TEM). The cross-section specimen preparation is generally time consuming, specimen dependent and consequently a trial and error method. But the features of XTEM observations are in results more informative and necessary in addition with the other methods of the observations, i.e. XRD, optical studies or micrographs of scratched samples from thin films.
The present work describes the preparation of thin film specimen, includes mechanical (pre-preparation) and Ar+ ion slicing (milling). It was successfully used for the preparation of a-Si:H/a-SiO2, nc-Si/a-SiO2 and ZnO thin film specimens for transmission microscopic analyses. The pre-preparation of sample for ion milling consists of cutting samples by diamond disc using low speed saw cutter (Buehler IsoMet) and mechanically thinning using JEOL Handy Lap to get the specimen dimension 2.5mm×500µm×100µm with plan-parallel to the surfaces. The ion slicing was carried out using JEOL IB-09060CIS Cryo Ion Slicer using Ar gas of purity 99.9999%. Finally thin regions of range from 100nm to 10nm were achieved over the thin film layers. A very thin cross-section of ~10nm could be used to obtain high resolution TEM images. 


The result was developed within the CENTEM project, reg. no. CZ.1.05/2.1.00/03.0088. 

Type of presentation: Poster

IT-2-P-1739 Development of the on-line DigitalMicrograph scripts for TEM imaging using the “Virtual TEM”.

Potapov P.1
1temDM, Dresden, Germany
info@temdm.com

Practical high resolution imaging still depends strongly on skills and smartness of operators. DigitalMicrograph scripting [1] can facilitate the practice of high resolution imaging in numerous ways: the fine tuning of the aberrations; the automatic compensation for the specimen and lenses drift; easy navigation over the area of interest; minimizing the applied electron dose and therefore reducing the radiation damage; the rapid switching between the imaging and spectroscopy modes.
DigitalMicrograph scripting provides users with a set of commands controlling the TEM hardware. However these commands are available in the on-line version of DigitalMicrograph only, i.e. they require the physical connection with TEM. This hinders the progress of the on-line scripting - for TEM time is expensive and should be used for imaging, not programming.
The present work introduces a plugin “Virtual TEM” that simulates all the scripting commands for communication between TEM and DigitalMicrograph. With this plugin, a user is able to edit, debug and roughly test the on-line scripts with no actual connection to the TEM; when being in office, at home or during air travel. The plugin imitates a simple TEM interface with the basic control of magnification, focus, stage, beam and stigmators (Fig.1). Depending on the instrumental settings, the “Virtual TEM” generates the image of the model object that can be captured by DigitalMicrograph and used as a feedback for the communication commands. The simplest example scripts - “Auto Acquisition”, “Focal Series”, “Correct Stage Drift” et cet - are provided as a part of the “Virtual TEM” package. The example scripts are aimed to be a seed for the development of more sophisticated customized tools.
Beyond the simplest examples, the advanced on-line DigitalMicrograph scripts are presented. The “Batch Recording” allows a user to shoot the images by a single button touch and automatically put them into the image container optimized for easy resizing and sorting (Fig.2). The “Click Mover” provides the convenient navigation over the place of interest by simple mouse double-click on the live high resolution image (Fig.3).
The plugins can be free downloaded from [2].

References:
[1] http://www.gatan.com/resources/scripting/
[2] http://www.temDM.com/


Fig. 1: “Virtual TEM” interface including the virtual camera menu, virtual TEM and virtual Filter control panels. The image and spectrum of the generated model object are displayed.

Fig. 2: “Batch Recording” tool provides convenient live imaging and storing the recorded images in the image container.

Fig. 3: “Click Mover” tool moves the feature of interest (mouse double-clicked) to the center of the live image.

Type of presentation: Poster

IT-2-P-1751 The atomic structure of epitaxially strained LaNiO3-LaGaO3 superlattices

Qi H. Y.1, Kinyanjui M. K.1, Biskupek J.1, Benckiser E.2, Habermeier H. U.2, Keimer B.2, Kaiser U.1
1University of Ulm, Central Facility of Electron Microscopy, Electron Microscopy Group of Materials Science, Albert Einstein Allee 11, D-89069 Ulm, Germany, 2Max Planck Institute for Solid State Research, Heisenbergstrasse 1, D-70579 Stuttgart, Germany
haoyuan.qi@uni-ulm.de

Many functional properties of ABO3 perovskite oxides are closely coupled to slight structural distortions in the perovskite lattice, thus defined symmetry constraints in oxide heterostructures can be used to access novel properties that are not found in bulk constituents [1].
Here, we present our study of an epitaxially strained [4 unit cell (u.c.)//4 u.c] х8 LaNiO3-LaGaO3 (LNO-LGO) superlattice grown on (001) SrTiO3 (STO) substrate (see the model shown in Figure 1). Due to the lattice mismatch, the superlattice is subject to tensile strain. We focus on the determination of the strain-induced distortions (changes in Ni-O bond length) and tilts (changes in Ni-O-Ni bond angle) of the corner-sharing octahedral network, as they may drastically influence the functionalities of the heterostructure. In order to discover the correlation between the NiO6 octahedra rearrangement and the functional properties of the material system, it is essential to study the interfacial structure with atomic-level accuracy. We studied the atomic structure of the octahedral network by means of aberration-corrected high-resolution transmission electron microscopy (AC-HRTEM). In order to enhance the image contrast, negative Cs imaging (NCSI) was applied [2].
Figure 2 is an experimental image of the LNO-LGO superlattice acquired in the vicinity of the top surface in [110] projection. It is clearly seen that the LNO and LGO layers manifest difference in both the zigzagness (out-of-plane corrugation) of BO2 layers and the image contrast. The smooth oscillation of the tilt angles indicates: 1) dissimilarity in tilt systems of each material, 2) proximity effect between adjacent layers. As a result of coherent epitaxial growth, the in-plane lattice parameter d220 remains constant while only the out-of-plane lattice parameter d001 varies. We will discuss further investigation at the substrate-layer interface and answer the question on the assignment of the layers. However, from merely HRTEM imaging, it is difficult to find out which of the layers, LNO or LGO shows the higher out-of-plane corrugation.

[1] H.Y. Hwang, Y. Iwasa, M.Kawasaki, B. Keimer, N.Nagaosa and Y. Tokura, Nat. Mater. 11, 103 (2012).

[2] C. Jia, M. Lentzen and K. Urban, Microsc. Microanal. 10, 174 (2004).


We are grateful to S. Grözinger for assistance with TEM specimen preparation and the German Research Foundation (DFG) for financial support (project DFG: KA 1295/17-1)

Fig. 1: Atomic structure model of a [4 unit cell (u.c.)//4 u.c] х8 LNO-LGO superlattice grown on STO substrate (viewed in [110] orientation). Unit cells are defined by pseudo-cubic symmetry axes.

Fig. 2: Experimental aberration corrected image of the LNO-LGO superlattice acquired in [110] projection in the vicinity of the top surface at 300kV. The oscillation of the tilt angles shown in the right diagram indicates non-identical tilt systems in LNO and LGO layers, which is clearly visible from the magnified areas (red and yellow rectangles).

Type of presentation: Poster

IT-2-P-1996 Structural and spectroscopic analyses of exfoliated 2-D transition metal dichalcogenides nanosheets with special emphasis on TEM

Pokle A. S.1, Coelho J.2, Mendoza B.2, Nicolosi V.1, 2
1School of Physics, Advance Microscopy Lab, Trinity College Dublin, Ireland, 2School of Chemistry, CRANN, Trinity College Dublin, Ireland
poklea@tcd.ie

Having unique physical, chemical and structural properties, 2-D nanomaterials such as the Transition Metal Dichalcogenides (TMD’s) have attracted considerable attention. Similar to graphene, TMD’s are atomically thin two dimensional materials with electronic properties different from their bulk counterparts. Graphene’s vanishing band-gap for semiconductor application poses a major setback. As a result, it is not suitable for logic applications, because devices cannot be switched off. On the other hand, 2D TMDs (i.e. MoS2, NbSe2, MoSe2, etc.) are semiconductors with a variety of tunable bandgaps. This property makes them perfect contenders for replacing Silicon in the semiconducting industry. In addition, different classes of 2-D materials such as like Transition Metal Oxides (TMOs) have shown to exhibit excellent electrical, optical and electrochemical properties. In virtue of this properties they have become excellent candidates for applications in energy storage devices such as lithium-ion batteries and supercapacitors.

Few layered or single-layered TMDs and TMOs can be obtained either through exfoliation of bulk material or by a bottom-up synthetic approach. The approached used in our group is the synthesis of 2D materials by liquid-phase exfoliation. This method produces atomically-thin and few-layers sheets dispersed in a solvent media. In order to apply these materials to feasible applications it becomes crucial to analyse their structure and correlate that to the ultimate properties when these materials are used in devices.

In this work we present a structural and spectroscopic characterization of a range of liquid-phase exfoliated 2D materials. Major focus is given to the study of their crystallographic structure, presence of defects, possible oxidative processes, and edge-effects. For that we use a combined approach, where by X-ray diffraction (XRD), scanning electron microscopy (SEM), high-resolution (scanning) transmission electron microscopy - HR(S)TEM, energy dispersive X-ray spectroscopy (EDX), electron energy loss spectroscopy (EELS) and X-ray photoelectron spectroscopy (XPS) are all used to obtain a throughout characterization of the materials.


The authors gratefully acknowledge funding from the FP7 People Network – ITN: Initial Training Network and Science Foundation Ireland – European

Research Council: SFI - ERC Support Program

Fig. 1: TEM image of MoSe2 (LHS) and WSe2 with their respective SADP.

Fig. 2: HR-TEM image of MoSe2 (LHS) and WSe2 with FFT 

Type of presentation: Poster

IT-2-P-2016 Surface Science of Metal Oxides by High-resolution TEM

Yu R.1, Zhan W.1, Lu S. R.1, Zhu J.1
1Tsinghua University, Beijing, China
ryu@tsinghua.edu.cn

Surfaces of metal oxides are of crucial importance for a variety of technological applications such as heterogeneous catalysis, thin film growth, gas sensing, and corrosion prevention [1]. Due to the complexities of oxides in crystal structure and electronic structure, however, the surface science of oxides lags far behind that of metals or semiconductors. Conventional surface-science techniques, typically scanning tunneling microscopy (STM) and low energy electron diffraction (LEED), are usually limited to surfaces of single crystals with relatively simple structures. Metal oxides are usually good insulators, either band insulators or Mott insulators, making them not suitable for STM, LEED, and most of spectroscopic methods using low energy electrons as probes. On the other hand, the complex atomic structures of oxides results in too many structural parameters to be determined by spectroscopy or diffraction methods. Recent developments in high-resolution transmission electron microscopy (TEM) provide us opportunities to overcome the above difficulties. With the realization of aberration-correction, the point resolution of TEM has been improved into the milestone 1 Angstrom scale. In addition, the correction of the spherical aberration has almost eliminated the contrast delocalization in high-resolution images. Therefore, high resolution TEM becomes an even more powerful tool than before for materials research at a truly atomic-scale. Here, we will present our recent works on atomic and electronic structure of oxide surfaces [2-4]. We will show that the structure and dynamics of oxide surfaces can be directly imaged and measured at the sub-angstrom scale with an accuracy of picometers, comparable to that obtained by conventional surface science techniques on single crystals. Special attention will be on line defects at the surfaces of MgO and Fe2O3.

References:

1. V. E. Henrich, and P. A. Cox, The Surface Science of Metal Oxides (Cambridge University Press, Cambridge, 1994).

2. R. Yu, L.H. Hu, Z.Y. Cheng, Y.D. Li, H.Q. Ye, J. Zhu, Phys. Rev. Lett., 105, 226101 (2010).

3. M.R. He, R. Yu, J. Zhu, Angew. Chem. Int. Ed., 124, 7864 (2012).

4. S.R. Lu, R. Yu, J. Zhu, Phys. Rev. B, 87, 165436 (2013).


This work was supported by National Basic Research Program of China (2011CB606406), NSFC (51071092, 51371102, 11374174, 51390471, 51390475), and Program for New Century Excellent Talents in University. This work used the resources of the Beijing National Center for Electron Microscopy and Shanghai Supercomputer Center.

Fig. 1: (a) HRTEM image of the α-Fe2O3 (-1102) surface defect viewed in the [1-101] direction. The inset shows the simulated image of the relaxed structure  by DFT calculations (b).

Type of presentation: Poster

IT-2-P-2042 Contrast Investigation of Annular Bright-Field Imaging in Scanning Transmission Electron Microscopy of LiFePO4

Zhou D.1, Sigle W.1, Müller K.2, Rosenauer A.2, Zhu C.3, Kelsch M.1, Maier J.3, van Aken P. A.1
1Max Planck Institute for Intelligent Systems, Stuttgart Center for Electron Microscopy, Heisenbergstraße 3, 70569 Stuttgart, Germany, 2Institute of Solid State Physics, University of Bremen, Otto-Hahn-Allee 1, D-28359 Bremen, Germany, 3Max Planck Institute for Solid State Research, Heisenbergstraße 1, 70569 Stuttgart, Germany
danzhou@is.mpg.de

Light elements, such as lithium, are difficult to detect using high-angle annular dark-field imaging (HAADF) in STEM because of their weak atomic scattering. Recently, a novel imaging mode in aberration-corrected STEM was presented which uses an annular detector spanning an angular range mainly within the illumination cone of the focused electron beam1. It was shown that due to the smaller dependence on atomic number Z, approximately Z1/3 compared to Z2 in HAADF, the resultant images enable one to visualize the light element columns2. This imaging mode has been called annular bright-field (ABF) imaging. The contrast differences are clearly visible in Figure 1 which compares HAADF (b) and ABF (c) images of LiFePO4 structure in [010] orientation (a).

In this work, we studied the contrast of ABF imaging in STEM on LiFePO4, correlating results of experiments from the newly installed probe-aberration-corrected JEM-ARM 200CF and simulations using the STEMsim program3. Previous work briefly demonstrated the possibility to acquire direct images of LiFePO4 and partially de-lithiated LiFePO4 at atomic resolution4. Figure 2 presents an experimental comparison of the visualization of lithium in LiFePO4 with HAADF and ABF. The present work aims at presenting a more detailed description of the contrast dynamics of ABF imaging of LiFePO4 with a view to its interpretation, and optimization. Thickness, defocus, angular range, and possible contamination introduced by sample preparation are taken into account to understand the image contrast. In particular, the probe and detector configurations in the microscope are taken into consideration to step from qualitative to quantitative contrast evaluation.

1. Okunishi, E.; Ishikawa, I.; Sawada, H.; Hosokawa, F.; Hori, M.; Kondo, Y., Visualization of Light Elements at Ultrahigh Resolution by STEM Annular Bright Field Microscopy. Microsc Microanal 2009, 15, 164-165.

2. Findlay, S. D.; Shibata, N.; Sawada, H.; Okunishi, E.; Kondo, Y.; Ikuhara, Y., Dynamics of annular bright field imaging in scanning transmission electron microscopy. Ultramicroscopy 2010, 110 (7), 903-923.

3. Rosenauer, A.; Schowalter, M., STEMSIM-a New Software Tool for Simulation of STEM HAADF Z-Contrast Imaging. Springer Proc Phys 2008, 120, 169-172.

4. Gu, L.; Zhu, C. B.; Li, H.; Yu, Y.; Li, C. L.; Tsukimoto, S.; Maier, J.; Ikuhara, Y., Direct Observation of Lithium Staging in Partially Delithiated LiFePO(4) at Atomic Resolution. J Am Chem Soc 2011, 133 (13), 4661-4663.

 


The research leading to these results has received funding from the European Union Seventh Framework Programme [FP7/2007-2013] under grant agreement n°312483 (ESTEEM2).

Fig. 1: (a) Projection of the LiFePO4 crystal structure in [010] orientation. Simulated HAADF (b, angular range 40-100 mrad) and ABF (c, angular range 11-22 mrad) imaging using STEMsim. The parameters used for simulations are: high voltage 200 kV, convergence angle 22 mrad, spherical aberration 0 mm, defocus 0 nm, specimen thickness 30 nm.

Fig. 2: As acquired experimental results on the visualization of Li in LiFePO4 with HAADF (a, 90-370 mrad) and ABF (b, 11-22 mrad) imaging on the probe-aberration-corrected JEOL JEM-ARM 200CF microscope using a convergence angle of 22 mrad and a probe size of about 0.08 nm. (c) Assignment of atomic column in [010] orientation.

Type of presentation: Poster

IT-2-P-2046 HAADF STEM characterization of BST-MgO interface structure

Kuskova A. N.1, Zhigalina O. M.1, Khmelenin D. N.1
1Shubnikov Institute of Crystallography, Russian Academy of Sciences
xorrunn@gmail.com

The properties of thin perovskite ferroelectric films can be different from those of bulk materials, that is caused by the mechanical stress at the film–substrate interface [1]. Such stress is usually relaxed by the formation of misfit dislocations at the heterostructure interface.
It has been shown [2] that the degree of stress in epitaxial BST thin films is a function of thir thickness. In this study we have introduced a high angle annular dark field (HAADF) scanning transmission electron microscopy (STEM) investigation of this heterostructure interface combined with the modelling of the HAADF STEM images, a geometry phase analysis [3] and a statistical quantitative analysis [4].
The perovskite structure of BST (Ba0.8Sr0.2TiO3) allows two types of starting planes for growth on (100) cubic MgO substrate: the Ba(Sr)O or the TiO2 planes, which could enable different chemical bonding at the interface. Since the HAADF STEM image intensity is proportional to Z2 of the scanned crystal (Z is average atomic number of atomic columns), the Ba(Sr) atomic columns where are Z=52 observed as the brightest dots in the image (Fig.1). TiO columns with Z=15 have a lower brightness and MgO columns where are Z=10 demonstrated the least brightness in the images. The pure O columns (Z=8), located between the brightest Ba(Sr) columns, are not visible in the image. The misfit dislocations marked by arrows on Figure 1(a) were visualized by the geometric phase analysis [3]. Figure 1(b) illustrates the enlarged dislocation core and its Burgers vector identified as ½аBST[010].
It is obvious that to obtain the information about the chemical interface structure based only on the direct observation of changes in the image contrast is not correct. Model-based statistical quantitative analysis has quantified the chemical composition of ‘unknown’ atomic columns at the interface based on a comparison of their scattered intensities with ones of ‘known’ columns located far from the interface [4]. Figure 1d illustrates the estimated peak volumes for Figure 1c. This analysis of intensities of different types of planes has indicated that the first atomic layer of the film does not lie on top of the substrate, but is embedded into the upper layer of MgO.

1. Y.S. Kim, D.H. Kim, J.D. Kim, et al., Appl. Phys. Lett., 86 (2005), p.102907
2. O.M. Zhigalina, A.N. Kuskova, R.V. Gaynutdinov, et.al, Journal of Surface Investigation: X-Ray, Synchrotron and Neutron Techniques. 4 (2009). p. 542-547.
3. A.K. Gutakovskii, A.L. Chuvilin, Se Ahn Song, Izvestiya RAS, ser. phys., 71 (2007). p.1464-1470-
4. S. Van Aert, J.Verbeeck, R.Erni et al., Ultramicroscopy, 109 (2009) p. 1236-1244
5. J.W. Reiner, F.J. Walker, & C.H. Ahn. Science 323 (2009), p. 1018–1019.


This work was done using IC RAS Research Center equipment and supported by the Ministry of Education and Science of the Russian Federation and the grant RFBR №14-02-31223-mol_a.

Fig. 1: HAADF STEM images of the 120 nm BST film. Misfit dislocations marked by arrows and numbers of half-planes between them (a), an enlarged part of the interface with one misfit dislocation and its Burgers vector (b),an enlarged part of the interface and corresponding map with estimated scattered intensities (c) and (d), respectively

Type of presentation: Poster

IT-2-P-2162 A channelling based approach for scattering cross sections of mixed columns in HAADF STEM images

van den Bos K.1, Van Aert S.1
1Electron Microscopy for Materials Science (EMAT), University of Antwerp, Antwerp, Belgium
Karel.vandenBos@uantwerpen.be

HAADF STEM is used to determine structure parameters of nanostructures, such as the number of atoms and the atomic column positions. In order to quantitatively evaluate HAADF STEM images different performance measures including peak intensities and scattering cross sections have been introduced [1-3]. Here, a channelling based approach is proposed to predict these measures for mixed columns.

In the analysis of experimental images performance measures which are sensitive for the parameter of interest are desirable. A comparison between scattering cross sections, determined using statistical parameter estimation theory [2], and peak intensities shows that peak intensities level off at a relatively low number of atoms whereas scattering cross sections increase nearly linearly up to relatively large thicknesses (Fig. 1). This is in agreement with the scattering cross sections computed by using the probe-position integrated cross sections [3]. For that reason, the number of atoms of monotype atomic columns has been successfully determined from experimental scattering cross sections [4]. However, in case of mixed columns the analysis is more complicated since more structure parameters are involved. Therefore, it is desirable to be able to predict performance measures as a function of composition and thickness. Often the assumption of longitudinal incoherence is considered where the scattering intensity of an atomic column is written as the sum of the scattering intensities of the individual atoms constituting this column. However, the non-linear behaviour of peak intensities as well as scattering cross sections makes it impossible to make a valid prediction using this assumption (Fig. 2). A more accurate prediction is obtained based on the channelling theory in which it is assumed that each atom acts as a lens focussing the electrons on the next atom [5]. In this approach the change in scattering intensity with thickness of monotype atomic columns is taken with respect to that of a single atom to estimate the scattering intensity of mixed columns. This approach leads to a significant improvement in the prediction of both performance measures (Fig. 2) and is especially accurate for scattering cross sections. This is an important step forward for the quantitative analysis of complex hetero-nanostructures.

In conclusion, scattering cross sections of mixed columns can be predicted more accurately using a channelling based approach as compared to assuming longitudinal incoherent modelling.

References

[1] Erni et al., Ultramicroscopy 94 (2003), p. 125
[2] Van Aert et al., Ultramicroscopy 109 (2009), p. 1236
[3] E et al., Ultramicroscopy 133 (2013), p. 109
[4] Van Aert et al., Nature 470 (2011), p. 374
[5] Van Aert et al., Ultramicroscopy 107 (2007), p. 551


The authors kindly acknowledge funding from the Fund for Scientific Research, Flanders (FWO).

Fig. 1: Simulations of the scattering cross sections and peak intensities of a single atomic column of (a) Al, (b) Ag, (c) Cd and (d) Pb with respect to thickness. Simulations were carried out using an aberration corrected system with a convergence angle of 21.78 mrad and a detector covering an area of 90-158 mrad.

Fig. 2: Prediction models of simulated scattering cross sections and peak intensities for 17 atom thick mixed columns. In (a) and (c) the centre of an Al column is replaced by Ag atoms keeping the thickness at 17 atoms. In (b) and (d) the centre of a Cd column is replaced by Pb atoms. The parameters for the simulations were the same as in Fig. 1.

Type of presentation: Poster

IT-2-P-2255 Probability of error for counting the number of atoms from high resolution HAADF STEM images

De Backer A.1, De wael A.1, Van Aert S.1
1Electron Microscopy for Materials Science (EMAT), University of Antwerp, Antwerp, Belgium
Annick.DeBacker@uantwerpen.be

During the last years, different quantification methods to count the number of atoms in an atomic column based on HAADF STEM images were developed [1-4]. These methods can be applied to a wide variety of experiments. However, to go beyond the current state-of-the-art, all sources contributing to errors in atom counting need to be understood. Therefore, we discuss a theoretical tool that can be used to quantitatively determine the probability of error for counting the number of atoms.
In principle, expressing the reliability in atom counts can be simplified to discussing the ability of distinguishing between n and n+1 atoms, i.e. the possibility to detect the difference of 1 atom in an atomic column. Using the principles of detection theory [5], this problem is written as a binary hypothesis test with the hypotheses corresponding to atomic columns having n (null hypothesis) and n+1 (alternative hypothesis) atoms. The goal is to minimise the probability of assigning the wrong hypothesis. This is illustrated in Fig. 1. For both hypotheses a so-called log likelihood ratio distribution can be defined. For a given atomic column, the log likelihood ratio then determines which of these hypotheses is decided. If this log likelihood ratio is larger than 0, the alternative hypothesis is decided; otherwise the null hypothesis is decided. From Fig. 1, it is clear that the probability of error is defined by the overlap of log likelihood distributions. This overlap can be calculated numerically.
As a preliminary example, the probability of error is calculated as a function of electron dose and number of atoms using a simple Gaussian model that linearly increases with the number of atoms. The analysis is shown in Fig. 2. As expected the probability of error increases for decreasing electron dose. Furthermore, it is shown that distinguishing between n and n+1 atoms in an atom column becomes more difficult for increasing n. One of the possible applications is to apply this method to realistic simulations in order to optimise the experiment design. This can be realised by minimising the probability of error as a function of a variety of parameters of interest, such as magnification, acceleration voltage, and inner and outer detector angle.
In conclusion, the method quantifies the error for counting the number of atoms as a function of the parameters of interest and enables us to understand the origin of miscounting the number of atoms.

References

[1] Erni et al., Ultramicroscopy 94, 125 (2003)
[2] LeBeau et al., Nanoletters 10, 4405 (2010)
[3] S Van Aert et al., PRB 87, 064107 (2013)
[4] A De Backer et al., Ultramicroscopy 134, p 23 (2013)
[5] den Dekker et al., Ultramicroscopy 134, p 34 (2013)


The authors kindly acknowledge funding from the Fund for Scientific Research, Flanders (FWO).

Fig. 1: The probability of error is depends on the electron dose D, the number of atoms n, the total scattered intensity of a column CS and the pixel size in the simulated HAADF STEM image dx; F denotes the cumulative distribution function of the normal distribution having mean μ and standard deviation σ.

Fig. 2: (a) Probability of error as a function of electron dose for choosing between 10 and 11 atoms in a column (b) Probability of error as a function of number of atoms for a constant electron dose (200 electrons per pixel) (c) Probability of error as a function of number of atoms and electron dose.

Type of presentation: Poster

IT-2-P-2256 STEM Optical Sectioning for Imaging Screw Displacements in Dislocation Core Structures

Yang H.1, Lozano J. G.1, Pennycook T. J.1,2, Hirsch P. B.1, Nellist P. D.1,2
1University of Oxford, Department of Materials, Oxford, UK, 2EPSRC SuperSTEM Facility, Daresbury Laboratory, Warrington, UK
hao.yang@materials.ox.ac.uk

Aberration corrected transmission electron microscopes have advanced our knowledge of the atomic structure of edge dislocations, which are viewed end-on with the tensile or compressive strain normal to the dislocation being clearly visible. Atomic displacements associated with screw dislocations however cannot be observed end-on because the helical screw displacements are parallel to the viewing direction. In this paper the helical displacements around a screw can be imaged with the dislocation lying transverse to the electron beam by “optical sectioning” in annular dark-field scanning transmission electron microscope imaging. In optical sectioning the few nanometer depth of focus is utilized to extract information along the beam direction by focusing the electron probe at specific depths within the sample. This novel technique is applied to the study of the c-component in the dissociation reaction of a mixed [c+a] dislocation in GaN that has previously been observed end-on [1].

Figure 1 shows atomic layers from different depths of a c-type screw dislocation aligned along the c axis [0001] in GaN. Each layer consists of a (2-1-10) plane, and is parallel to the dislocation line. In layers far from the screw dislocation, the displacements vary slowly across the field of view as expected from the lower shear strain that exists further from the dislocation core. Layers close to the screw dislocation core show displacements that very rapidly in the vicinity of the dislocation core, with a rapidly varying shear of the (0002) planes to given an apparent displacement of c/2 across the dislocation core (as expected for a total Burgers vector of c).

A focal series of experimental images were recorded using a Nion UltraSTEM100 aberration-corrected STEM operating at 100 kV (Figure 2). A 1μm thick sample of GaN, grown by metalorganic vapour phase epitaxy on a sapphire substrate, was thinned to be viewed along the a crystallographic axis. A dislocation was found lying in the plane of the sample, and characterized using weak-beam imaging to be of a mixed [c+a] type along [0001]. As the electron beam is focused closer to the dislocation from (a) to (e), the shearing of the (0002) planes becomes more localized in the image, and a more detailed observation of the screw displacements shows that the shearing occurs equally along two distinct lines along [0001], indicated by the arrows in Figure 2. It is therefore apparent that the screw component of the dislocation has dissociated according to the reaction c= c + ½c] confirming the assumption made in previous end-on observations [1,2].

[1] P.B. Hirsch et al., The dissociation of the [a+c] dislocation in GaN, Philosophical Magazine, 93 (2013) 3925.
[2] H. Yang et al., manuscript in preparation.


The authors would like to acknowledge financial support from the EPSRC (grant number EP/K032518/1) and the EU Seventh Framework Programme: ESTEEM2.

Fig. 1: Figure 1. Structure model of a 10nm thick GaN c type screw dislocation viewed at different depths along direction, with the screw dislocation lying in the middle. From the top to the bottom panel, the distances from the dislocation core are +5, +1, +0.3, -0.3, -1 and -5nm, respectively.

Fig. 2: Figure 2. ADF STEM optical sectioning of a [c+a] dissociated screw dislocation viewed perpendicular to the dislocation line along the a direction. The focal step between (a),(c) and (e) is 4nm. (b,d,f) contain Fourier filtered versions of the corresponding images using just the (0002) Fourier components to highlight the shearing of the planes.

Type of presentation: Poster

IT-2-P-2265 Adding the Third Dimension to Atomic Resolution Spectrum Imaging

Pennycook T. J.1,2, Lewys J.1, Cabero M.3, Ribera-Calzada A.3, Leon C.3, Varela M.4,3, Santamaria J.3, Nellist P. D.1,2
1Department of Materials, University of Oxford, Parks Road, Oxford OX1 3PH, UK , 2SuperSTEM Laboratory, STFC Daresbury, Keckwick Lane, Warrington WA4 4AD, UK, 3Grupo de Fisica de Materiales Complejos, Universidad Complutense, 28040 Madrid, Spain , 4Materials Science and Technology Division, Oak Ridge National Laboratory, Oak Ridge, TN 37831, USA
timothy.pennycook@materials.ox.ac.uk

Aberration correction made possible two-dimensional atomic resolution spectrum imaging in the scanning transmission electron microscope (STEM). It also led to a significantly reduced depth of field which has been utilised to perform optical sectioning with atomic number contrast annular dark field (ADF) imaging and determine the positions of individual dopant atoms in three dimensions. Here we combine these corollaries of aberration correction to demonstrate three dimensional elemental mapping with atomic resolution electron energy loss spectroscopy (EELS). Atomic lateral resolution is critical as the optical transfer function of the STEM has a large missing cone, leading to excessive depth elongation for laterally extended objects [1]. The longitudinal resolution varies with the lateral spatial frequency. In regions relevant to lateral resolutions achievable today, the missing cone causes the longitudinal resolution to vary as approximately d/α where d is the characteristic spacing of the object and α is the convergence angle. If the highest spatial frequency resolved is for example the width of a 3 nm nanoparticle, the depth resolution will be around 136 nm with a 22 mrad convergence angle. If instead we were able to resolve atomic columns with a spacing of 0.3 nm and use a 30 mrad convergence angle the depth resolution improves to 10 nm. This relationship between lateral spatial frequency transfer and depth resolution applies both to ADF and EELS imaging. However, although theoretical simulations suggested it was possible to perform optical sectioning with EELS, it had not previously been demonstrated experimentally.

Using a Nion UltraSTEM 100 operated at 100 kV with a 30 mrad convergence angle we acquired successive spectrum images from the same area of a sample, but with the probe focused to different depths. The EELS optical sectioning revealed the presence of a network of yttria-stabilized zirconia (YSZ) islands buried beneath strontium titanate (STO). These regions appear perovskite like from 2D imaging focused at the entrance surface, emphasising the importance of considering the possibility of three dimensional inhomogeneity. The results also highlight the unambiguous nature of EELS elemental mapping, revealing 3D compositional changes that cannot be determined through ADF optical sectioning with complete certainty.

[1] G. Behan et al, Phil. Trans. R. Soc. A 367 (2009), p. 3825.


SuperSTEM is the EPSRC UK National Facility for Aberration-Corrected STEM. Research at ORNL was sponsored by the U.S. DOE, Office of Science, Materials Sciences and Engineering Division (MV). The research leading to these results has received funding from the European Union Seventh Framework Programme under Grant Agreement 312483 - ESTEEM2 (Integrated Infrastructure Initiative–I3).

Fig. 1: Optical sectioning with atomic resolution spectrum imaging. The Sr (M-edge), Ti (L-edge) and Zr (M-edge) maps have been denoised with principle component analysis, and are shown alongside the simultaneously acquired high angle annular dark field (HAADF) images.

Type of presentation: Poster

IT-2-P-2307 Observing depth dependent strain via optical sectioning in the STEM

Lozano J. G.1, Yang H.1, Guerrero-Lebrero M. P.2, Yasuhara A.3, Okunishi E.3, Zhang S.4, Humphreys C. J.4, Galindo P. L.2, Hirsch P. B.1, Nellist P. D.1
1Department of Materials, University of Oxford, Oxford (UK), 2Departamento de Lenguajes y Sistemas Informaticos, CASEM, Universidad de Cadiz, Puerto Real (Cadiz), Spain, 3JEOL Ltd., Tokio (Japan), 4Department of Materials Science and Metallurgy, University of Cambridge, Cambridge (UK)
juan.lozano@materials.ox.ac.uk

The development of spherical aberration correctors for the scanning transmission electron microscopes (STEM) has led to a reduction in the depth of field, which can be of just a few nanometers in a modern instrument. Since this value is smaller than the typical sample thickness, it creates an opportunity to optically section the sample simply by imaging with the focal plane set to a specific depth within the sample [1]. By recording a series of images over a range of focus values, a full three-dimensional image can be obtained. Here we demonstrate that optical sectioning in the high angle annular dark field mode (HAADF)-STEM mode has a sufficiently small depth of field to detect depth-dependent atomic displacements associated with dislocations in GaN, in particular the so-called Eshelby twist [2]. The Eshelby twist is a consequence of the relaxation of the stresses on the free surfaces of the thin TEM sample in dislocations with a screw component of the Burgers vector normal to the foil. It can be seen as an apparent rotation of the lattice in one surface, with displacements that decrease with increasing depth below the surface until the mid-plane of the foil is reached where no displacements should be observed. A rotation in the opposite sense occurs at the opposite surface.
HAADF images of a GaN crystal containing the displacements associated with a right-handed screw dislocation and the Eshelby twist were simulated at different focus value using SICSTEM [3]. The results indicate that the displacements due to the Eshelby twist become larger as the distance from the core increases within the field of view of the simulations, leading to an apparent overall rotation of the lattice around the dislocation core (Figure 1). This allows the twist to be measured on the lattice planes that are parallel to the fast scan direction in the STEM image, providing an approach very robust to the effects of sample drift and scan distortion.
We show that the Eshelby twist can be experimentally measured, by recording focal series in dislocations with a screw component (either pure screw or mixed) imaged end-on. The rotation of the lattice with respect to the image of the crystal at the entrance, as a function of defocus, was measured using the Radon transform [4] (Figure 2), which allowed us to quantify the rotation rate in the different types of dislocations (Figure 3) and us to determine the sign of the screw component of their Burgers vector.
[1] G. Behan et al., Phil. Trans. R. Soc. A 367, 3825 (2009)
[2] J. D Eshelby and A. N. Stroh, Phil. Mag. Series 7 42, 1401 (1951).
[3] J. Pizarro et al. , Appl. Phys. Lett. 93, 153107 (2008).
[4] S. R. Deans, The Radon Transform and Some of Its Applications, New York: John Wiley & Sons (1983)


We gratefully acknowledge financial support from the EPSRC, the Spanish MINECO and the Junta de Andalucía.

Fig. 1: Displacements maps due to the Eshelby twist for focus values a) 0 nm (entrance surface), b) -10 nm (exit surface) for a GaN crystal with a screw dislocation normal to the foil with the Eshelby twist, and the atomic positions of the infinite crystal. For better visualization, a factor of 10 has been applied.

Fig. 2: (a) and (b) show two HR-STEM images of a screw dislocation taken at focus values 4 nm apart, and (c) and (d) their Radon transform maps .(e) and (f) represent the variance of each intensity profile in (c) and (d), where the maximum variance corresponds to the angle at which the horizontal a-planes are aligned with the Radon projection.

Fig. 3: Experimentally measured rotation angle as a function of defocus for three types of dislocations: a pure screw, a pure edge and two neighbouring mixed dislocations. The rotation is clockwise for the screw and one of the mixed dislocations and anticlockwise for the other. No significant rotation is observed for the pure edge dislocation.

Type of presentation: Poster

IT-2-P-2322 Comparison of intensity and absolute contrast of simulated and experimental high-resolution transmission electron microscopy images for different multislice simulation methods

Krause F. F.1, Müller K.1, Zillmann D.1, Jansen J.2, Schowalter M.1, Rosenauer A.1
1Institut für Festkörperphysik, Universität Bremen, Otto-Hahn-Allee 1, 28359 Bremen, Germany, 2National Centre for HREM, Kavli Institute of Nanoscience, Delft University of Technology, Lorentzweg 1, 2628 CJ Delft, The Netherlands
f.krause@ifp.uni-bremen.de

Discrepancies between experimental and simulated images were often reported for high resolution transmission electron microscopy (HRTEM). Simulated contrasts deviate by a factor of up to 3 from experimental ones [1,2]. This disagreement, termed Stobbs factor, has prevented evaluation of HRTEM contrast by comparison to simulations as successfully realised in Z-contrast scanning TEM [3]. The mismatch is caused mainly by improper consideration of the camera modulation-transfer function (MTF) [4]. It was further proposed that contrast-overestimation could be attributed to the use of absorptive potentials (AP) for thermal diffuse scattering (TDS) and to the treatment of incoherence by coherent envelopes [5]. A frozen lattice (FL) simulation with incoherent summation of intensities simulated for various Gaussian-distributed incident angles is more adequate.

The influence of each of these simulation methods on HRTEM contrast was examined by studies of simulated defocus series. Figure 1 shows the results for the simulation of 15 nm thick gold, the proper use of the MTF yields the largest contrast reduction by a factor of 140%. The consideration of TDS by FL instead of AP yields a small contrast decrease below 10%. Incoherent summation of different incident angles contributes a reduction of about 20%. Use of the coherent envelopes instead of the more accurate transmission cross coefficients (TCC) also causes overestimation of image contrast of 10%.

The mismatch of experiments and simulations, conducted with FL, incoherent summation and properly considered MTF, was investigated. Defocus series of a gold foil were acquired with a CS-corrected microscope. Specimen thickness and orientation were determined by diffraction pattern refinements and FL simulations were conducted for these parameters. With an aperture of 7 nm-1 radius, a very good agreement is achieved for image patterns and intensities quantitatively measured in units of incident intensity. The image contrast also coincides as shown in Fig. 2. The ratio of simulated and measured contrast is 0.98±0.07. For larger apertures a discrepancy of 20% is found and good agreement of intensities is observed. Without any aperture the difference amounts to a factor of 40%. Residual aberrations and drift as cause for this were ruled out.

The contrast mismatch between HRTEM simulations and experiments is definitely reduced by proper consideration of the camera MTF and FL simulations with incoherent summation but still remains observable with larger apertures.

[1] M.J. Hÿtch, W.M. Stobbs, Ultramicroscopy 53(1994) 191

[2] A. Howie, Ultramicroscopy 98(2004) 73

[3] J.M. LeBeau et al., Phys.Rev.Lett. 110(2008) 206101-1

[4] A. Thust, Phys.Rev.Lett. 102(2009) 22080-1

[5] D. V. Dyck, Ultramicroscopy 111(2011) 894


This work was supported by the Deutsche Forschungsgemeinschaft (DFG) under contract № RO2057/4-2.

Fig. 1: Contrast of images of a defocus series simulated for 15 nm gold in [100] direction with different techniques using an objective aperture of radius 14 nm-1: The red curve is the result of conventional MS and incoherence treated by coherent envelopes. For the following curves, the image formation was successively simulated more accurately.

Fig. 2: Comparison of experimental and simulated contrast of a defocus series of [100] oriented gold of 12 nm thickness and an objective aperture of 7 nm-1 radius. Both the values and the periodicity agree well.

Type of presentation: Poster

IT-2-P-2407 Studying ω to α Phase Transformation in Ti-15Mo alloy by Combination of Aberration-corrected Scanning Transmission Electron Microscopy and Ab-initio Calculations

Sung Jin Kang 1 Sung-Hwan Kim 1 Heung Nam Han 1 Min-Ho Park 2 Cheol-Woong Yang 2 Hu-Chul Lee 3 Yoon-Uk Heo 3 Miyoung Kim 1
Department of Materials Science & Engineering, Seoul National University, South. Korea 1 Department of Materials Science & Engineering, Sungkyunkwan University, South Korea 2 High Temperature Energy Materials Research Center, Korea Institute of Science and Technology, South Korea 3
kang123@snu.ac.kr

Compared to α -Titanium alloys (Hexagonal), β-Titanium alloys (BCC) are inherently ductile and have promising potentials to substitute new technology materials in daily life1. Mo is commercially added as β-Ti stabilizer to precipitate finely dispersed round shaped α-Ti phase in the β-Ti matrix to enhance the hardness of the Titanium. Interestingly, this α-Ti precipitate is not nucleated directly from the β-Ti phase but from the nucleation sites provided by ω precipitates2. Though there have been intensive studies on the phase transformation of β → α, detailed atomistic dynamics, including the ω phase, have rarely been investigated. We study Ti-Mo(15 wt%) alloy for the phenomena of α-Ti phase formation from the ω precipitate using aberration corrected high annular angle dark field scanning electron transmission microscopy (HAADF-STEM) and electron energy loss spectroscopy for chemical information as well as atomic structural information. We present direct images of the early stage in ω → α transition state exhibiting a metastable state. In bright and dark Z-contrast regions of ω precipitates, the atomic arrangement as well as Z-contrast seems very different from each other. Bright Z-contrast regions show an atomically resolved projected image of ω precipitate crystal structure in the [112 ̅0] zone axis (Fig. 1B), while dark Z-contrast regions show a burry image which is not easy to be interpreted. The image of dark Z-contrast region has a layered periodic pattern and the atoms on each layer are not well resolved as presented in Fig. 1A. The Ti atomic layers could be considered as a metastable phase which will finally develop to a part of α precipitate. We speculate that the metastable structure is formed by distortions caused by local defects of the ω phase. Substantial amount of stacking faults and dislocations in an extremely early stage of ω→ α phase transition supports this hypothesis. Using systematic ab-initio calculations, we found that there is a reasonably stable defective ω-Ti structure which is relaxed to the structure similar to the Z-contrast in HAADF-STEM images (Fig. 2). It is also confirmed that the defective Ti structure is relaxed to the stable hexagonal α-Ti structure with additional Ti atoms on the Ti deficient sites, converting the [112 ̅0]ω phase orientation to [0001]α directly. This study demonstrates that ω → α transition in Ti-Mo alloy system is governed by defect mediated phase transformation. References 1. Ankem S, Green CA. Mater Sci Eng A 263 (1999) 127 2. S. Nag et.al, Acta Materialia 57 (2009), 2136-2147
This work was supported by the National Research Foundation of Korea (NRF) grant funded by the Korea government (MSIP) (NRF 2013034238)
Fig. 1: Enlarged HAADF-STEM image of [11-20]ω//[110]β orientation in rapidly cooled Ti-15wt% Mo sample after aging at 400°C. The Z-contrast of region A is apparently different from the atomic array of ω phase in region B.
Fig. 2: Position of considered layered defects (left) and the Ab-initio calculation result of fully relaxed structure with defects (right).
Type of presentation: Poster

IT-2-P-2425 Sensitive X-ray analysis system on an automated aberration correction FE-STEM

Inada H.1, Hirayama Y.1, Tamura K.1, Terauchi D.1, Namekawa R.1, Shichiji T.1, Sato T.1, Suzuki Y.1, Konno M.1, Nakamura K.1, Hashimoto T.1
1Science & Medical Systems Design Div. Hitachi High-Technologies Corp.
inada-hiromi@naka.hitachi-hitec.com

In recent years the aberration-correction technique has brought a revolution in analytical microscopy by making atomic-resolution imaging and analysis routinely achievable in transmission and scanning transmission electron microscopy (TEM and STEM) 1). We have developed an aberration corrected STEM (Hitachi HD-2700) with an automated aberration correction function. The HD-2700 is equipped with a large solid angle Energy dispersive X-ray spectrometry (EDX) detector which enables an atomic spatial resolution and high sensitivity for EDX analysis.
In the new auto-aberration correction function, the aberration coefficients are measured from a Ronchigram image recorded on a CCD camera for an amorphous sample. Using the coefficients, the software determines the aberrations that need to be corrected and proceeds to correct them. The aberrations are measured and corrected repeatedly and automatically until they are settled under the thresholds. Experiment tests revealed that it took approximately 11 minutes to complete up to the 3rd order aberration2). Figure 1 shows an example of image comparison before and after auto correction for a silicon single crystal imaged along the <110> direction. The Si dumbbell structure is clearly observed after the correction.
To improve the X-ray detection efficiency of the STEM-EDX system, we adopted a design of a windowless3), large area (100mm2) silicon drift detector (SDD). The detector is located closely to the specimen to realize a large solid angle of 1.1sr. The detection sensitivity of the light element (Nitrogen) is more than 10 times higher than that of the 30 mm2 SDD. Figure 2 shows STEM-EDX mapping results for a Pd-Pt catalyst particle specimen using a 30mm2 Si(Li) conventional detector and the new 100mm2 SDD, respectively. The beam energy of 200kV, probe current of 800pA, and acquisition time of 3min. were used. Clearly the maps obtained using the new SDD show much better signal to noise ratio for both nano-particles and carbon support. The high-speed X-ray analysis with the new 100mm2 windowless SDD also largely reduces the beam irradiation damage to the specimen4).

1) H.Inada et al., J. Elec. Microsc., 58 (2009), 111.
2) Y. Hirayama et al., JSPS132 congress (2013).
3) S.Isakozawa et al., J. Elec. Microsc., 59 (2010), 469.
4) K. Tamura et al., Microsc. and Microanal., S2 19 (2013) 1192.

 


Fig. 1: Image comparison between before and after auto Cs correction of silicon dumbbells.

Fig. 2: EDX mapping comparison of Pt and Pd catalyst particles (a) Conventional 30mm2 Si(Li), (b) Newly developed 100mm2 SDD.

Type of presentation: Poster

IT-2-P-2453 3D-structural elucidation of highly ordered mesoporous TiO2 thin film by the method of electron crystallography

Xu B. B.1,2, Feng Z. D.1, Zhou H.1, Wang C.1
1College of Materials, Xiamen University, Xiamen, China, 2Center of Instrumental Analysis, College of Chemistry and Chemical Engineering, Xiamen University, Xiamen, China
zdfeng@xmu.edu.cn

At present, transmission electron microscopy 3D reconstruction has become an important technique to elucidate the 3D structure of materials. Electron tomography, single particle analysis and electron crystallography are the common methods of 3D reconstruction. Thereinto, electron crystallography, comprising transmission electron microscopy and electron diffraction, determines the 3D structure of materials via their 2D structure information. This method is used to investigate the structure of crystal in general. In fact, due to the similar structure between highly ordered and crystal, highly ordered mesoporous structure can also form electron diffraction patterns, so the method of electron crystallography can be used to study the 3D structure of highly ordered mesoporous TiO2 thin film as well. In this experiment, we obtained TEM images and their corresponding slected-area electron diffraction (SAED) patterns from different zone axes. Especially, the corresponding SAED patterns were recorded at an instrument camera length of 200 cm, which is much longer than the average length. The top view TEM image of the film shows a highly ordered honeycomb arrangement with a nearly perfect hexagonal disposition. Its corresponding SAED pattern exhibits a 6-fold symmetry, which is compatible with the [001] zone axis of the hexagonal structure. Mainly, the diffraction patterns taken from three directions ([001], [1-10] and [121]) can be indexed in the P63/mmc space group. And the cross-section TEM image regarded as viewed from [100] zone axes shows an ABAB stacking sequence. Moreover, the diameter of the pores can be directly measured to be ~10nm. In summary, the method of electron crystallography implements an effective explanation of highly ordered mesoporous TiO2 thin film with 3D hexagonal structure.


Fig. 1: The TEM images along [001] with SAED

Fig. 2: The TEM images along [1-10] with SAED

Fig. 3: The TEM images along [121] with SAED

Fig. 4: The cross-section TEM images along [100]

Type of presentation: Poster

IT-2-P-2571 Optimization of imaging conditions for atomic resolution in Titan TEM to minimize radiation damage and to study low angle boundaries in graphene-like materials

Lopatin S.1, Chuvilin A.2
1FEI Company, Eindhoven, Netherlands, 2CIC nanoGUNE, Donostia - San Sebastian, Spain
sergei.lopatin@fei.com

   Recent advances in spherical aberration (Cs) correction for TEMs in a combination with monochromated electron sources enabled imaging of single and bilayer graphene with atomic resolution [1]. Newly developed TEM techniques such as a single atom or single-atomic-column spectroscopy [2, 3] and atomic resolution electron tomography [4] drive the need for increased electron radiation doses applied to samples. The radiation damage started to be the key limitation factor for high-resolution TEM [5].

   For graphene-like (light element) materials [6] the radiation dose limitation is particularly severe. First, the knock-on damage cross section is higher for low atomic number elements [7]. Second, light elements produce less contrast than heavier elements, so even higher doses are needed to obtain a sufficient signal-to-noise ratio (SNR). Finally, the graphene-like materials appear in the form of low dimensional allotropes that have only one or a few atoms in a typical projection of a high-resolution image.

   To minimize the electron dose the optimization of acquisition parameters is needed. Here we present an extensive study of TEM tuning to obtain high quality HRTEM images of graphene. We used a Titan TEM (FEI Co) equipped with a Cs image corrector, a super-high brightness gun and a monochromator (energy spread better than 0.15eV). Tuning of the Cs corrector is based on measurement of images defocus (df) and astigmatism while recording so-called Zemlin tableau [8]. It was demonstrated that proper accounting for Cs of 3rd and 5th order (C3 and C5) and systematic error of C3 measurement results in more than 2 times increase of contrast, meaning more than 4 times decrease in dose needed for the same SNR (Fig.1).

   The optimal settings found were applied to study low angle boundaries (LAB) in graphene. LAB is a row of edge dislocations, separation of those defining the boundary angle. LABs are not visible directly on the image but can be identified by methods such as geometrical phase analysis (GPA), see Fig.2. Physically LAB may be interesting as they represent a perfect discontinuous layer with periodically spaced singularities.

[1] K. W. Urban, Nature Materials, 10 (2011) 165.
[2] P. E. Batson, Nature (London), 366 (1993) 727.
[3] D. A. Muller, et al, Science, 319 (2008) 1073.
[4] M. B. Sadan, L. Houben, S. G. Wolf, A. Enyashin, G.Seifert, R. Tenne, and K. Urban, Nano Lett., 8 (2008) 891.
[5] R. F. Egerton, P. Li, and M. Malac, Micron, 35 (2004) 399.
[6] K. S. Novoselov, A.K. Geim, S.V. Morozov, D. Jiang, Y. Zhang, S.V. Dubonos, I.V. Grigorieva, and A. A. Firsov, Science, 306 (2004) 666.
[7] F. Banhart, Rep. Prog. Phys., 62 (1999) 1181.
[8] F. Zemlin, K. Weiss, P.Schiske, W. Kunath, K. -H. Herrmann, Ultramicroscopy 3 (1978) 49.


Fig. 1: Simulation verification of the impact of optimum conditions for 0.1nm transfer: a) Scherzer conditions optimized; b) C5+C3+df optimized; c) C5+C3+df optimized and systematic error from Zemlin tableau is accounted; d) the intensity profiles across simulated images; e) an experimental image acquired at approximately optimum conditions.

Fig. 2: LAB in graphene: a) original HRTEM image; b) dislocations identification by GPA (rotation map).

Type of presentation: Poster

IT-2-P-2594 Aberration correction through auto-iteration system utilizing diffractogram analysis by profile fitting technique

Morishita S.1,3, Nakamichi T.1, Takano A.1, Satoh K.1, Hosokawa F.1, Suenaga K.2,3, Sawada H.1,3
1JEOL Ltd., 2National Institute of Advanced Industrial Science and Technology, 3Research acceleration program, Japan Science and Technology Agency
shmorish@jeol.co.jp

Spherical-aberration-corrected TEM/STEM has become widely used in the past decade. To automatically correct the aberrations in the correction system, a precise measurement of residual aberrations and an optimized correction procedure are crucial. Several methods have been reported for quantitative measurement of the aberrations [1-5]. We have developed corrector control software JEOL COSMO (Corrector System Module), in which aberrations are measured by diffractogram tableau method in TEM and SRAM method [6] in STEM. In the diffractogram tableau method, measurement precision for measurable defocus (df) and two-fold astigmatism (A2), at diffractograms with tilted illuminations, determines the final precision of the correction, since residual aberrations and the intrinsic A2 and df to be corrected are calculated from these measurable parameters. This paper reports a profile fitting technique to analyze the diffractograms incorporated into our developed auto-iteration system, which enables us to correct aberration with an improved precision.
In the diffractogram analysis, radial intensity profiles are used. Each of the radial profile is fitted with a phase contrast transfer function to pick up a parameter of the first-order components, that is, amount of defocus in particular azimuth. For searching the local minima in the profile that determine the parameters of the transfer function, profiles around local minima instead of simple detection of local minima are utilized to reduce affection of noise on the profile in our system. With thus obtained first-order components at many azimuths, the intrinsic df, A2 and other aberrations are calculated. Figure 2 compares the plots of A2 obtained using only the position of first zero and using the devised fitting method over 20 diffractograms. With this method, standard deviation of measured intrinsic A2 is improved from > 1 nm to a few angstroms.
Next, we developed the auto-iteration system for aberration correction using a script language integrated in the JEOL COSMO. The system automatically chooses the next targets of aberration to be corrected and corrects them. Our algorithm for correcting procedure preferentially corrects aberrations of lower-order or large higher-order to minimize the phase disturbance. Finally, we successfully performed the auto aberration correction in TEM with the improved procedure, which results in the residual third-order aberrations from about 10 μm to < 1 μm within 15 min.

[1] F. Zemlin, et al., Ultramicros. 3, 49 (1978).
[2] S. Uhlemann, et al., Ultramicrosc. 72, 109 (1998).
[3] M. Haider, et al., Ultramicrosc. 81, 163 (2000).
[4] J. Barthel, et al., Ultramicrosc. 111, 27 (2010).
[5] M. Vulovic, et al., Ultramicrosc. 116, 115 (2012).
[6] H. Sawada, et al., Ultramicrosc. 108, 1467 (2008).


This work is supported by Japan Science and Technology agency, Research Acceleration Program.

Fig. 1: Example of a diffractogram and its line profile. The experimental profile indicated by solid line is used for profile fitting.

Fig. 2: (a) Results of two fold astigmatism measurements by using positions of first zero (gray) and by using profile fitting (black). (b) Standard deviation of (a), which includes both measurement error and actual fluctuation.

Type of presentation: Poster

IT-2-P-2830 Quantitative Low-dose HRTEM Imaging and Analysis of Radiation-sensitive Materials

Huang C.1, Borisenko K. B.1, Kim J. S.1, Berkels B.2, Kirkland A. I.1
1University of Oxford, 2RWTH Aachen University
chen.huang@spc.ox.ac.uk

It is well-known that radiation damage caused by fast electrons in the electron microscopes is a main obstacle for high resolution transmission electron microscopy (HRTEM) characterisation of materials that are easily damaged by the exposure to electrons. Although extensive research has been carried out on damage mechanisms, critical doses, and low-dose imaging techniques for decades, the search for the ultimate resolution for radiation-sensitive materials imaging and the most optimal imaging conditions is still continuing.

Due to the usually simultaneous existence of more than one kind of damage mechanism and the innate complexity of each damaging process, theoretical predictions of the dose-limited resolution for many materials are still only qualitative. When it comes to the design of quantitative low-dose experiments, whether it is single-shot imaging or image series acquisition, a more accurate knowledge of the effects of dose rate, total dose, accelerating voltage and microscope aberrations on the resolution is needed to achieve the optimal resolution for a particular sample.

In this work we demonstrate an experimental approach to determining the dose-limited resolution of radiation-sensitive materials. Apart from the established low-dose imaging methods, we apply multiple related techniques, such as exit wave reconstruction and non-rigid image registration to improve the quantitative data analysis and interpretation. It is shown that with careful calibration, the suggested quantitative low-dose high resolution imaging and data processing procedures should be easily adaptable to any specific transmission electron microscope equipped with standard instrumentation.


Type of presentation: Poster

IT-2-P-2922 Real Space Characterization of the Finite Shape of the STEM Probe

Grimley E. D.1, Sang X.1, LeBeau J. M.1
1North Carolina State University, Department of Materials Science and Engineering, Raleigh, North Carolina, United States
edgrimle@ncsu.edu

The shape and size of the electron probe impact many aspects of scanning transmission electron microscopy (STEM), including resolution in imaging and related STEM spectroscopies. This work introduces the projective standard deviation (PSD) as a tool for examining the shape of the STEM probe from atomic resolution images of a reference material. The PSD possesses sensitivity to the shape and intensity of atom columns in a STEM image and has already proven pivotal in the recently developed Revolving-STEM (RevSTEM) technique for eliminating drift related image distortions [1].

The PSD utilizes the normalized Radon transformation mathematical construction that projects the normalized integrated intensities of an image onto an orthogonal vector; this transformation oriented along a lattice vector results in a profile with sharp, periodic oscillations while a projection away from a lattice vector forms a profile with nearly flat intensity. Fig. 1 (a) shows transformations over a 180° range for a simulated Si <100> image. Fig. 1 (b) displays the transformations at 45° (red oscillating line) and 60° (blue dashed line with comparatively flat intensity) from Fig. 1 (a) which are oriented along and away from lattice vectors, respectively. The standard deviation of a normalized Radon transformation over a desired angle range generates a PSD plot with standard deviation as a function of angle. Projection perpendicular to lattice vectors results in large standard deviations, while angles oriented away from lattice vectors result in small standard deviations, as exemplified by the peaks and flat regions of the PSD plot in Fig 2 (a).

We demonstrate herein that the comparison between the PSD of an experimental image and the PSDs of simulated images enables the identification of the probe shape for a given microscope and image. Fig. 2 (a) illustrates this principle as the PSD of a simulated Si <100> image convolved with a Gaussian function (FWHM 0.10 nm by 0.10 nm; red line) differs substantially from the PSD of the same simulated image convolved with an astigmatic Gaussian function (FWHM 0.10 nm by 0.07 nm; dashed blue line). Iterative least-squares fitting of PSD plots allows determination of a probable probe shape for given conditions. Fig. 2 (b) shows the fitting of a PSD of a background subtracted experimental Si <100> RevSTEM micrograph (red line) with the PSD of a simulated Si <100> image convolved with a Gaussian function of FHWM 0.102 nm by 0.092 nm and ~45° rotation (blue dashed line). Strikingly, the simulated Si <100> image with added noise (Fig. 2 (d)) is almost identical to the background noise subtracted experimental RevSTEM Si <100> image (Fig. 2 (c)).

Reference: [1] X Sang and JM LeBeau, Ultramicroscopy 138 (2014), p. 28


The authors acknowledge the use of the Analytical Instrumentation Facility (AIF) at North Carolina State University, which is supported by the State of North Carolina and the National Science Foundation, and they acknowledge use of the High Performance Computing (HPC) services at North Carolina State University.

Fig. 1: (a) Normalized Radon transformation spanning 180° of a simulated Si <100> image (~ 3 x 3 cells; 240 x 240 pixel) convolved with a Gaussian (FWHM 0.10 nm by 0.10 nm). (b) 45° (red) and 60° (dashed blue) projection angle line profiles from (a) showing oscillation at 45° and comparatively flat intensity at 60° due alignment with a lattice vector.

Fig. 2: (a) PSD of simulated Si <100> convolved with symmetric (red) and astigmatic (blue dashed) Gaussian functions. PSD (red trace in (b)) and image of (c) a background noise subtracted RevSTEM Si <100> image and PSD (dashed blue trace in (b)) and image of (d) simulated Si <100> image (noise added to (d)).

Type of presentation: Poster

IT-2-P-2932 Complementary Nature of Microscopy Techniques for Understanding Materials Phenomena

Ghosh C.1, Basu J.1, Divakar R.1, Mohandas E.1
1Materials Synthesis and Structural Characterisation Division, Physical Metallurgy Group, Metallurgy and Materials Group, Indira Gandhi Centre for Atomic Research, Kalpakkam-603102, Tamil Nadu, India
chanchal@igcar.gov.in

Advent of nanotechnology over last couple of decades has not only reduced the size of basic building block of materials by several orders of magnitude, complexity of materials architecture and chemistry has also increased enormously. In such a scenario, judicious application of suitable microscopy technique can only provide answer to a particular question. Even though, resolution limits, capabilities and instrumentation of microscopes have improved to a large extent, relatively older techniques even today are found to be equally relevant. In this paper, relative merits of several microscopy techniques will be compared with reference to materials problems encountered in our laboratory.
Understanding anti phase domain boundaries in polar crystals has been a major challenge. Till date, all the interface models for APB are based on (10-10) interfaces. However, anti phase domains apart from this interfaces and complex interaction of these interfaces with other defects are observed very frequently. It is almost impossible to suitably image such boundaries at atomic resolution as crystallography of the interfaces is not known. Diffraction contrast imaging is probably the starting point for understanding APB in polar crystals. In case of phase contrast microscopy the contrast is generated by the interaction of specimen potential with the incident e- wave which is further modified with the instrument CTF. Whereas, during STEM a very fine e- probe scans over the atomic columns and during scanning either electron energy-loss spectra or the X-ray signal is used to understand the chemistry. A number of complex oxides have been studied by this method. During incoherent imaging of complex oxides, heavier cations act as strong scattering centers while the relatively lighter oxygen anions scatter weakly. So the image contrast is mostly dominated by the scattering from the cations. The structural imaging of complex intermetallics e.g. V-doped TiCr2 Laves phase is quite different in nature from the complex oxides. The atomic numbers of V, Ti and Cr are pretty close and all of them will scatter almost equally. As a result differential contrast as is generated in a complex oxide will not happen for V-doped TiCr2 phase. Zero loss phase contrast microscopy with image simulation has been proven useful for structural imaging of Ti and Cr atomic columns and also providing V occupancy information.
Though all of these techniques fetch the materials information at the atomistic scale, still each one of them is unique by electro-optical configuration and interaction with the materials. In another words all of these techniques are complementary to each other and only a combination of all of these techniques can provide the complete solution of the materials related issues.


The authors would like to acknowledge UGC-DAE-CSR for providing the experimental support.

Type of presentation: Poster

IT-2-P-2949 Development of highly stabilized water chiller for atomic resolution microscope

Hamochi M.1, Ishii T.1, Chisaka S.1, Okunishi E.1, Sawada H.1, Wakui S.2
1JEOL ltd., 2Tokyo University of Agriculture and Technology
hamochi@jeol.co.jp

With progress of high resolution imaging using an aberration corrected STEM/TEM, fluctuation less than 50 pm is detectable at a high magnification. As electrical and mechanical stability of the microscope has been improved, the disturbance, due to change of environmental conditions, becomes more crucial at the higher magnification. Among the conditions, a temperature of cooling water is extremely important because it directly affects the temperature of a lens or a column. Magnitude of the temperature fluctuation on a specimen can be roughly estimated as follows. Thermal expansion coefficients for typical metallic materials are in order of 10-5, so that the temperature fluctuation less than 0.1 degree C changes a length of nm-order for mm-size structures in a microscope. Thus, a highly stabilized water chiller system is desired for an atomic resolution electron microscope. The requirements for the microscope are temperature fluctuation of ± 0.05 degree C/min or less, and water temperature drift less than 0.2 degree C/h. We report a water chiller system developed by us with precise temperature control to realize very small fluctuation of water temperature.

Figure 1 (a) shows appearance of the developed water chiller. The size of the chiller is W550×L775×H1350 mm and the cooling power is 6 kW. The high stability of temperature for the cooling water was achieved by a developed temperature compensator with plural heat sources and an improved heat exchanger. Figure 1 (b) plots a temperature change of flowing water controlled by our system for an ultrahigh resolution Cs-corrected 300-kV microscope operated at 300 kV for an hour. The result shows that the fluctuation was less than ±0.01 degree C/min expressed as difference of maximum and minimum within a 60 sec time window, and the temperature drift was less than 0.006 degree C/h expressed as moving average deviations with 60 sec time window.

We recorded high-resolution images with long acquisition times to evaluate the performance of the chiller. Figures 2 (a) and (b) show HAADF STEM images of a Si[110] with an acquisition times of 10 s and 80 s. Fig. 3 (a) shows a high resolution STEM image of Si3N4 of 4k × 4k pixels with 160 s. The images with long acquisition times showed small distortion due to small sample drift, indicating that the temperature of flowing water was sufficiently stabilized for an atomic resolution imaging with long acquisition time. It should be noted that electrical and mechanical stabilities for this microscope were also devised against the other environmental disturbances such as external magnetic field, an external temperature change and so on. As we reported, we have successfully developed a highly stable water chiller that is compatible for an atomic resolution microscopy.


Fig. 1:  (a) Appearance of the developed precise water chiller, (b) Example of temperature measurement for a flowing water cooling an ultra high resolution Cs-corrected 300kV TEM/STEM

Fig. 2: HAADF STEM image of Si[110], (a) with acquisition time 10 s, (b) 80 s (512 x 512 pixels)

Fig. 3:  HAADF STEM image of Si3N4. Inset image is magnified from area indicated by dotted rectangle.

Type of presentation: Poster

IT-2-P-3171 Quantitative study of defocus-dependent annular bright field images

Lee S.1,2, Oshima Y.2,3, Takayanagi K.1,2
1Tokyo Institute of Technology, Tokyo, Japan, 2JST-CREST, Tokyo, Japan, 3Osaka University, Ibaraki, Japan
slee@surface.phys.titech.ac.jp

Previously, we found that the lithium column intensity of an annular bright field (ABF) image varied by a step of a single lithium atom in correlation with the thickness change of the LiV2O4 crystal [1]. But, ABF imaging mechanism has not been investigated quantitatively. In this study, we observed ABF and high angle annular dark field (HAADF) images of very thin specimen simultaneously and investigated defocus dependency of visibility of atomic columns [2].
By using a spherical aberration corrected electron microscope (R005), both ABF and HAADF images were taken simultaneously for very thin lithium manganese oxide, LiMn2O4 specimen from the [001] view direction. The incident convergent semi-angle was 30 mrad, and the detector semi-angles were 15-30 mrad for ABF and 102-272 mrad for HAADF. Fig.1 shows the through focus series of ABF and HAADF images obtained from 10 nm over-focus to -10 nm under-focus condition. In the ABF images, the atomic column had dark contrast at over-focus, but the contrast reversed into bright one when the defocus condition was changed to under-defocus. While, the HAADF image showed the bright column contrast which did not reverse regardless the focus change. It indicates that ABF image is a kind of phase contrast image, and could be explained by weak-phase-object approximation (WPOA).
We measured visibility in the ABF and HAADF images in order to estimate the defocus for obtaining the maximum contrast and the depth of focal (DOF). The optimum defocus was different between both images: the maximum contrast of the atomic columns was obtained at a few nm over-focus in the ABF image, while it, at a few nm under-focus in the HAADF image. And, DOF was determined to be 4 and 8 nm in the ABF and HAADF image, respectively. DOF of ABF image is obviously narrower than one of HAADF image. ABF imaging which has a narrow DOF could be used for visualizing light elements three-dimensionally.

[1] S. Lee, et al., J. Appl. Phys. 109 (2011) 113530.
[2] S. Lee, et al., Ultramicroscopy 125 (0), 43-48 (2013).


Fig. 1: (a) A structure model of very thin lithium manganese oxide (LiMn2O4) specimen. Through focus (b) ABF and (c) HAADF images are shown from 10 nm (over-focus) to -10 nm (under-focus). The maximum contrast is obtained at the defocus indicated by red rectangle.

Type of presentation: Poster

IT-2-P-3187 Quantitative analysis of CeO2 and Gd-doped CeO2 nanocrystals by HRTEM focal series restoration

Stroppa D. G.1 2, Dalmaschio C. J.3, Thust A.2, Lentzen M.2, Barthel J.2 4, Houben L.2
1International Iberian Nanotechnology Laboratory (INL), Braga, Portugal, 2Forschungszentrum Jülich, Jülich, Germany, 3Federal University of Espírito Santo, São Mateus, Brazil, 4Rheinisch-Westfälische Technische Hochschule Aachen (RWTH), Aachen, Germany
daniel.stroppa@inl.int

High resolution transmission electron microscopy (HRTEM) is a well-known characterization technique with atomic resolution imaging capability, and different approaches have been explored to extend its use with aim at the retrieval of quantitative information with high spatial resolution. A particular example is the restoration of the electron exit-plane wave-function (EPWF) from HRTEM focal series, as it ideally contains information on the atomic species and their position along individual atomic columns [1]. In this work, we explore the phase shifts on restored EPWFs from CeO2 and Gd-doped CeO2 nanocrystals aiming to extract their local chemical composition with atomic resolution.
HRTEM focal series from nanocrystals in <100> zone axis orientation were obtained using an aberration corrected TEM microscope at 300 kV under a negative Cs-imaging (NCSI) condition [2]. This procedure allowed the imaging of O and Ce-Gd atomic columns with enhanced contrast, as shown in Figures 1a and 2a. Electron EPWF from the two investigated samples were restored from the focal series taking into account the residual aberrations, the instrument instabilities and the detection system (CCD) modulation transfer function (MTF) [3]. Figures 1b and 2b show the respective electron EPWF phases reconstructed for CeO2 and Gd-doped CeO2 nanocrystals. Finally, the local EPWF phase shifts corresponding to the individual atomic columns centers were analyzed and compared to results of multislice image calculations.
The results show that the EPWF phase shift differences between columns containing heavy atoms (Ce or Ce-Gd) are significant with respect to the background, indicating that they can be used to map the atomic columns thicknesses and to infer nanocrystals 3D morphology [4]. Even though the phase shift differences between columns containing light atoms (O) are appreciable, they are on the same range of the background phase fluctuation. This indicates that spurious contributions to the EPWF, probably related to the carbon support film, the CCD read-out noise and the phase-shift tail from heavy elements, limit the direct extraction of quantitative information for the current experimental setup.


The authors would like to thank Dr. Antonio J. Ramirez and Prof. Edson R. Leite for the fruitful discussions that contributed to this project.

Fig. 1: Figure 1: a) HRTEM image and b) phase from the reconstructed EPWF after a HRTEM focal series from a CeO2 nanocrystal. While the direct identification of Ce and O atomic columns can be ambiguous from a HRTEM image, the reconstructed EPWF phase clearly shows Ce atomic columns with brighter intensity with respect to the O atomic columns.

Fig. 2: Figure 2: a) HRTEM image and b) phase from the reconstructed EPWF after a HRTEM focal series from a Gd-doped CeO2 nanocrystal.

Type of presentation: Poster

IT-2-P-3197 Temperature dependence of Z-contrast in InGaN

Mehrtens T.1, Schowalter M.1, Tytko D.2, Choi P. P.2, Raabe D.2, Hoffmann L.3, Jönen H.3, Rossow U.3, Hangleiter A.3, Rosenauer A.1
1Institute of Solid State Physics, University of Bremen, Bremen, Germany, 2Max-Planck-Institut für Eisenforschung GmbH, Düsseldorf, Germany, 3Institute of Applied Physics, TU Braunschweig, Braunschweig, Germany
mehrtens@ifp.uni-bremen.de

Scanning transmission electron microscopy (STEM) combined with a high angle annular dark field detector (HAADF) gives rise to image contrast strongly depending on the nuclear charges of the scattering specimen atoms and is often referred to as Z-Contrast. The comparison of HAADF-STEM image intensities with simulated intensities from multislice calculations allows determining specimen thickness or material composition for each atomic column [1, 2] in a high resolution HAADF-STEM micrograph.

The main contribution to the measured signal in HAADF-STEM stems from thermal diffuse scattering (TDS) due to the thermal vibrations of the specimen atoms. An increase of temperature should result in a higher HAADF-signal due to the larger amount of generated TDS.

In this contribution we have studied the thermal dependence of Z-contrast for InGaN/GaN and specimen temperatures between 300K and 600K. Fig. 1 shows the intensity profile of an five-fold InGaN/GaN multi-quantum well structure measured at different temperatures. We observed an increase of the HAADF-STEM intensity with increasing temperature for GaN as well as for InGaN. However, we also noticed that the material contrast in the image (ratio between intensities for InGaN and GaN) decreased with temperature (Fig. 2).

In order to understand this effect, multislice simulations were carried out in the frozen phonon approach for different specimen temperatures using the STEMsim program [3]. A temperature dependent parameterization of Debye-Waller-Factors derived from density functional theory calculations [4] was used. The simulated material contrast in dependence of specimen thickness is shown in Fig. 3 for an indium concentration of 30% and for three different temperatures. The contrast decreases with temperature as is it was found in the experiment and is related to the effect of static atomic displacements (SAD). SADs occur, if atoms with different covalent radii share the same crystal lattice or sublattice., which is the case for In and Ga in InGaN. A rise of temperature increases the contribution of thermal diffuse intensity, whose material contrast is smaller than that of Huang-scattering [5] caused by SADs. Thus, the material contrast decreases with increasing temperature.

[1] Rosenauer et al., Ultramicroscopy 109, 1171-1182 (2009)
[2] Rosenauer et al., Ultramicroscopy 111, 1316-1327 (2011)
[3] Rosenauer and Schowalter, Springer Proc. in Phys. 120, 169-172 (2007)
[4] Schowalter et al., Acta Cryst. A 65, 227-231 (2009)
[5] Z. L. Wang, Acta Cryst. A 51, 569-585 (1995)


This work was supported by the Deutsche Forschungsgemeinschaft under Contract No. RO 2057/8-1 and the Bundesministerium für Bildung und Forschung (BMBF) in the frame of the “ERA-SPOT True Green (13N9634)” project.

Fig. 1: HAADF-STEM intensity profile of a five-fold InGaN/GaN multi-quantum well structure grown on GaN measured at specimen temperatures of 300K, 450K and 600K. The intensity is increasing with increasing temperature.

Fig. 2: Intensity profile of an InGaN quantum well normalized with respect to the intensity of the neighboring GaN for different specimen temperatures.

Fig. 3: Simulated material contrast (IInGaN/IGaN) of In0.3Ga0.7N for different specimen temperatures.

Type of presentation: Poster

IT-2-P-3202 Sub-angstrom resolution realized with super high-resolution aberration corrected STEM at 300 kV

Sawada H.1, Okunishi E.1, Shimura N.1, Satoh K.1, Hosokawa F.1, Kaneyama T.1
1JEOL Ltd.
hsawada@jeol.co.jp

An aberration corrected scanning transmission electron microscopy (STEM) enables us to perform a structural analysis at sub-angstrom resolution [1-3]. By high angle annular dark field (HAADF) STEM method, a resolution of 47 pm was achieved using a Ge [114] [4,5]. For light elements, sub-angstrom distance between Si-N atomic columns in a β-Si3N4 was resolved by an annular bright filed (ABF) imaging technique [6]. Recently, we developed a 300-kV super high-resolution aberration corrected microscope. The stability of the electronic power supplies, mechanical stiffness, and optical parameters such as chromatic aberration coefficient were improved in the microscope. The microscope was equipped with a cold field emission gun to realize high brightness and smaller energy spread. In this paper, we report the results on observations of atomic dumbbells separated by sub-angstrom with the developed microscope.

We observed a GaN [211] [2,7]. Next, sub-50 pm imaging was performed using a Ge [114] and a Si [114] [7]. Figures 1(a, d) and 1(e, f) show HAADF and ABF images of GaN [211]. The Ga-Ga atomic columns separated by 63 pm was resolved by HAADF STEM in Fig. 1(a). Spatial information better than (63 pm)-1 was confirmed in the Fourier transform shown in Fig. 1(c). The intensity profile in Fig. 1(d) shows that 63-pm separated atomic dumbbell of the Ga-Ga was clearly resolved. The ABF was utilized for a light element imaging with sub-angstrom resolution. N-N atomic dumbbell with 63-pm separation was resolved as a gray contrast in Figs. 1(e) and 1(f). The intensity profiles in Figs. 1(e, f) show that the 63-pm is clearly resolved.

Figure 2(a) shows a HAADF STEM image of a Ge [114], which shows 47-pm separated Ge-Ge dumbbells. The Fourier transform in Fig. 2(d) shows -8-84 spots, which correspond to (47 pm)-1. The line profiles in Fig. 2(c) show separations of 47 pm. Next, we challenged a Si [114], which shows 45-pm separated Si-Si dumbbells. The resolution has never been reported. Fig. 2(f) shows HAADF image of the specimen. The -8-84 spots and the line profiles in Figs. 2(d) and 2(h) confirm the resolution of 45 pm.

In conclusion, we have successfully demonstrated that light element imaging with sub-angstrom resolution by ABF, and sub-50 pm resolution by HAADF with the developed super high-resolution microscope at 300kV with CFEG.

References:

[1] P. Nellist, et al., Science 305: 1741 (2004).

[2] H. Sawada et al., Jpn. J. Appl. Phys. 46: L568 (2007).

[3] C. Kisielowski et al., Microsc. & Microanal. 14: 469 (2008).

[4] H. Sawada et al., J. Electron Microsc. 58: 357 (2009).

[5] R. Erni, et al., Phys. Rev. Lett. 102: 096101(2009).

[6] E. Okunishi, et al., Micron 43: 538 (2012).

[7] M. O'Keefe et al., J. Electron Microsc. 54: 169 (2005).


The authors thank Professor Y. Ikuhara, and Associate Professor N. Shibata (The University of Tokyo) for collaboration and instrumental supports.

Fig. 1: (a) Raw HAADF STEM image of GaN [211] taken at 300 kV. The convergence semi-angle was 30 mrad. (b, c) Intensity histogram and Fourier transform of (a). (d) Raw HAADF image and intensity profile from dotted rectangle area. (e,f) Raw and filtered ABF images and their intensity profiles.

Fig. 2: (a, b, f, g) Raw and low-pass filtered dark field STEM images of Ge [114] and Si [114] taken at 300 kV with simulated images. (c, h) Intensity profiles from the dotted rectangles in (a, b) and (f, g). (d, e) Fourier transforms and intensity histograms from (a) and (f).

Type of presentation: Poster

IT-2-P-3462 The mini-TEM: highquality imaging and analysis of biological specimen

Sintorn I.1,2, Kylberg G.2, Nordström R.2, Uppström M.2, Danielsson K.2, Fulin J.2, Åkesson J.2, Coufalova E.3, Drsticka M.3, Kolarik V.3, Stepan P.3
1Centre for Image Analysis, Uppsala University, Sweden, 2Vironova AB, Gävlegatan 22, Solna, Sweden, 3Delong Instruments, Brno, Czech Republic
ida.sintorn@it.uu.se

 

Traditional transmission electron microscopes are bulky and complex machines that are mostly operated by trained specialists. In this paper, we introduce the miniTEM, shown in Fig. 1, a desk-top instrument designed for imaging of biological as well as inorganic samples. The miniTEM has a high degree of automation in the microscope alignment, image acquisition, and analysis processes. The idea with the miniTEM is a small, cheap, robust, and easy to use system that requires no more training than any simple lab equipment, and can be hosted in any office or lab (even a mobile lab). Here we illustrate the imaging possibilities and show that it is good enough for quantitative and qualitative analysis.

The miniTEM microscope runs at 25 keV, which enables high-quality imaging of biological samples with a thickness up to at least 100 nm. The height of the microscope is only 70cm and it can sit and run on any desk in any lab or office space. It achieves a resolution sufficient for tasks such as virus identification in clinical samples, and morphological nanoparticle analysis. In addition to TEM functionality, the miniTEM can also run in STEM (scanning transmission electron microscopy) mode. We present here the achieved parameters of resolution, applicable sample thickness and image signal collection efficiency in both operating modes. One of the first images of inorganic nanoparticles acquired with a miniTEM microscope prototype in the TEM mode is shown in Fig. 2.

 

The graphical user interface is divided into three main views: live, edit, and analysis. It is developed for Windows 8, and designed for a touch screen, allowing convenient scrolling over the sample and zooming in (changing magnification). The live view, illustrated in Fig. 3, is used when manually investigating the sample by moving around, changing magnification and acquiring images. The edit view is for manually marking, drawing, measuring and annotating objects in the images. In the edit view the user can also manually correct analysis results, i.e., remove, add, and rename objects. The analysis view is where the user creates and applies automated image acquisition and/or image processing (GPU-accelerated) and analysis scripts. A graph-based interface is used to create scripts, which can be saved for future use and applied to multiple images.


This work is part of the miniTEM project funded by EU and EUREKA through the Eurostars programme. 

Fig. 1: The miniTEM microscope

Fig. 2: Image of inorganic nanoparticles acquired with the miniTEM, scalebar 200nm.

Fig. 3: The live-view in the miniTEM graphical user interface. In the default layout, thumbnails of acquired images, and the movement and position on the grid are shown in the left panel. The main window shows the current image. In the right panel the histogram and Fourier spectrum of the current view are shown.

Type of presentation: Poster

IT-2-P-5735 Quantitative measurement of electron magnetic circular dichroism on polycrystalline iron film

Muto S.1, Rusz J.2, Tatsumi K.1, Adam R.3, Arai S.1, Kocevski V.2, Oppeneer P. M.2, Bürgler D. E.3, Schneider C. M.3
1Nagoya University, Nagoya, Japan, 2Uppsala University, Uppsala, Sweden, 3Peter Grünberg Institute, Jülich, Germany
jan.rusz@physics.uu.se

Electron magnetic circular dichroism (EMCD) is a measurement technique, which allows to measure element-specific spin and orbital magnetic moments using a transmission electron microscope. Since its discovery in 2006 [1] it went through a rapid development, bringing improvements in spatial resolution and signal strength. Yet, quantitative measurements remained difficult using the classical approach because of the low signal to noise ratio (SNR). The classical approach suggested to orient the sample into a 2-beam or 3-beam orientation and acquire core-level spectra (ELNES) at two distinct positions in between Bragg spots (so called Thales circle positions). Here we describe a different approach that overcomes these limitations and can be applied on a polycrystalline sample without setting any particular orientation [2]. We use a polycrystalline iron sample, Fig. 1, and acquire few hundreds of Fe L2,3 ELNES spectra, while scanning with the beam over the sample using a megaelectronvolt Jeol JEM-1000 K RS microscope. The scanning step is set to a value close to an average size of a crystalline grain in the sample, therefore it is very likely that every spectrum is taken from a different grain and thus at a different crystal orientation. The detector orientation is shifted by approximately G(110)/2 from the transmitted beam direction. In the next step the dataset is statistically processed and an averaged EMCD signal is accumulated, Fig. 2. The reliability of the method was checked on a NiO control sample, which shows no net magnetization. Finally, EMCD sum rules [3] have been applied to extract the ratio of the orbital to spin moment, Fig. 3. The obtained value 0.0429 +/- 0.0075 is in close agreement with x-ray magnetic circular dichroism measurements [4].

[1] P. Schattschneider, S. Rubino, C. Hebert, J. Rusz, J. Kunes, P. Novak, E. Carlino, M. Fabrizioli, G. Panaccione and G. Rossi, Nature 441, 486 (2006).

[2] S. Muto, J. Rusz, K. Tatsumi, R. Adam, S. Arai, V. Kocevski, P. M. Oppeneer, D. E. Burgler, and C. M. Schneider, Nature Comm. 5, 3138 (2014).

[3] J. Rusz, O. Eriksson, P. Novak, P. M. Oppeneer, Phys. Rev. B 76, 060408(R) (2007).

[4] C. T. Chen et al., Phys. Rev. Lett. 75, 152 (1995).


A portion of this work was supported by a Grant-in-Aid on Innovative Areas "Nano Informatics" (grant number 25106004) and on Young scientist A (24686070) from the Japan Society of the Promotion of Science. J.R. and P.M.O. acknowledge the support from the Swedish Research Council, J.R. acknowledges support from STINT and P.M.O. from the European Commission (grant No. 281043).

Fig. 1: TEM image of the polycrystalline iron sample. Scale bar corresponds to 50nm.

Fig. 2: Accumulated ELNES (blue and red curves) and EMCD (black line) spectra from the entire dataset of 225 individual ELNES spectra.

Fig. 3: Extrapolation of the ml/ms ratio as a function of the low-pass filter width. Low-pass filter was applied to individual ELNES spectra prior to statistical extraction of the EMCD spectrum. WIthout low-pass filter the statistical extraction was numerically unstable.

Type of presentation: Poster

IT-2-P-5974 High quality FIB lamella preparation for wide area atomic resolution TEM investigations

Straubinger R.1, Beyer A.1, Gries K. I.1, Schneider C.2, Rohnke M.2, Mogwitz B.2, Janek J.2, Volz K.1
1Material Sciences Center and Faculty of Physics, Philipps-Universität Marburg, Germany, 2Physikalisch-Chemisches Institut, Justus-Liebig-Universität Gießen, Germany
rainer.straubinger@physik.uni-marburg.de

A large increase in research efforts on thermoelectric power generation is currently occurring because of the improved properties of various nano structured thermoelectric materials. NaXCoO2 is a thermoelectric material which for example makes the recovery of the waste heat emitted by vehicles and factories possible. In addition it can be used in electronic processors. The single phase NaXCoO2 crystals we are working with are grown by pulsed laser deposition on Al2O3 (001) or LaAlO3 (001). To improve this material transmission electron microscopy (TEM) investigations are indispensable. Especially for structures that reveal a lot of inhomogeneity it is necessary to have high quality focussed ion beam (FIB) TEM lamellas for wide area atomic resolution.
During our research we continuously improve the FIB preparation process. Because of the hardness of the sapphire substrate it is necessary to thin the lamella from the substrate side (shadow FIB). Especially when working with very ion beam sensitive structures this preparation technique is also very interesting for other materials, like even organic material.
In Fig. 1 a high angle annular dark field (HAADF) scanning transmission electron microscopy (STEM) image of a cross sectional TEM lamella is shown. This lamella consists of the Al2O3 substrate, the NaXCoO2 layer, a Pt protection layer to avoid oxidation and a W protection layer deposited in the FIB. The bright areas within the NaXCoO2 layer are CoO2 impurities.It gets obvious that the thickness of the lamella does not change significantly over the whole field of view with the width of approximately 4μm. Thus, a good overview of a big sample region showing different features can be created using FIB. In Fig 2 a high resolution HAADF STEM image of the same lamella is presented, showing the interface between the substrate and a CoO2 impurity. It can be seen, that FIB preparation is a useful method to obtain thin samples over a wide range. This enables in combination with (S)TEM the characterization of samples containing a lot of inhomogeneities. This presentation will summarize the necessary steps to optimize the FIB preparation to obtain optimal samples.


Fig. 1: HAADF STEM image of NaXCoO2 grown on Al2O3 with a Pt + W cap layer.

Fig. 2: High resolution HAADF STEM image of a CoO2 impurity on Al2O3 substrate.

Type of presentation: Poster

IT-2-P-6055 Visualizing and correcting dynamic specimen processes in TEM using a large-format Direct Detection Device

Bammes B. E.1, Spilman M.1, Chen D. H.1, Jin L.1, Bilhorn R. B.1
1Direct Electron, San Diego, CA
bbammes@directelectron.com

Multiple factors reduce the resolution and signal-to-noise ratio (SNR) of transmission electron microscopy (TEM) images, including the microscope instrumentation, dynamic specimen processes (e.g., drift, beam-induced motion, charging, radiation damage, etc.), and inefficient electron detectors.

With the goal of overcoming many of these obstacles, Direct Electron introduced the first large-format Direct Detection Device (DDD®) in 2008, as the culmination of academic and industrial partnerships. Development has culminated in the DE-20 (5k x 4k) in 2012 and the new DE-64 (8k x 8k) this year. These DDD cameras deliver dramatically improved performance compared to traditional electron detectors such as film or CCD cameras.

In addition to improved efficiency and resolution, the architecture of DDD cameras allows for continuous streaming of unbinned full-frame images at ~30 frames per second, with no dead time between consecutive frames. Many TEM methods require a static specimen image, such as low-dose electron cryo-microscopy of biological specimens. In these methods, dynamic specimen processes are detrimental, causing either non-isotropic resolution loss (i.e., specimen drift) or overall degradation of the SNR in each image (e.g., beam-induced motion, charging, radiation damage, etc.). We have developed methods and algorithms for exploiting the “movie mode” output from DDD cameras to correct for these dynamic processes and maximize the isotropic resolution and SNR of each image. Briefly, a “movie” is acquired of a specimen at 2-3× the normal total electron exposure. To correct specimen drift (which is consistent across the entire image), the frames from the movie are iteratively aligned, and to correct beam-induced specimen motion and charging (which are local effects that vary across the image), sub-regions for each frame are iteratively aligned. To correct radiation damage, low-pass filters are applied to each frame based on expected damage rate of the specimen.

We have demonstrated the benefits of this method by using images of frozen-hydrated Brome mosaic virus (BMV). Images generated based on our method show improved isotropic high-frequency SNR along with significantly improved low-frequency contrast compared to conventional imaging (Fig. 1). We processed a data set of ~32,000 particles using both the conventional method and our new “damage compensation” method to generate de novo three-dimensional reconstructions of BMV. Our new method improved the resolution significantly from 4.4 Å to 3.8 Å resolution, thus demonstrating the power of damage compensation with a direct detection camera for high-resolution structural studies.


We sincerely thank the National Center for Macromolecular Imaging (Baylor College of Medicine, Houston, TX) for images and reconstructions of BMV. We also acknowledge funding from the National Institutes of Health (Grant #8R44GM103417-03).

Fig. 1: BMV imaging on a 300 kV TEM at 1 μm underfocus. (A) A particle with 20 e-/Å2 exposure, and (B) with the new method with a total exposure of 36 e-/Å2 with correction of dynamic specimen processes. (C) The Fourier transform of the image with the traditional method, and (D) the new method. (E) Comparison of the spectral SNR of (C) and (D).

IT-3. Super-resolution light microscopy and nanoscopy imaging

Type of presentation: Invited

IT-3-IN-1768 Superresolution Light Microscopy of nuclear Genome Organization

Cremer C.1,2,3
1Institute of Molecular Biology (IMB), D-55128 Mainz/Germany, 2Institute for Pharmacy and Molecular Biotechnology (IPMB), University Heidelberg, D-69120 Heidelberg/Germany, 3Kirchhoff-Institute for Physics (KIP), University Heidelberg, D-69120 Heidelberg/Germany
c.cremer@imb-mainz.de

The spatial organization of the genome in the interphase nucleus has far reaching functional consequences for gene regulation. Recently, various methods of superresolution light microscopy have been developed which made possible to enhance the spatial analysis of nuclear structures far beyond the conventional limits of about 200 nm in the object plane and 600 nm along the optical axis. Here, we report on quantitative nuclear nanostructure analysis based on Structured Excitation Illumination/Structured Illumination Microscopy (SEI/SIM), and on Spectrally Assigned Localization Microscopy (SALM), respectively. Presently, these approaches realized with custom-built systems allow us to optically resolve nuclear structures down to the range of ca. 120 nm laterally/350 nm axially using structured illumination, and few tens of nanometer in 3D using a special variant of localization microscopy, Spectral Precision Distance/Position Determination Microscopy (SPDM). In addition, both SIM and SPDM techniques were combined in a  single microscope setup. Application examples will be presented on the use of such ‘nanoscopy‘ approaches to perform quantitative analyses of individual small chromatin domains labelled by Fluorescence-in situ Hybridization (FISH); fluorescence-labelled replication units; of repair foci induced by individual accelerated heavy ions; of Fluorescent-Protein (GFP/YFP/mRFP) tagged histones and chromatin remodeling proteins; or of immunolabelled transcription/splicing related nanostructures. In addition, we report on the direct high resolution SPDM of nuclear DNA distribution, localizing more than 1 million individual DNA sites in an optical section of various types of mammalian cell nuclei. Some perspectives of these novel, quantitative “superresolution” microscopy methods for deciphering the „4D Nucleome“ will be discussed.


The support of the Boehringer Ingelheim Foundation, and of Heidelberg University is gratefully acknowledged.

Type of presentation: Invited

IT-3-IN-2856 Smart NanoBioImaging: multimodal correlative nanoscopy.

Diaspro A.1, 2, 3
1Department of Nanophysics, Istituto Italiano di Tecnologia, Genoa, Italy, 2Department of Physics, University of Genoa, Genoa, Italy, 3Nikon Imaging Center, NIC@IIT, Genoa, Italy
alberto.diaspro@iit.it

Nanoscopy and super resolution localization microscopy changed the paradigm in optical microscopy and increased the portfolio of applications along with new developments. Among biological applications, the demand for imaging cell aggregates (i.e., tumor spheroids) or tissues/organs and small organisms (i.e. zebrafish) and for performing multimodal investigations is challenging. Light scattering, polarization properties and other than light-based mechanisms of contrast can represent an important issue for further advances. Within such a framework, mixed technologies for investigating biological systems (and not only) at the nanoscale will be outlined. Specifically, the possibility of utilizing a Mueller matrix approach for scattering and polarization dependent data - also exploiting optically active biological structures, with particular interest in chiral objects - could lead to improve informative content of the formed images. For example, fluorescence and SHG data can be enriched by Mueller matrix signature and polarization considerations. As it was early demonstrated the possibility of getting ultrastructural information about chromatin-DNA organization by means of circular intensity differential light scattering makes the Mueller matrix integrated approach an effective good candidate projected to label free high-resolution imaging. To this end , a Mueller Matrix polarimetry integrated architecture will be outlined, based on photoelastic modulation. A Classical electrodynamics model can be the starting point to decipher high resolution information due to light scattering.
Moreover, although optical methods are a comparatively safe way to probe a biological system without substantial perturbation, scanning/surface probe microscopy had a relevant impact on biological imaging after the advent of atomic force microscopy (AFM). Force mapping and curves can be analyzed in order to obtain, for example, local elasticity information (Young’s modulus evaluation pixel by pixel) or performing molecular nanomanipulation, with a high specificity that generally lacks in atomic force microscopy. A hybrid modality, coupling super resolution methods based on individual molecule localization (IML, PALM, STORM) and on optical nanoscopy (STED, RESOLFT) with AFM will be critically discussed.
Multimodal and multidimensional correlative super-microscopy launches a new trend in microscopy. The focus is on asserting that the key elemental differences in the superresolution hyrbid approaches can be perceived as a modern overture for addressing old and new biological biological questions.


The auhor is indebted with members of the NanoBioImaging and NanoBIoPhotonics - LAMBS IIT research team. This work partially funded by  the Italian Programmi di Ricerca di Rilevante Interesse Nazionale PRIN 2010JFYFY2-002 grant.

Type of presentation: Oral

IT-3-O-1457 Imaging of cleared biological samples with the Ultramicroscope

Dodt H. U.1, Becker K.1, Hahn C.1, Saghafi S.1
11Vienna University of Technology, Chair of Bioelectronics, 1040 Vienna, Austria
dodt@tuwien.ac.at

In the last years we have developed a special Ultramicroscope (light-sheet microscope) for visualizing neuronal networks in whole brains. In the Ultramicroscope whole cleared brains are illuminated with a sheet of light and the optical sections are used for 3D reconstructions. This approach allows one to employ also low power, wide field objectives for imaging of large samples.

By clearing neuronal tissue with organic solvents (BABB) after dehydration, we could visulalize GFP-labelled neuronal networks in the whole brain [1.[ Improving our clearing technology by using tetrahydrofuran for dehydration and dibenzylether (THF/DBE) for clearing we were able to image GFP-labelled axons even in heavily myelinated spinal cord [2,3]. Also nervous and muscle structures in drosophila melanogaster can be imaged [4]. Our and other clearing solutions have non standard refractive indices. Due to a heavy refractive index mismatch imaging in these solutions with e.g. air or water immersion objectives gives therefore suboptimal results. We thus developed special objective devices that allow refractive index matched imaging. We show that high resolution imaging through 10 mm clearing medium is possible (Fig.1).

Furthermore we substantially increased the axial resolution of our light-sheet microscope by developing completely new optics for light sheet generation. These optics create an extremely thin light sheet by the use of a Powell- and several aspheric lenses. As light sheet thickness determines the axial resolution it is of pivotal importance for the performance of the light-sheet microscope. Our light sheet is static and will thus in future allow combination with other microscopic techniques which need constant nonscanned illumination. Examples for the application of the ultramicroscope as imaging of mouse brain, spinal cord and whole drosophila are given.

References

1. H.U. Dodt, U. Leischner, A. Schierloh, N. Jährling, C.P. Mauch, K. Deininger, J.M. Deussing , M. Eder, W. Zieglgänsberger, and K. Becker, Nat. Meth., 2007, 4, 331-336.

2. A. Ertürk, C.P. Mauch, F. Hellal, F. Förstner, T. Keck, K. Becker, N. Jährling, H. Steffens, M. Richter, M. Hübener, E. Kramer, F. Kirchhoff, H.U. Dodt, and F. Bradke, Nat. Med., 2012, 18, 166-171.

3. A. Ertürk, K. Becker, N. Jährling, C.P. Mauch, C.D. Hojer, J.G. Egen, F. Hellal, F. Bradke, M. Sheng, and H.U. Dodt, Nat. Protoc., 2012, 7, 1993-95.

4. C. Schönbauer, J. Distler, N. Jährling, M. Radolf, H.U. Dodt, M. Frasch, F. Schnorrer, Nature, 2011, 479, 406-409.


FWF

Type of presentation: Oral

IT-3-O-1814 Large parallelization of STED nanoscopy using optical lattices

Yang B.1,2, Przybilla F.1,2, Mestre M.1,2, Trebbia J. B.1,2, Lounis B.1,2
1Univ Bordeaux, LP2N, F-33405 Talence, France , 2Institut d’Optique & CNRS, LP2N, F-33405 Talence, France
jean-baptiste.trebbia@institutoptique.fr

Recent developments in super-resolution microscopy techniques achieved nanometer scale resolution and showed great potential in live cell imaging. STED (Stimulated Emission Depletion) and more generally RESOLFT (REversible Saturable OpticaL Fluorescence Transitions) need parallelization in order to fully benefit from this spatial resolution for fast wide-field imaging. An approach for parallelization is based on structured illumination pattern. RESOLF parallelization has been proposed using 1D interference pattern, but the resolution improvement is only obtained along one direction. Recently methods for massive parallelization of RESOLFT with photo-switchable proteins and STED nanoscopy based on the use of 2D structured illumination have been reported. Larger field of view could be achieved for parallelized RESOLFT using photo-switchable fluorescent proteins, because it requires less intensity to switch. However, protein switching is a relatively slow on-off process (~10 ms), which sets a limit to the imaging acquisition rate. Moreover RESOLFT with photo-switchable fluorescent proteins is constrained in its versatility by the need for genetic modification and transfection.

We show how well-designed optical lattices, created by multi-beam interference can provide efficient depletion patterns with moderate laser power and can be used for large parallelization of STED, so far the most important and widely used RESOLFT technique. The stimulated emission depletion being an ultrafast on-off switching process (~1 ns), its imaging speed is therefore only limited by the number of detected photons and fast large field of view super-resolution imaging can be achieved. With optical lattices, acquisition of large field of view super-resolved images only requires scanning over a single unit cell of the optical lattice which can be as small as 290 nm x 290 nm. STED images of 2.9 µm x 2.9 µm with resolution down to 70 nm are obtained at a frame rate of 12.5 images/s (figure 1).

Photobleaching might be a constraint in STED nanoscopy when recording a large number of frames. We reduce the photobleaching (i.e. the probability of a molecule to get promoted to a highly excited and reactive levels), by structuring both excitation and depletion beams. In this case, we found a decay time two times longer than in the homogeneous excitation case. We demonstraste the high-speed capability of our STED microscope by imaging the movement of 20 nm fluorescence particles in a Carbopol gel [1] (figure 2). We clearly show that recording of fast relative movement of two particles separated by a distance well below the diffraction limit.

Reference :

[1] B. Yang et al.,“Large parallelization of STED nanoscopy using optical lattices”, Optics Express, In Press (2014).


We acknowledge financial support from the ANR, Région Aquitaine, the French Ministry of Education and Research, the ERC and FranceBioImaging (Grant N° ANR-10-INSB-04-01).

Fig. 1: Structured excitation (a) and depletion (b) patterns for OL-STED (Optical lattice STED). (c) Fluorescence signal decays in the case of structured excitation and homogeneous excitation. (d) diffraction limited image of 20 nm fluorescent beads, (e) Super-resolved OL-STED image. (f) Normalized fluorescence intensity profiles.

Fig. 2: Diffraction limited (left side) and OL-STED (right side) successive images of 20 nm fluorescent beads. Images are taken at the rate of 80 ms per image with PSTED = 280 mW, Pexc = 2mW. The squares indicate the region where one bead is moving around another one.

Type of presentation: Oral

IT-3-O-2173 Effects of optical aberrations in single molecule and super resolution imaging

Coles B. C.1, Schwartz N.2, Rolfe D. J.1, Guastamacchia M.1,3, Lo Schiavo V.1, Martin-Fernandez M. L.1, Webb S. E.1
1Science & Technology Facilities Council, Rutherford Appleton Laboratory, Didcot, UK, 2Science and Technology Facilities Council, Astronomy Technology Centre, Edinburgh, UK, 3Heriot-Watt University, Edinburgh, UK
benjamin.coles@stfc.ac.uk

The images created in a fluorescence microscope are imperfect because of aberrations caused by the various objects in the optical path. These aberrations may be due to both the microscope optics and the sample itself. More specifically, they may be due to imperfect optic design and manufacturing processes, imprecise alignment of optical components and mismatched and inhomogeneous refractive indices within the sample specimen. Aberrations caused by the sample are necessarily greater as we go deeper, particularly in heterogeneous biological samples. As a consequence, most single-molecule tracking and single-molecule based super-resolution imaging is performed in microscopes equipped with total internal reflection illumination, which limits fluorescence excitation to a thin layer a few hundred nanometres thick. The aberrations impose limitations on the resolving power of a fluorescence microscope, the localisation precision of super resolution imaging and the proportion of features detected in single molecule tracking analysis. Achieving the same high resolution throughout a 3D sample such as biological cells depends on correcting these aberrations and recovering a high signal-to-noise ratio in deeper layers, which can be achieved using adaptive optics. Correcting the aberrations is particularly important for measuring accurate distances within a 3D volume. Note that it is not, however, trivial to determine what the aberrations from each point in the sample are, particularly in widefield microscopes.

There are approximately ten significant low-order Zernike modes of aberration present in a standard fluorescence microscope. We have investigated the effects of these different aberrations on data analysis in single molecule imaging techniques such as tracking in live cells. In particular, we have studied the consequences of aberrations on single molecule detection rates and apparent intensities in the poor signal-to-noise ratio environment of biological cells. These have wide consequences for the accurate determination of, for example, stoichiometry, FRET and diffusion rates from single-molecule measurements.

We have also studied the impact of aberrations on data analysis in single molecule localisation super-resolution techniques, such as STORM/PALM imaging, with varying levels of noise. Aberrations directly affect the signal-to-noise ratio and hence the achievable localisation precision.

Lastly, both Gaussian and astigmatic point spread functions were considered in order to extend the improvements to three dimensional super resolution imaging and single molecule tracking.


This work was funded by the UK's Medical Research Council under grant number MR/K015591/1

Type of presentation: Oral

IT-3-O-2789 Identify and Localise: Algorithms for Single Molecule Localisation Microscopy

Best G.1 2 5, Prakash K.3 4 5, Hagmann M.1 2, Cremer C.1 3 4, Birk U.1 4
1Kirchhoff-Institute for Physics (KIP), University of Heidelberg, Heidelberg, Germany, 2University Hospital Heidelberg, University of Heidelberg, Heidelberg, Germany, 3Institute for Pharmacy and Molecular Biotechnology (IPMB), University of Heidelberg, Heidelberg, Germany, 4Institute of Molecular Biology (IMB), Mainz, Germany, 5equal contribution
K.Prakash@imb-mainz.de

Single Molecule Localisation Microscopy (SMLM) is increasingly viewed as one of the major tool for analysis of biological processes on a high resolution level in the range of 10 to 50 nm. The procedure relies on sequential detection of (a subset of) individual fluorophores. For dense regions (fluorophores with significant overlap), a compromise between fluorescence labelling density and the photoswitching behaviour of fluorophores is needed to have an optical isolation i.e. sparse distribution of molecules in each acquired frame.

Algorithms used to precisely identify the locations of these fluorophores can be broadly classified into two categories, namely fitting based and non-fitting based (usually Centroid) methods. While iterative fitting-based methods can usually provide fitted parameters equal or close to the maximum likelihood estimate, ad hoc centroid based methods are usually very quick. However, all localisation methods struggle if the underlying model poorly represents the observed data e.g. background level, out of focus signals, noise, etc. A particular challenge for the exact fluorophore determination is posed by spatially as well as temporally fluctuating background intensities arising from out of focus blinking fluorophores. This is to some degree always given if the structure is not per se 2-dimensional (e.g. PALM using TIRF illumination).

Here, we present a comparative analysis of a range of available localisation algorithms on complexity, applicability and performance by testing them on both synthetic and experimental data that cover examples of both sparse and dense regions, with both low and high background levels to determine, which method is suited for a given set of data.


We gratefully acknowledge the colleagues at IMB who supported us with reagents. In particular, we would like to thank Aleksander Szczurek and Hyun-Keun Lee for samples, reagents and many interesting discussions. This work is supported by the Boehringer Ingelheim Foundation. The support of University Hospital Heidelberg (Prof. S. Dithmar) to G.B. and M.H. is also gratefully acknowledged.

Type of presentation: Poster

IT-3-P-1448 Nonlinear optical single-molecular image technique applying in nanostructure study of macromolecules from chicken egg

Wang X. M.1
1Hubei University of Chinese Medicine
foxglove@163.com

Nonlinear optical single-molecular image technique is a new technique which is not widely recognized in scientific community. It is our patent technique (Chinese patent 200910060951.7, PCT /CN2010/000138 ). It is a new innovative principle to get a profile image of tiny material noninvasively. by a series of lenses which diameter from small to large, adjusting each lens move back and forward carefully along a straight line ,the two different direction rays from the same point of out edge of objects can focus on a plate to form an image. This technique can magnify profile images of small samples. Its x-y resolution breakthrough the limit of Abbe’s diffraction law. Nowadays ,it resolution can reach 1-3nanometer . It can get single molecular image in water. It can be used to trace the trajectory of single molecule in living cell. The principle of technique will be broad application in many areas in near future. In this paper, we gave more supporting evidences that Nonlinear optical single-molecular image technique is a practical tool. The photos of our experimental results demonstrated that the principle of Nonlinear optical single-molecular image technique is correct. The photos showed the double helix structure of DNA extracted from chicken egg. This technique will bring human a lot of knowledge and information about molecules, especially about biological macromolecules in living cell. It will improve drug research and human understand micro world. With computer image reconstruct technique , Nonlinear optical single-molecular image technique will become more powerful tool for scientific research.


Gratful the Financial support of The nature science foundation of Hubei province 2006ABA060,The Education Department of Hubei province  D200516013

Fig. 1: macromolecules extracted from chicken egg white

Fig. 2: macromolecules deployed by double distilled water on glass slide 

Fig. 3: The photo showed double helix structure of macromolecules

Fig. 4: Diphenylamine assay demonstrated the macromolecules was DNA

Type of presentation: Poster

IT-3-P-2060 Sub resolution spectral discrimination of Lipufuscin-granules inside human RPE cells

Schock F.1,2, Best G.1,2, Celik N.2, Heintzmann R.4,5, Sel S.2, Birk U.3,6, Dithmar S.2, Cremer C.1,3,6
1Kirchhoff Institute for Physics, University of Heidelberg, Im Neuenheimer Feld 227, 69120 Heidelberg, Germany, 2Department of Ophthalmology, Universityhospital of Heidelberg, Im Neuenheimer Feld 400, 69120 Heidelberg, Germany, 3Institute of Molecular Biology, Ackermannweg 4, 55128 Mainz, Germany, 4Institute for Physical Chemistry, University of Jena, Lessingstr. 10, 07743 Jena, Germany, 5Institute of Photonic Technology, Albert-Einstein-Straße 9, 07745 Jena, Germany, 6Department of Physics, University of Mainz, Staudingerweg 7, 55128 Mainz, Germany
florian.schock@kip.uni-heidelberg.de

In the last decades, a variety of new microscopy methods has been developed to circumvent the classical resolution-limit[1]. Furthermore, combinations of confocal-microscopes and spectrometers have proven useful and are already commercially available. However, presently no existing device is used to combine spectrometry with superresolution.

Here we present a first attempt to analyze the emission spectrum of super resolution images in a clinically important field of application. We used a custom-made Structured Illumination Microscope (SIM) equipped for multicolor imaging[2]. At an excitation wavelength of 488 nm, this instrument provides an optical resolution down to about 120 nm in the object plane and 350 nm along the optical axis. To obtain spectral information we used different emission filters and calculated the resulting spectral bands.

We used this method on human retinal pigment epithelial (RPE) tissue sections and present first superresolution images on spectral information. We also present that this method is able to separate the intracellular autofluorescent Lipufuscin-granule-types that are connected to age related maculadegeneration. All work on human tissue was done according to the Declaration of Helsinki.

[1] Christoph Cremer and Barry R. Masters; Resolution enhancement techniques in microscopy; The European Physical Journal H; 2013
[2] Sabrina Rossberger, Thomas Ach, Gerrit Best, Christoph Cremer, Rainer Heintzmann, Stefan Dithmar; High-resolution imaging of autofluorescent particles within drusen using structured illumination microscopy; Br J Opthalmol 2013, 97, 518-523


Type of presentation: Poster

IT-3-P-2190 Super-resolution imaging of 3D-cultured cells by SAX microscopy

Yonemaru Y.1, Yamanaka M.1, Uegaki K.1, Smith N. I.1, Kawata S.1, Fujita K.1
1Osaka University, Osaka, Japan
fujita@ap.eng.osaka-u.ac.jp

We have proposed the use of saturated excitation (SAX) of fluorescence to improve the spatial resolution of confocal fluorescence microscopes [1]. SAX introduces the highly nonlinear relation between excitation and emission intensities. With using a focusing laser for fluorescence excitation, the nonlinear response is localized within the laser focus, therefore the extraction of fluorescence signal, which nonlinearly responds to the excitation intensity, can realize the spatial resolution beyond the diffraction limit [2]. Imaging of biological samples have been demonstrated firstly with a fixed sample [3] and recently applied to live cell observation with fluorescence proteins [4].

Since the spatial resolution in SAX microscopy is improved by the nonlinear relationship between the excitation and the emission, SAX microscopy has the imaging property similar to two-photon microscopy. The nonlinear relation between emission and excitation allows us to remove the background fluorescence signal generated at out-of-focus planes. With using this benefit, we have applied SAX microscopy to observed 3D-cultured HeLa cells [5, 6]. A cell cultured on a flat substrate expands its body and forms a thin and flat shape with approximately 10 µm thickness. On the other hand, a cell cultured in a 3D matrix (such as gel) has more freedom to expand in the 3D space and shape their bodies with a thickness of several tens of micrometers. Since many differences has already observed between cell functions in 2D and 3D cell culture and the 3D cell culture can provide a condition for cell growth closer to the nature, super-resolution imaging of 3D-cultured cell may become important in biological and medical researches in the near future.

Fig. 1a shows a SAX fluorescence image of actin in fixed HeLa cells cultured in 3D. The cells were cultured in gel and stained with ATTO Rho6G phalloidin. We scanned the entire cell cluster to obtain a 3D data set of fluorescence distribution in the sample to construct Fig. 1a as a projection of the data set. Fig. 1b shows the enlarged view of the dotted rectangle area in Fig. 1a. Fig. 1c shows the same area as Fig. 1b, but obtained by a typical confocal microscope without SAX. The comparison of Fig. 1b and 1c confirms the improvement of the spatial resolution and the image contrast by SAX.

Reference

[1] K. Fujita et al., Phys. Rev. Lett., 99, 228105 (2007).
[2] M. Yamanaka et al., Biomed. Opt. Express, 2, 1946 (2011).
[3] M. Yamanaka et al., J. Biomed. Opt., 13, 050507 (2008).
[4] M. Yamanakaet al. Interface FOCUS, 3, 2013007 (2013).
[5] M. Yamanaka  et al., J. Biomed. Opt., 18, 126002 (2013).
[6] Y. Yonemaru et al., ChemPhysChem, in press (2014).


This study was supported by Next-Generation World-Leading Researchers (NEXT program) of the Japan Society for the Promotion of Science (JSPS).

Fig. 1: a) SAX images of HeLa cells cultured in 3D matrix. Actin filaments were stained. b) the enlarged view of the area in the doted rectangle in a). c) the same area observed by a typical confocal microscope.

Type of presentation: Poster

IT-3-P-2200 Structured Illumination Microscopy With Multifrequency Patterns and LED Light Sources

Švindrych Z.1, Křížek P.1, Ovesný M.1, Borkovec J.1, Smirnov E.1, Raska I.1, Hagen G. M.1
1First Faculty of Medicine, Charles University in Prague
deden@seznam.cz

Structured illumination microscopy (SIM) has grown into a family of methods which achieve optical sectioning, resolution beyond the diffraction limit, or a combination of both these effects in optical fluorescence microscopy. SIM techniques rely on illumination of a sample with patterns of light which must be shifted between each acquired image. The patterns are typically created with physical gratings or masks, and the final optically sectioned or high resolution image is obtained computationally after data acquisition. Here we used a high speed ferroelectric liquid crystal microdisplay together with incoherent LED illumination to generate the illumination patterns and a sCMOS camera for widefield image acquisition. The high precision and flexibility of the generated patterns allowed us to use advanced processing techniques relying on the precise knowledge of the display-camera mapping, such as scaled subtraction in the case of optical sectioning SIM [1] and precise determination of spectral parameters (modulation period, direction and phase) in the case of super-resolution SIM. The freedom in choosing the illumination patterns also allows to tune the spatial frequencies and orientations of the patterns. Here we demonstrate the use of multi-frequency one-dimensional patterns to achieve both increased lateral resolution and high contrast optical sectioning with incoherent illumination and two-dimensional data processing in the Fourier domain (see inset in Fig. 1 C). We have also evaluated the impact of incoherent illumination on the SNR (signal to noise ratio) of the recovered high-frequency image components [2].

[1] P. Křížek, I. Raška, and G. M. Hagen, Opt. Express 20 (2012), p. 24585.
[2] M. G. Somekh, K. Hsu, and M. C. Pitter, J. Opt. Soc. Am. A. Opt. Image Sci. Vis. 26 (2009), p. 1630.


This work was supported by the Grant Agency of the Czech Republic [P304/09/1047, P205/12/P392, P302/12/G157 and 14-15272P], by Charles University in Prague [Prvouk/1LF/1, UNCE 204022], and by European Union Funds for Regional Development [OPPK CZ.2.16/3.1.00/24010].

Fig. 1: Fixed HT-1080 cells labeled with EdU-Alexa647 (fluorescently tagged nucleotide that incorporates into newly replicated DNA, red) and fluorouridine (synthetic nucleotide which is incorporated into active transcription sites, green), maximum projections of 20 sections, 0.1 µm step. A – widefield, B – optical sectioning SIM, C – super-resolution SIM.

Type of presentation: Poster

IT-3-P-2211 Flexible Structured Illumination Microscope with a Programmable Illumination Array

Švindrych Z.1, Křížek P.1, Ovesný M.1, Borkovec J.1, Raška I.1, Hagen G. M.1
1First Faculty of Medicine, Charles University in Prague
deden@seznam.cz

Structured illumination microscopy (SIM) has grown into a family of methods which achieve optical sectioning (OS-SIM), resolution beyond the Abbe limit (SR-SIM), or a combination of both effects in optical microscopy. SIM techniques rely on illumination of a sample with patterns of light which must be shifted and/or rotated between each acquired image. The patterns are typically created with physical gratings or masks, or by laser interference, and the final optically sectioned or high resolution image is obtained computationally after data acquisition. We used a flexible, high speed ferroelectric liquid crystal display for definition of the illumination pattern coupled with widefield detection and subsequent image processing. Focusing on optical sectioning, we developed a unique and highly accurate calibration approach which allowed us to determine a mathematical model describing the mapping of the illumination pattern from the microdisplay pixels to the camera sensor pixels. This is important for higher performance image processing methods such as scaled subtraction of the out of focus light, which require knowledge of the illumination pattern position in the acquired data. The calibration is also advantageous for SR-SIM reconstruction, as it provides precise information about reconstruction parameters (pattern period, orientation and phase) [1]. We evaluated the signal to noise ratio and the sectioning ability (see Fig. 1) of the OS-SIM reconstructed images for several data processing methods and illumination patterns with a wide range of spatial frequencies [2].

[1] M. G. L. Gustafsson, “Surpassing the lateral resolution limit by a factor of two using structured illumination microscopy,” J. Microsc., vol. 198, pp. 82–87, 2000.
[2] P. Křížek, I. Raška, and G. M. Hagen, “Flexible structured illumination microscope with a programmable illumination array.,” Opt. Express, vol. 20, no. 22, pp. 24585–99, Oct. 2012.


This work was supported by the Grant Agency of the Czech Republic [P304/09/1047, P205/12/P392, P302/12/G157 and 14-15272P], by Charles University in Prague [Prvouk/1LF/1, UNCE 204022], and by European Union Funds for Regional Development [OPPK CZ.2.16/3.1.00/24010].

Fig. 1: Comparison of different optically sectioning microscopes on a pollen grain sample. The SIM system achieves an optical sectioning thickness of 300 nm, much better than is possible in CLSM (Confocal Laser Scanning Microscopy).

Type of presentation: Poster

IT-3-P-3025 Measurement and simulation of PSF

Nahlik T.1, Stys D.1
1University of South Bohemia in Ceske Budejovice, Faculty of Fisheries and Protection of Waters, Institute of Complex Systems, Nové Hrady, Czech Republic
nahlik@frov.jcu.cz

Goal of the microscopy is to observe nature and man-made products which are smaller than resolution of human eyes. Microscope should help us to see more details but it has its own limits. One of these limitation is PSF (Point Spread Function). Microscope is device with many lenses and each of the lens can bring some distortion to the final image. When the object is so small that it can be consider as a point the image of this point will be not only the point but set of points. This image is called PSF. It is important to know and describe behavior of the PSF because it helps us to understand how the microscope transfers and shows the image.
Many simulation and measurements were done on the fluorescent microscope, because it is much easier. The difference between bright field and fluorescent microscopy is in the light. In BF the light is scattered by the sample, in fluorescence the light is emitted out of the sample. We tried to use ENZ (Extended Nijboer-Zernike theory) simulation (See Fig. 1). This works for fluorescent microscopy but it fails in BF (See Fig. 2). We tried to measure the PSF in BF (See Fig. 3) under different condition; with different particles (15nm and 200nm) and under different light condition (Light intensity is expressed by current on the LED used for illumination). We came to conclusion that the shape and size of PSF depends not only on parameters described in ENZ like aberration, light wavelength, numerical aperture, size of the particle but also on light intensity.
The size and shape of PSF is connected with the real resolution of the microscope and terms like distinguishability and discriminability. We propose that measuring of PSF should be the basic experiment done with every microscope. Measured PSF should be used as a kernel for deconvolution function for improving microscopy images.


CENAKVA CZ.1.05/2.1.00/01.0024, and CENAKVA II (The results of the project LO1205 were obtained with a financial support from the MEYS under the NPU I program); GA JU 134/2013/Z;

Fig. 1: Different aberration in ENZ Simulation. Parameters of simulation were – wavelength = 200nm, NA (numerical aperture) was 0.5, diameter was 400nm; A – no aberration; B – Coma; C – Tilt; D – Astigmatism

Fig. 2: ferent Light condition – Upper row - 15nm gold particles, Bottom row – 200nm latex particle. Light is increasing from left to right (starts at 1000mA, step 500mA, end at 3000mA)

Fig. 3: Comparing of ENZ simulation and real PSF. Red line shows position of focus. Right column shows original images in different distance from focus and appropriate Z position in PSF.

Type of presentation: Poster

IT-3-P-3035 Drift Correction Strategies for Single Molecule Localisation Microscopy

Hagmann M.1 2 6, Prakash K.3 4 6, Best G.1 2, Kaufmann R.5, Birk U.1 4, Cremer C.1 3 4
1Kirchhoff­-Institute for Physics (KIP), University of Heidelberg, Heidelberg, Germany, 2University Hospital Heidelberg, University of Heidelberg, Heidelberg, Germany, 3Institute for Pharmacy and Molecular Biotechnology (IPMB), University of Heidelberg, Heidelberg, Germany, 4Institute of Molecular Biology (IMB), Mainz, Germany, 5Division of Structural Biology, Wellcome Trust Centre for Human Genetics, University of Oxford, Oxford, UK, 6equal contributions
K.Prakash@imb-mainz.de

The correct position determination of fluorescent molecules is crucial for the interpretation of localisation microscopy data, e.g. of the biological structure investigated. The relative position of fluorophores with respect to the detector is highly sensitive to environmental disturbances (e.g. acoustic vibrations) and to mechanical instabilities of the microscope hardware (e.g. thermal expansion or mechanical relaxation). These disturbances cause distortion in the recorded image which pose a new natural limit to the localisation accuracy, especially for new acquisition protocols that allow acquisition times in the order of hours.

Here, we present two drift correction strategies based solely on data already acquired without any fiducial markers. We found that in some flavours of SMLM, many of the biological samples exhibit enough permanent (photostable) structure to reveal information about the sample location, or if this is not the case, reconstructions of a subset of the complete localisation data stack can be used at several time points to gain information about the sample drift. In both cases, the drift is found by determining the peak of the computed 2D-autocorrelation function. A polynomial or a set of Fourier functions is fitted through the data, based on which the dislocation of every localised fluorophore in a given frame of the acquired image stack is subtracted.

Using this approach, we successfully corrected localisation microscopy data down to a final drift less than 5 nm, which is comparable with fiducial markers based strategies. We demonstrate that with this procedure the resolution of the final reconstructions is substantially enhanced and the theoretical limit of localisation accuracy is almost restored.


We gratefully acknowledge the colleagues at IMB who supported us with reagents. In particular, we would like to thank Aleksander Szczurek and Hyun-Keun Lee for samples, reagents and many interesting discussions. This work is supported by the Boehringer Ingelheim Foundation. The support of University Hospital Heidelberg (Prof. S. Dithmar) to G.B. and M.H. is also gratefully acknowledged.

Type of presentation: Poster

IT-3-P-5736 Three-Dimensional Surface Modelling Of Cellular Structures And Intracellular Proteins Expression In Urinary Bladder Cancer Cells.

ELKABLAWY M. A.2
1Pathology Department, Faculty of Medicine, Menoufyia University, Egypt; , 2Pathology Department, Faculty of Medicine, Taibah University, Almadinah Almonourah, Saudi Arabia.
elkablawy@hotmail.com

Introduction: Visualizing and describing structure and function at a subcellular level is of fundamental importance in the molecular pathological sciences. The majority of microscopic techniques give a single two-dimensional representation of often complex three-dimensional (3D) structures. In many cases it would be useful to be able to view objects seen under the microscope in three dimensions, an approach which has clear advantages for the user in terms of comprehension and interpretation.
The aim of this study: was to produce 3D models of data from confocal laser scanning microscopy (CLSM) with a user-friendly PC software package. The study would examine the spatial localization of cellular proteins p53 and Bcl2 in apoptotic human urinary bladder cancer cell lines.

Material and method: Various combinations of immunofluorescence for the proteins of interest and propidium iodide staining for nuclear delineation were carried out in urinary bladder l cancer cells using contrasting fluorochromes. Specimens were examined with CLSM and stacks of optical sections from each light channel saved to disk. Using a desktop computer, the volumes of data were merged and rendered using gradient shading ray tracing. Surfaces of nuclei and protein aggregations were made transparent and the subcellular structures could be viewed from any angle. Series of renderings could be tagged together to produce a 3-D video moving sequences (build-up, fly-by and clipping) for better studying the protein expression.

In conclusion: the rendered images and videos were of high quality and were extremely helpful in studying the intracellular proteins expression. The technique both aided the research process and improved the means of scientific interpretation of gene and protein localization at the subcellular level.


the Author Thank Sarah Elkablawy, Ayah Elkablawy and Mona Mawlana for technical help and support. Also Thank dr. Nashaat Shawky for revising the language  and grammer of the abstract.

Fig. 1: this Images show the conversion of 2D image into series of 3D images

Type of presentation: Poster

IT-3-P-5751 Light induced in situ observation technology in EM

Huang M. R.1, Lin C. L.1, Chen K. F.1, Liu S. Y.1, Haung T. W.1, Tseng F. G.1, Chen F. R.1
1Engineering and System Science Department/National Tsing Hua University, Hsinchu, Taiwan
sgps51238@yahoo.com.tw

Compare with optical microscope (OM), electron microscope (EM) has the advantage with higher resolution. Therefore, when Hans Busch developed the first electromagnetic lens in 1926 [1], the technique of electron microscope has attracted increasing attention. The first commercial transmission electron microscope (TEM) were produced in 1939 by Siemens. [2] In today’s development, EM is important for the use of observing smaller sample which like the microstructure of the materials, the device in MEMS manufacturing and microbiological analysis. On the other hand, since the rise of environmental awareness, green energy become more and more concerned recently and solar energy is one of the candidates. As a result, a growing number of researchers are studying on using catalysts to transfer light into chemical energy. However, the efficiency of commercial solar energy devices still lower than 30 %, we thick it may be limited by the comprehension and observation of reaction mechanisms. Due to the importance of solar energy, we designed an experiment combined the observation of light-induced reaction and electron microscope.
In our experiment, two kinds of light-induced system using in electron microscope are produced. They are named “LED-based system” and “fiber-based system”. Figure 1(a) and (b) showed the”LED-based system” and “fiber-based system” respectively. The ”LED-based system” contains a light-emission diode (LED) put under the sample holder and controlled by a power supply. This kind of light-induced system is used in Hitachi TM-1000. Another “fiber-based system” includes a UV light source with optical fiber to induce UV light through the EDS hole for JEM-2010. Then we use the wet-cell sealing technology [3] to fabricate the wet sample containing TiO2, H2O, CH3OH and H2PtCl6 solution. Finally, we successfully observe the in-situ light-induced deposition reaction in EM. Figure 2 showed the deposition process in scanning electron microscopy (SEM). When light Illuminate about 3 minutes, the Pt particles began to emerge in the upper right corner.

Reference:
[1] Mathys, Daniel, Zentrum für Mikroskopie, University of Basel: Die Entwicklung der Elektronenmikroskopie vom Bild über die Analyse zum Nanolabor, p. 8
[2] "James Hillier". Inventor of the Week: Archive. 2003-05-01. Retrieved 2010-01-31.
[3] Tsu-Wei Huang, Shih-Yi Liu, Yun-Ju Chuang, Hsin-Yi Hsieh, Chun-Ying Tsai, Yun-Tzu Huang, Utkur Mirsaidov, Paul Matsudaira, Fan-Gang Tseng, Chia-Shen Chang and Fu-Rong Chen, Lab Chip, 2012,12, 340-347.


Fig. 1: Figure 1 Two light-induced system development which named (a) the”LED-based system”, and (b) the“fiber-based system”.

Fig. 2: Figure 2 The deposition process in scanning electron microscopy (SEM).

Type of presentation: Poster

IT-3-P-5777 ThunderSTORM: a comprehensive ImageJ plugin for PALM and STORM data analysis and super-resolution imaging

Křížek P.1, Ovesný M.1, Borkovec J.1, Švindrych Z.1, Hagen G. M.1
1Institute of Cellular Biology and Pathology, First Faculty of Medicine, Charles University in Prague
pavel.krizek@lf1.cuni.cz

We present ThunderSTORM, an open-source, interactive, and modular software designed for automated processing, analysis, and visualization of data acquired by single molecule localization microscopy (SMLM) methods including STORM, dSTORM, SPDM, PALM, and FPALM. ThunderSTORM was developed using a home-built SMLM system, but the software has been tested, and works well with data acquired using commercially available Nikon N-STORM and Zeiss Elyra systems. Our philosophy in developing ThunderSTORM has been to offer an extensive collection of processing and post-processing methods which were developed based on extensive testing with both real and simulated data. We also provide a very detailed description of the implemented methods and algorithms as well as a detailed user’s guide. ThunderSTORM is written in Java and distributed as a plug-in for ImageJ. This enables users to run the software on computers with different operating systems, and to use all of the advantages of ImageJ including its rich collection of plug-ins. The latest version of ThunderSTORM and the source code are freely available at https://code.google.com/p/thunder-storm/.


ThunderSTORM can process data for both 2D and 3D SMLM imaging, including data with high spatial molecular density, which is known as the “crowded field" problem. The steps involved in SMLM data processing are shown in Figure 1. Several algorithms for each of the processing steps have been implemented so experienced users have many options to adapt the processing to their data. However, the default settings perform very well on many of the SMLM data sets we have experimented with.


ThunderSTORM is also capable of generating simulated SMLM data and of evaluation of the performance of localization algorithms based on the ground-truth positions of the molecules. This allows users to perform Monte Carlo simulations and to quantitatively evaluate the performance of the applied algorithms. Localization results, as well as the ground-truth positions of molecules, can be imported/exported to/from ThunderSTORM in a variety of data formats, allowing compatibility with other SMLM localization software.


We also present our preliminary data of replication and transcription processes in the cell nucleus of HeLa cells, see Figure 2. Our labeling strategy results in brightly stained cells revealing two distinct nuclear activities. Investigating how these processes are organized within the cell nucleus relates to the larger issue of understanding how DNA replication is regulated on a cellular level. The super-resolution data were processed by ThunderSTORM.


This work was supported by the Czech Science Foundation [P304/09/1047, P205/12/P392, P302/12/G157, 14-15272P]; by European Union Funds for Regional Development [OPPK CZ.2.16/3.1.00/24010]; and by Charles University in Prague [Prvouk/1LF/1, UNCE 204022].

Fig. 1: Data processing pipeline for single molecule super-resolution imaging.

Fig. 2: Replication and transcription, detail of nucleolus. HT-1080 cells labeled with EdU-Alexa647 (a fluorescently tagged nucleotide that incorporates into newly replicated DNA, red) and fluorouridine (a synthetic nucleotide which is incorporated into active transcription sites, green).

Type of presentation: Poster

IT-3-P-5793 Nonlinear   optical single-molecular image technique

wang X. M.1
1Hubei University of Chinese medicine
foxglove@163.com

Nonlinear optical single-molecular image technique is a new technique which is not widely recognized in scientific community. It is our patent technique( Chinese patent 200910060951.7, PCT /CN2010/000138 ). It is a new innovative principle to get a profile image of tiny material noninvasively. by a series of lenses which diameter from small to large, adjusting each lens move back and forward carefully along a straight line ,the two different direction rays from the same point of the out edge of objects can focus on a plate to form an image. This technique can magnify profile images of small samples. Its x-y resolution breakthrough the limit of Abbe’s diffraction law.It can get single molecular image in water. It can be used to trace the trajectory of single molecule in living cell. The principle of technique will be broad application in many areas in near future. In this paper, we gave more supporting evidences that Nonlinear optical single-molecular image technique is a practical tool. The photos of our experimental results demonstrated that the principle of Nonlinear optical single-molecular image technique is correct. The photos showed the double helix structure of DNA extracted from chicken egg. This technique will bring human a lot of knowledge and information about molecules, especially about biological macromolecules in living cell. It will improve drug research and human understand micro world. With computer image reconstruct technique , Nonlinear optical single-molecular image technique will become more powerful tool for scientific research.


Fig. 1: single suger molecular image in double-distilled water

Fig. 2: starch chain  molecular image

Fig. 3: DNA double helix image

Fig. 4: our equitment of nonlinear single molecular image

Type of presentation: Poster

IT-3-P-5941 Development of the pump-probe nanoscopy architecture.

Korobchevskaya K.1, Bianchini P.1, Scotto M.1, Sheppard C.1, Diaspro A.1
1Nanophysics, Istituto Italiano di Tecnologia, 16136 Genova, Italy
ksenia.korobchevskaya@iit.it

Modern super-resolution microscopy techniques can provide high quality images with sub-diffraction resolution. However, a significant limitation for most of the methods is the use of fluorescence as readout, which results in photo-bleaching, and reduction of penetration depth, and significantly complicates the deep tissue imaging process due to strong scattering inside the sample. These issues can be avoided by using infra-red (IR) vibrational spectroscopy, but in that case, the image resolution is very low, due to the diffraction limit. As a new solution, the infrared absorption microscopy method was recently proposed [1, 2].

Strictly speaking the IR absorption microscopy method combines two well known techniques – transient absorption spectroscopy and stimulated emission depletion microscopy (STED) [3, 4]. In particular, it is based on a saturation effect, where the first laser pulse creates contrast within the sample, and the second pulse, at different frequency, detects the change. This allows label free transient images to be obtained. By introducing an additional doughnut-shaped depletion pulse, the excited area can be transiently saturated in the periphery of the focal spot, allowing collection of a signal from the central sub-diffraction area. By choosing the pumping and probing wavelengths, we can image different non-fluorescent species. The possibility to use IR light for sub-diffraction imaging is especially vital for deep tissue imaging, because the IR spectrum lies in the transparency window of biological tissues.


To verify the concept, we designed and assembled prototype of such a system presented in this work. The microscope combines two home-built setups – a pump-probe spectroscope and STED nanoscope. This configuration allows achieving high temporal and spatial resolution in the very same instrument. The setup is based on a femtosecond laser coupled with an OPA, which can generate laser pulses in a broad spectral range from visible to near IR, according to experimental needs. Therefore, explicit spatial and dynamical information about the sample can be obtained. This is very useful for cell biophysics and nanochemistry applications. We also present the possibility of implementing a 3D super-resolution capability.


The research leading to these results has received funding from the European Community’s Seventh Framework Programme: LANIR (FP7/20012-2015) under grant agreement n⁰ 280804.

Type of presentation: Poster

IT-3-P-5986 Inherently coaligned dual color STED microscopy

Göttfert F.1, D'Este E.1, Lukinavičius G.2, Hell S. W.1
1Max-Planck-Institute for Biophysical Chemistry, Göttingen, 2Ecole Polytechnique Fédérale de Lausanne, Lausanne
fabian.goettfert@mpibpc.mpg.de

Stimulated Emission Depletion (STED) Microscopy provides a versatile tool for investigating cellular structures on the nanoscale: While preserving the advantages of fluorescence microscopy, such as easy sample preparation and live cell imaging, the resolution is not limited by diffraction.


Here, we demonstrate the capabilities of STED microscopy for protein mapping and colocalization analysis. Applying a two color, pulse interleaved excitation and detection scheme combined with a single STED beam, two spectrally distinct dyes can be imaged almost simultaneously at a resolution down to 20nm. As the fluorescent volume in the sample is defined by the STED laser, the described setup provides a colocalization accuracy well below its resolution, insensitive to moderate misalignment and drift.


We imaged various proteins of the nuclear pore complex (Figure 1). Following standard immunolabeling protocols the staining was done with primary antibodies targeting the protein and secondary antibodies, tagged with the dye, targeting the primary antibody. Although the spectra of the dyes used for two-color imaging are overlapping, the crosstalk between the channels is below 20% and no post processing is required. These properties open the possibility to map even closely neighboring proteins, as exemplified on gp210 and NUP133.


Until recently, the lack of a suitable dye in the red wavelength region confined live cell STED imaging to wavelengths below 600nm. Long wavelengths however have the tendency to be less phototoxic and reduce background by autofluorescence. Here, we demonstrate live cell imaging with a newly developed silicon-rhodamine dye. Using the STED system described above we achieve a resolution better than 40nm in living cells, imaging microtubules and actin (Figure 2).


Fig. 1: STED image of nuclear pores in Xenopus cells. The protein gp210 (red) is arranged symmetrically around the central pore channel. As the diameter of the ring is approximately 160nm, the structure would not be resolvable with conventional fluorescence microscopy. In green, various FG repeat nucleoporins were labeled with a pan-specific antibody.

Fig. 2: Live rat primary hippocampal neurons stained with the recently developed fluorogenic dye SiR-actin. Actin in axons can arrange in rings with a spacing of 180nm. The used silicon-rhodamine fluorophore is cell permeable and minimally toxic to the cell. Its high photostability makes it a versatile probe for superresolution microscopy.

IT-4. Scanning electron microscopy

Type of presentation: Invited

IT-4-IN-1844 Ultra-low-energy STEM in SEM

Frank L.1, Nebesářová J.2, Müllerová I.1
1Institute of Scientific Instruments, Brno, Czech Republic, 2Biology Centre, České Budějovice, Czech Republic
ludek@isibrno.cz

Examination of thin samples in TEM or STEM has been performed at hundreds of keV. This energy range offered high resolution but low contrasts which meant that tissue sections had to be contrasted with heavy metal salts. Recent TEM with aberration correctors preserve an acceptable resolution down to 20 keV and provide enhanced contrasts [1]. The LVTEM device is operated at 5 keV on samples thinner than 20 nm [2]. STEM attachments to SEMs have become widespread [3] profiting from an image contrast substantially increasing even for light elements at tens or units of keV. The methods for the preparation of ultrathin sections of various substances are capable of producing layers at and even below 10 nm [4,5] which enables one to further decrease the energy of the electrons provided the image resolution is maintained. When using the STEM technique virtually all transmitted electrons can be utilised for imaging, while in TEM we are restricted to using electrons capable of forming the final image at acceptable quality. This forces us to narrow the ranges of the angular and energy spreads of electrons that enter the image-forming lenses. Consequently, the STEM technique promises higher contrasts at comparable resolutions. Unlimited reduction of the energy of the illuminating electrons is possible by employing the cathode lens principle [6]. This consists of biasing the sample together with its holder (made flat on both sides) to a high negative potential that retards the incident electrons before they land on the sample surface and accelerates backscattered and transmitted electrons to their respective detectors above and below the sample (Fig. 1). Calculations have shown a final spot size only moderately extended even at units of eV so that quality-consistent micrographs can be recorded over the full energy scale [7].

Ultra-low-energy STEM at hundreds of eV can be successfully applied to the examination of ultrathin tissue sections free of any heavy metal salts (Fig. 2) [8] or to 2D crystals. Single atomic steps are revealed at high contrast on multilayer graphene samples and transmittance of electrons at tens or units of eV can serve as a tool for “counting” the graphene layers (Fig. 3).

[1] Kaiser, U. et al., Ultramicroscopy 111 (2011) 1239.

[2] Drummy, L.F., Yang, J., Martin, D.C., Ultramicroscopy 99 (2004) 247.

[3] Morandi, V., Merli, P.G., Journal of Applied Physics 101 (2007) 114917.

[4] Riedl, T. et al., Microscopy Research and Technique 75 (2012) 711.

[5] Nebesářová, J., Vancová, M., Proceedings of IMC16, Sapporo 2006, Vol. 1, 500.

[6] Müllerová, I., Frank, L., Advances in Imaging and Electron Physics 128 (2003) 309.

[7] Müllerová, I., Hovorka, M., Frank L., Ultramicroscopy 119 (2012) 78.

[8] Frank, L. et al., Ultramicroscopy, submitted.


Support by the Technology Agency of the Czech Republic under no. TE01020118 and the institutional support RVO:68081731 are acknowledged.

Fig. 1: Trajectories of signal electrons toward transmitted (TE) and backscattered (BSE) electron detectors and through-the-lens detector (TLD) with the specimen immersed in the field of the open objective lens (a), with the biased sample retarding the beam 11 times (b), and with a combination of both (c).

Fig. 2: Section of mouse heart muscle, free of osmium tetroxide post-fixation and any staining, estimated thickness 5 nm, micrograph taken at 500 eV (a), electron energy dependence of the average edge resolution (b), and electron energy dependence of the relative variance contrast (c).

Fig. 3: Commercial CVD multilayer graphene imaged at 220 eV (a), total transmittance of extremely slow electrons through varying number of graphene layers (b).

Type of presentation: Invited

IT-4-IN-5709 A review of SE and BSE imaging in SEM and Variable Pressure SEM

Griffin B. J.1, Joy D. C.2, Michael J. R.3, Gauvin R.4
1Centres for Forensic Science, and for Microscopy, Characterisation and Analysis, The University of Western Australia, Perth, Australia , 2Centre for Nanophase Materials Science, Oak Ridge National Laboratory, Oak Ridge, USA, 3Sandia National Laboratories, PO Box 5800, Albuquerque, USA, 4Department of Materials Engineering, McGill University, Montreal, Quebec, Canada
brendan.griffin@uwa.edu.au

The SEM has evolved to a complex platform with multiple in-column and chamber-mounted SE and BSE detectors. With monochromated electron sources and biased stages the SEM can routinely provide sub-nanometre (0.4-7 nm) SE images and high voltages (15-30 kV), and in some cases low voltages (1 kV), on suitable samples. Live FFT analysis also now available on some instruments and it is particularly useful for high resolution operation. Another interesting and important development is the ability to collect images from the different detectors simultaneously, allowing consideration of the full range of sample information [figure 1], a consequence of higher resolution digital displays. The information content of the SE image is also much better understood, the in-column SED avoiding the well-documented but often forgotten “swamping” of the SE1 by SE2, SE3 and BSE through filtering and physical placement. Figure 1 also provides one example where surface contamination is most evident in the image from the in-column SED relative those from the BSED images and chamber-mounted SED image, even at high accelerating voltage (20 kV) and with C-coating. Stage tilt, reflecting angular selection, remains useful to enhance surface effects (such as relief in polished samples [figure 1b]). Electrostatic imaging of the chamber is a useful tool to illustrate the contribution of BSE generating SE3 on the pole piece and chamber walls in the vicinity of the E-T SED [figure 2b]. The angular sensitivity of BSE imaging has also been explored recently, particularly through the research of the late Heiner Jaksch in exploring the low angle BSE signals with ‘unconventionally’ short working distances. Two technologies have emerged, the use of in-column BSE detectors and selectable, segmented annular BSED. For example, the latter approach using a ‘tiled’, annularly-segmented single crystal Si diode detector [figure 2a] allows rapid switching through a range of collection angles, depending on the sample working distance. SE images in the variable pressure mode of SEM operation are strongly filtered by the positive ion cloud near and above the sample surface. The consequence is that such images lose the high resolution, near-surface component of the emitted SE. The collected signal is consequently dominated by higher energy SE and BSE and the deeper, delocalised ‘material’ contrast from the sample, even at the lower beam energies. It is currently difficult to envisage conditions in variable pressure mode where SE images will be directly comparable to those collected in high vacuum mode. The presence of a thin metal coating will assist. This filtering is not present where localized gas leakage is used for surface charge cancellation. 


FEI, TESCAN, Hitachi HT, and JEOL have all generously provided time on instruments.

Type of presentation: Oral

IT-4-O-1411 Influence of the work-function changes on the contrast of images in SEM.

Cazaux J.1, Sato K.2, Kuwano M.3, Ikatura N.4
1Physics Department, Faculty of Sciences, BP 1039, 51687 Reims Cedex 2, France, 2. JFE Steel Corp., 1 Kawasaki-cho, Chuo-ku Chiba Japan, 3MJIIT, Universiti Teknologi Malaysia, Jalan Semarak, 54100 Kuala Lumpur, Malaysi, 4Kyushu-University, Kasuga, Fukuoka 816-8580, Japan
jacques.cazaux@orange.fr

In the present contribution the role of the local change of the work function on some contrasts in SEM is suggested and illustrated. When electrons, BSEs or SEs, escape from the sample they are partly refracted at the sample/vacuum interface. The refraction effect is given by [1]: √ES sin β= √Ek sin α (1) with ES: inner kinetic energy (referred to the bottom of the conduction band); β: inner incident angle to the normal; Ek: kinetic energy into the vacuum (referred to the vacuum level); α: emission angle into the vacuum. The relationship between ES and Ek obeys to ES = Ek+EF+φ (2a) for metals of  work function φ and  Fermi energy EF and to ES = Ek+χ (2b) for semiconductors or insulators of affinity χ. In addition, the transmission probability of the escaping electrons, T(α), differs from 100%. Then a local change of φ or χ with the crystalline orientation or oxidation or contamination will change the SE or BSE yields δ or η.

Fig. 1a shows the calculated change of T(α) when χ changes from 4.05 eV (Si; Ge; SiC) to 4.55 eV for SEs of an initial inner kinetic energy of ES= 5.05 eV. Fig. 1b shows the corresponding distortion of the spectral distribution of the emitted SEs, ∂δ/∂Ek, when χ changes by steps of 0.2 eV and the corresponding change in δ, from 100% to 69%, is indicated in caption. Performed for the calculated dependence of few-layer graphene on SiC [1] this type of evaluation applies also  to contamination effects on synthetic diamond:Fig. 2a.

Fig. 2b shows the contrast of dendritic SiGe crystals embedded in a SiGe amorphous matrix[2]. Such a contrast may be explained from the local change of χ between the two crystalline forms of SiGe.

The same analysis applies to the angular selective detection of the BSEs where the refraction effects increase with the detection angle α –Eq. (1)-.This point should be considered for the interpretation of Fig. 3 for two different Fe grains, A and B [3].

The same analysis may be transposed to the reflectivity of Very Low Energy Electrons, R(α)=1-T(α) [4] , a reflectivity changing rapidly when the incident beam energy, E°, varies from 1 to 10 eV: a point fairly illustrated by Frank et al.[5].

In conclusion, some material and crystalline contrasts reported in SEM using SE or BSE detection as well as in LVSEM may be explained from the local change of the work function or the electron affinity.

1. J. Cazaux Appl. Phys. Lett. 98 (2011) p 013109 1

2. M. Itakura, N. Kuwano, K. Sato and S. Tachibana, Journal of Electron Microscopy 5 (2010) pS165

3. J. Cazaux , N. Kuwano and K. Sato, Ultramicroscopy, 135 (2013) 43

4. J. Cazaux J. Appl. Phys. 111 (2012) DOI: .1063/1.3691956

5. L. Frank, S. Mikmekova, M. Hovarka, Z Pokorna and I. Müllerova; Proceed. 15 th European Microscopy Congress, Manchester, (2012) PS2.2


Fig. 1:  a: Refraction effects of SE’s at vacuum/sample interface. b: Distortion of the spectral distribution of the emitted SE’s as a fonction of the change of χ from 4.05 to 4.65 eV [1].

Fig. 2:  a: Contamination contrast on diamond. b: SE Crystalline contrast of SiGe[2].

Fig. 3: BSE Crystalline contrast for E°=5 keV [3]. From top left to bottom right, the angle of detection with respect to the normal, α, changes: 58.5°±4.5; 46°±5; 35°±5; 14.5°±5; 0°.

Type of presentation: Oral

IT-4-O-1437 Advancements in Integrated Micro-XRF in the SEM

Witherspoon K. C.1, Cross B. J.2, Lamb R. D.1, Sjoman P. O.1, Hellested M. D.1
1IXRF Systems, Inc., 3019 Alvin De Vane Blvd, Suite 130, Austin, Texas, 78741, USA, 2CrossRoads Scientific, P.O. Box 1823, El Granada, CA 94018
mandih@ixrfsystems.com

In recent years, small x-ray tubes have been modified for mounting on Scanning Electron Microscopes. There have been two main types: low-power miniature tubes mounted re-entrantly within the SEM [1], and higher-power tubes with integrated x-ray optics to produce smaller beam spots at the sample, yet with intensities high enough for routine analytical work [1,2]. This addition allows samples to be analyzed both by X-Ray Fluorescence (XRF), and by the electron beam (SEM-EDS), as illustrated by the two spectra in FIG. 1.

Both techniques can be used independently or together by taking sequential e-beam and x-ray excited spectra. Quantitative analysis using this combined approach was first shown at the IMC16 conference in Sapporo [3]. This approach uses the advantage of e-beam excitation for lighter elements below 2.0 keV, and the more-efficient XRF excitation for x-ray lines above 2.0 keV. Micro-XRF with X-Y stage scanning can be used to collect x-ray elemental maps similar to those collected with e-beams, except the stage is moved versus scanning of the beam. This Micro-XRF mapping method has been proposed for some time [e.g.4], and was first commercially demonstrated in 1986 [5]. It is possible to collect e-beam and x-ray excited maps simultaneously for combined qualitative x-ray elemental mapping.

Currently 40µm and 10µm x-ray beam spot sizes are available inside the SEM. The 10µm beam has shown count rates that exceed 2000cps on steel. Future expectations are of even smaller excitation areas, with “useful” x-ray count rates. To create a smaller spot the polycapillary optic needs to be more tightly focussed. This means that the Focal (working) Distance (FD) of the XRF source must be shorter. For example, for a 40µm spot, an FD of 11 mm is typical. With a 10µm excitation spot, an FD of about 4.5 mm is required, making the integration of the x-ray beam a bit more of a challenge (FIG. 2).

It is now possible to use primary filters (thin foils) in front of the x-ray source. Using an automated filter wheel, allows in situ tuning of the x-ray source spectrum [e.g. 4], with improved elemental detection limits. An automated filter wheel between the x-ray source and sample provides comparable capabilities to those in benchtop XRF. FIG. 3 shows a comparison of unfiltered and filtered spectra, showing how the overall “shape” can be varied to optimize sensitivities and peak-to-background ratios.


Fig. 1: EDS (top) and Micro-XRF (bottom) spectra of NIST SRM 610

Fig. 2: Illustrates Focal Point and Working Distance

Fig. 3: Filtered Micro-XRF in SEM Spectrum

Type of presentation: Oral

IT-4-O-1674 Innovation possibilities of scintillation electron detector for SEM

Schauer P.1, Bok J.1
1Institute of Scientific Instruments of the AS CR, v.v.i., Brno, Czech Republic
petr@isibrno.cz

To evaluate performance of a scintillation detection system for SEM, it is necessary to consider many scintillator parameters. Various attributes of the scintillator for the SEM electron detector are listed in Fig. 1. The very important parameters are those affecting the detective quantum efficiency (DQE) which is primarily a measure of image noise. Not a less important indicator of image quality is the modulation transfer function (MTF) which describes the ability to show fine image details. Therefore, using a scanning imaging system, the detector bandwidth, which is given especially by the scintillator decay time, is the key to the good MTF. Currently, the YAG:Ce single crystal scintillator (introduced already in 1978 [1]) having somewhat limiting decay characteristic is the most frequently used scintillator in the SEM. The aim of this paper is to outline possibilities of scintillator innovation to get the improved MTF and DQE.

A database containing scintillation properties of various materials excited by hard x-rays and/or g-rays, taken from the literature, was established and is maintained at our laboratory. Among collected scintillators is only very limited selection of those that meet requirements for the SEM scintillation detector. For example, all hygroscopic materials must be excluded. Excluded must be also materials that have a low light yield and/or high luminescence decay. Thus the only suitable scintillators are those based on Ce-activated oxides characterized by a very fast 4f-5d emission as selected in Fig. 2.

Current research carried out in our laboratory tries to get faster scintillators by applying substitution of Y and/or Al in the garnet structure on the one hand and by increasing Ce-activator concentration on the other hand. Unfortunately, the Ce concentration increase is not an easy task for the Czochralski grown single crystals because of a sharp decrease of the distribution coefficient at crystal growth. But the development of optical ceramics is promising technology to get a more activated scintillator [2]. Our recent research includes the cathodoluminescence (CL) study of the commercial single crystal scintillators such as CRYTUR CRY18 and CRY19 as well as promising multicomponent garnet films grown by liquid phase epitaxy, for example GdGaLuYAG:Ce (formula (Gd,Lu,Y)3(Al,Ga)5O12:Ce3+). The new studied scintillators are quite fast as shown in Fig. 3. Their CL emission spectra show acceptable PMT matching as seen in Fig. 4.

References

[1] Autrata R., Schauer P., Kvapil Jos., Kvapil Ji.: A single crystal of YAG:Ce - new fast scintillator in SEM., J. Phys E: Sci. Instrum., 11 (1978), 707.
[2] Miyata T., Iwata T., Nakayama S. and Araki T., Meas. Sci. Technol. 23 (2012), Article No: 035501, DOI: 10.1088/0957-0233/23/3/035501.


The authors thank CRYTUR comp. for the supply with single crystal scintillators. They also thank Charles University, Faculty of Math. & Phys., for the supply with film scintillators. The work was supported by the Technology Agency of the Czech Republic (TE01020118). It was also supported by the European Commission and Ministry of Education, Youth, and Sports of the Czech Republic (EE.2.3.20.0103).

Fig. 1: Influence of various scintillator attributes on the choice of the best scintillator for the SEM electron detector.

Fig. 2: Compilation of x-ray and/or g-ray excited rare-earth activated oxides having the light yield ≥ 10 photons/keV and the decay time (τ1/e) ≤ 100 ns.

Fig. 3: CL intensity decay characteristics of the new scintillators: CRY18 single crystal and GdGaLuYAG:Ce garnet film. For comparison the decay of the improved YAP:Ce single crystal scintillator is also shown.

Fig. 4: Normalized CL intensity spectra of the new scintillators: CRY18 single crystal and GdGaLuYAG:Ce garnet film. For comparison the spectrum of the improved YAP:Ce single crystal scintillator is also shown.

Type of presentation: Oral

IT-4-O-1791 Quantitative interpretation for angle selective backscattering image of iron oxide on steel

Aoyama T.1, Nagoshi M.2, Sato K.2
1JFE Steel Corporation, Fukuyama, Japan, 2JFE Steel Corporation, Chiba, Japan
to-aoyama@jfe-steel.co.jp

The contrasts in backscattered electron (BSE) images were studied from the cross section of a heat-treated steel sheet using a scanning electron microscope (SEM) equipped with a conventional annular BSE detector (Σigma, Carl Zeiss NTS GmbH). The specimen used was heat-treated low carbon steel with an oxide layer mainly composed of magnetite (Fe3O4). A cross-sectional specimen was prepared by argon ion irradiation (IB-09010CP, JEOL Ltd.) after polishing with diamond suspension. BSE images were observed at primary electron energies (Eps) of 2 keV, 5 keV, 10 keV and 15 keV at various working distance from 2 to 15 mm for an identical area of the specimen (cross section). The take-off angles (θ; measured from the specimen surface) of the detector were estimated to be 35-45°, 39-53°, 50-63°, 66-75° and 73-79° (except 2 keV) from the geometry of the detector and the specimen. The variation of BSE intensities between crystal grains was calculated from the images. According to the results, high Ep enhances bulk information and Z contrast, whereas low Ep improves surface information and channeling contrast. High θ also enhances bulk information and Z contrast, whereas low θ improves surface information and channeling contrast. In the case of the lowest θ, topographic information was enhanced by shadowing effect on BSEs, in addition to the amplification of channeling contrast. These results regarding channeling contrast and Z contrast can be understood by the ratio of low-loss electrons (LLEs) to the inelastic BSE components detected; LLEs contribute to channeling contrast, and their ratio increases with decreasing Ep and θ. The systematic results obtained in this study are useful for controlling SEM conditions in order to select Z and crystallographic information separately in BSE images for practical materials of interest.


Authors appreciate Dr. Š. Mikmeková of JFE Steel Corporation for her detailed advice on quantitative analyses of the image contrasts. And we would also like to thank Mr. M. Yamashita and Ms. K. Takase of JFE Steel Corporation for their technical supports.

Fig. 1: Schematic diagram showing dependencies of the BSE contrast on the θ and Ep. The areas where channeling contrast and Z contrast are enhanced in the BSE images are indicated by shaded and unshaded areas, respectively. The area where topographic information and channeling contrast are enhanced is indicated by dotted area.

Type of presentation: Poster

IT-4-P-1456 Photonic Crystal Structure of Butterfly Wing Scales Exhibiting Selective Wavelength Iridescence

Matějková-Plšková J.1, Mika F.1, Jiwajinda S.2, Dechkrong P.2, Svidenská S.3, Shiojiri M.4
1Institute of Scientific Instruments of the ASCR, v.v.i., Královopolská 147, Brno 612 64, Czech Republic, 2Bioresources and Biodiversity Section, Central Laboratory and Greenhouse Complex, Kasetsart University, Kamphaengsaen Campus, Nakhonpathom 73140, Thailand, 3Institute of Cellular Biology and Pathology, First Faculty of Medicine, Charles University in Prague, Albertov 4, 12801 Praha 2, Czech Republic, 4Professor Emeritus of Kyoto Institute of Technology, 1-297 Wakiyama, Kyoto 618-0091, Japan.
filip.mika@isibrno.cz

Characteristic patterns and the vivid coloration of the wing scales of butterflies have lately attracted considerable attention as natural photonic crystals. The coloration of butterflies that exhibit human visible iridescence from violet to green has been elucidated. A Sasakia charonda (S. charonda) or ‘great purple emperor’ butterfly (Fig. 1a) was sampled in a woodland in Japan, and an Euploea mulciber (E. mulciber) or 'striped blue crow’ butterfly (Fig. 1b) was reared from an egg at the Environmental Entomology Research and Development Center, Kasetsart University. SEM observations, with the aid of the optical reflectance measurement, revealed that highly tilted multilayers of cuticle on the ridges in their iridescent scales (Figs.1e-1h and Fig. 2a-h) cause a dark zone where no reflection occurs (Fig. 2i)1-3 and produce a limited-view, selective wavelength iridescence (ultraviolet (UV)~green) as a result of multiple interference between the cuticle-air layers (Fig. 2j).3,4 TEM observation of S. charonda’s iridescent scales, sectioned with an ultramicrotome confirmed these results (Fig.2j and 2k). The iridescence from Chrysozephyrus ataxus (C. ataxus) or Thermozephyrus ataxus butterflies (Fig. 1c), which were sampled in Japan, originates from multilayers in the groove plates between the ridges and ribs (Fig. 3a-3f).3,5 The interference takes place between the top and bottom surfaces of each layer and incoherently between different layers. Consequently, the male with the layers that are ~270 nm thick reflects light of UV~560 nm (green) and the female with the layers that are ~191 nm thick reflects light of UV~400 nm (violet). A Troides aeacus (T. aeacus) or ‘golden birdwing’ butterfly (Fig. 1d) also grew in Kasetsart University, The butterfly does not produce any iridescent sheen which Troides magellanus does.3,4 No iridescent sheen is ascribed to microrib layers, which are perpendicular to the scale plane (Fig. 3g-3j), so that they cannot reflect any backscattering. The structures of these butterflies would provide us helpful hints to manipulate light in photoelectric devices, such as blue or UV LEDs.

1J. Matějková-Plšková et al., J. Micros. 236, (2009) 88.

2J. Matějková-Plšková et al., Mater. Trans. 51, (2010) 202.

3F. Mika et al., Materials 5, (2012) 754.

4P. Dechkrong et al., J. Struct. Biol. 176, (2011) 75.

5J. Matějková-Plšková et al., Mater. Trans. 52, (2011) 297.

Present address of J. Matějková-Plšková: Sadovského 14, Brno 612 00, Czech Republic.


Presenting author acknowledges the support from MEYS CR (LO1212) together with EC (ALISI No. CZ.1.05/2.1.00/01.0017).

Fig. 1: (a) S. charonda. (b) E. mulciber. (c) C. ataxus. (d) T. aeacus. (e-g) SEM images of iridescent white scales of the S. charonda. (h) Multiple cuticle-air gap arrangement on the rides of the scales.

Fig. 2: OM (a) and SEM images of iridescent white and blue scales of the E. mulciber (b-h). (i) Dark zone appearing on the S. charonda wing. (j) Iridescent reflection. TEM images of cross- sections of white (k) and blue scales of the S. charonda (l).

Fig. 3: (a-e) SEM images of iridescent scales of the male C. ataxus. (f) Rides and groves. (g-i) SEM images of yellow scales of the T. aeacus, which has not cuticle layers on the ridges but has microribs on the side of ridges. The microribs normal to the wing plane do not cause any backreflection as indicated by blue arrows.

Type of presentation: Poster

IT-4-P-1423 Development of high-efficiency DF-STEM detector

Kaneko T.1, Saitow A.1, Fujino T.1, Okunishi E.1, Sawada H.1
1JEOL Ltd. 1-2 Musashino 3-Chome, Akishima, Tokyo 196-8558, Japan
takekane@jeol.co.jp

Most recently, observations at low accelerating voltages have been increasingly popular for carbon-based materials such as carbon nanotubes or graphenes, to reduce knock-on damage due to irradiation of an electron beam. High-resolution dark-field (DF) imaging in a scanning transmission electron microscope (STEM) has been widely used for structural analysis in materials science. In conventional system, a STEM signal is detected as light intensity emitted from a scintillator which is hit by electrons. The conventional detector showed low signal conversion efficiency from an electron to a STEM image signal at low accelerating voltages, since the STEM detector is optimized for high energy electrons. Thus, a detector with good efficiency from low to high accelerating voltages is sought after. The STEM detector is consisted of a scintillator, a glass light guide and a photomultiplier tube. Many materials for scintillator were tested to improve the efficiency. The scintillator of the STEM detector is selectable either a powder scintillator or a single crystal scintillator. However, the good efficiencies from low to high voltages were not found yet so far. We measured the efficiency for powder and single crystal scintillators whose chemical compositions were the same, depending on the accelerating voltages. The measured results showed that the scintillation efficiency for the single crystal becomes higher than that of powder at accelerating voltage greater than 100 kV. Combining these features, we have developed a hybrid type scintillator, which consisted of powder deposited layer and a single crystal substrate. The luminescent quantum efficiency of the hybrid scintillator was measured to be twice as large as that of the single crystal at 60 kV and was about 8 times higher than that of the powder at 300 kV, and covers the observation at the accelerating voltages from low to high voltages. Especially, it is useful for low voltage observations of carbon-based materials consisted of few atomic layers that produces weak scattering of electron.


This work was supported by JST under the Research Accelerating Program (2012–2016).

Type of presentation: Poster

IT-4-P-1467 Energy Low Loss Backscattered Electrons Imaging in Material Characterization and Analysis

Liu X.1
1Carl Zeiss Microscopy GmbH
xiong.liu@zeiss.com

With the continuous size and structure shrinkage in semiconductor and electronic devices, the final performance and properties of the materials are dominated by the surface and interface layers. This requires scanning electron microscope (SEM) as a most conventional technical method in material characterization and analysis not only to be able to visualize and image such nanostructures with the secondary electron imaging under a low energy beam but also to analyze the tiny compositional differences like doping contrast, oxidation states of elements, small phases of hybrids or function group in polymers etc., which are not available via the classical backscattered electron imaging or other Energy-dispersive X-ray spectroscopy methods. Although the classical backscattered electron (BSE) imaging are from the multiple inelastic scattering process which could provide density related contrast like channeling contrast at high energy beam, the backscattering coefficient shows non-linear behavior and get very complicated.
In the classical backscattering process (Rutherford scattering), the backscattered electrons are mainly from the scattering of the high energy primary electrons with the nucleus charge or inner electron shells of the material. In such a case the contrast or brightness of the BSE imaging scale with material density, atomic number (Z). However the scattering between the primary beam with the outer electron shells of the materials at low impact energy (below 3 kV) region is not any more negligible which even becomes more dominant where the surface plasma resonance and ionization loss could happen and contribute to in the total contrast mechanism.
The unique design of the Gemini® lens integrated with a beam booster in the beam path not only maintains the brightness of the downward primary electron beam at low energies but also has a dispersion function for the generated reverse electron signals backward into the column. It means that the secondary electrons and backscattered electrons with a small energy and angle differences could be amplified and separated by the Gemini® lens in real time and space without converting the signals or by applying any additional stage bias. The separated backscattered electrons could be further filtered with an energy filtering grid and projected back into the corresponding detector. Backscattered electrons with a specific energy low loss could be picked out for imaging by setting an appropriate threshold potential to the filtering grid. After the grid filter the multiple inelastic scattered electrons could be cut away and the signal is consisted of the so-called energy low loss backscattered electrons which reveals some characteristic resonance of the materials.


Fig. 1: The SE1 image (left) and corresponding LL-BSE image (right) of the Ceincorporated into mesoporous silica as catalyst where the Ce ions andnanoclusters give high brightness.

Fig. 2: The SE1 image (left) and corresponding LL-BSE image (right) of the ZnSxO1-x thin film on Al2O3 substrate where the LL-BSE image is from the low loss BSEs with an energy between 700 eV and 800 eV.

Type of presentation: Poster

IT-4-P-1503 Using an EBSD Detector as a Microstructural Imaging Device

Wright S. I.1, Nowell M. M.1, de Kloe R.3, Camus P. P.2
1EDAX, Draper, Utah, United States, 2EDAX, Mahwah, New Jersey, United States, 3EDAX, Tilburg, The Netherlands
stuart.wright@ametek.com

Electron Backscatter Diffraction (EBSD) has proven to be a useful tool for characterizing the crystallographic orientation aspects of microstructures at length scales ranging from tens of nanometers to millimetres in the the scanning electron microscope (SEM). With the advent of high-speed digital cameras for EBSD use, it has become practical to the EBSD detector as an imaging device similar to a backscatter (or forward-scatter) detector [1-3]. When the EBSD detector is used in this manner, images exhibiting topographic, atomic density and orientation contrast can be obtained at rates similar to slow scanning in the conventional SEM manner. The same high-tilt (~70°) sample geometry is used and the camera is binned considerably – to a 5x5 “super-pixel” image - in order to get extremely fast acquisition rates. At such high binning, the captured patterns are not suitable for indexing. However, no indexing is required to for using the detector as an imaging device. Rather, a 5x5 array of images is formed by essentially using each super-pixel as an individual scattered electron detector as shown in Figure 1. The images formed in this way can then be combined in a variety of ways to form composite images of the microstructure as shown in Figure 2. The flexibility to combine these images together allows different contrast mechanisms to be emphasized in the composite images. While images formed in this manner lack the quantitative nature of the maps formed by using EBSD in the traditional manner, they still provide a wealth of information that can be obtained at rates much faster than the quantitative EBSD maps and with much less EBSD expertise required by the operator.

References

[1] S. I. Wright & M. M. Nowell (2006) “Microstructure Characterization Using EBSD Image Quality Mapping”, Presentation at THERMEC, Vancouver, Canada.
[2] R. Schwarzer, J. Sukkau & J. Hjelen (2011) “Imaging of topography and phase distributions with an EBSD detector in the SEM”, Poster presentation at Microscopy Conference, Kiel, Germany.
[3] E. J. Payton, L. Agudo Jácome & G. Nolze (2013) “Phase Identification by Image Processing of EBSD Patterns” Presentation at Microscopy & Microanalysis, Indianapolis, USA.


Scott Lindeman of EDAX is gratefully acknowledged for his assitance.

Fig. 1: A five by five array of images formed using a heavily binned EBSD detector as an array of twenty-five individual scattered electron detectors from a Mylonite sample.

Fig. 2: A composite image formed by combining the individual images shown in Figure 1together in the manner shown in the accompanying schematic. The circular outline in the schematic shows the position of the phosphor screen relative to the 5x5 pixel array.

Type of presentation: Poster

IT-4-P-1551 A simple way to obtain BSE image in STEM

Tsuruta H.1, Tanaka S.2, Tanji T.2, Morita C.3
1Department of Electronics, Nagoya University, 2EcoTopia Science Institute, Nagoya University, 3Meijo University
s-tanaka@esi.nagoya-u.ac.jp

Backscattered electron (BSE) signal has been used to image small objects in a liquid phase. A thin film such as silicon nitride (SiN) film was used to seal the liquid solution, and imaging electron beam was incident through the film, and BSE signals were detected for imaging. Such a technique is usually based on SEM, thus the accelerating voltage available is up to about 30 kV. In this paper, we introduce a simple BSE detector that is easily incorporated into a scanning transmission electron microscope (STEM) sample holder, and present some results for BSE imaging using STEM electron beam up to 200 kV.

Fig. 1 is a schematic representation of our BSE detector. The BSE detector is consisted of p-type silicon (Si) and Schottky contact. A dimple was made from one side, and a through-hole with a diameter of about 200μm was created at the bottom of the dimple. A thin Schottky electrode was made on this side. On the other side, an ohmic electrode was made. TEM grids were used to hold particle objects, and the grid was placed just below the detector. This was conveniently done with a silver paste. Observation experiments were performed using Hitachi H-8000 STEM (accelerating voltage 75 - 200 kV). The beam current was about 1.5 nA.

We used two types of samples. One was latex (Φ90 nm) and Au (Φ60 nm) particles on a carbon film coated grid. The other was Au (Φ60 nm) particles confined between two SiN membrane window grids (fig.2). The Au particles of this sample were in air atmosphere.

Figs. 3(a) and 3(b) are dark-field (DF) STEM and BSE images, respectively, of the latex and Au sample taken at 75 kV. Both latex and Au particles are visible in the BSE image, and they are distinguishable according to their intensity. Au particles appear brighter than latex particle. The detector current at bright Au particles was about 130 nA. On the other hand, for the STEM image, the difference of the intensities is not so noticeable, and it is difficult to distinguish latex and Au particles. Fig. 4 shows a BSE image of Au particles confined between the two SiN membrane window grids, taken at 200 kV. These particles were located on the upper membrane. In spite of the presence of 100 nm thick membrane, we can see each particle clearly, owing to the usage of a high accelerating voltage of 200 kV. The detector current at bright Au particles was about 8 nA. The low current is mostly due to the low backscattering probability as compared with 75 kV. And partially because of the fact that the BSEs at 200 kV are distributed higher angle than at 75 kV.

Our BSE detector was conveniently fixed to the sample grid with a silver paste. But it was able to remove the detector from the sample grid with tweezers. So, the detector was reusable until breakage which may happen by mistake.


Fig. 1: Schematic representation of the BSE detector.

Fig. 2: Au (Φ60 nm) particles were confined between two SiN membrane window grids. Au particles were in air atmosphere. This was fixed to the detector with a silver paste for BSE observation of Au particles.

Fig. 3: Latex and Au sample taken at 75 kV. (a) Dark-field (DF) STEM and (b) BSE images, respectively.

Fig. 4: BSE image of Au particles confined between the two SiN membrane window grids taken at 200 kV.

Type of presentation: Poster

IT-4-P-1565 Imaging of nanoparticles in cells with backscattered electrons in a scanning electron microscope

Müller E.1, Seiter J.1, Blank H.1, Gehrke H.2, Marko D.2, Gerthsen D.1
1Laboratory for Electron Microscopy, Karlsruhe Institute of Technology, Karlsruhe, Germany, 2Department of Food Chemistry and Toxicology, University of Vienna, Vienna, Austria
erich.mueller@kit.edu

Scanning electron microscopy (SEM) is an established technique for ultrastructure imaging of cells. Backscattered electrons (BSEs) yield subsurface information and atomic-number contrast [1] and are used in this work to image cellular structures and NPs incubated in cells. Specifically, optimum primary electron energies E0 for BSE imaging were determined for thin cell sections with thicknesses 100 nm ≤ t ≤ 1000 nm deposited on indium-tin-oxide (ITO-)covered glass slides which are interesting substrates for correlative light and electron microscopy imaging [2]. We also developed a technique to determine the information depth (ID) which denotes the maximum subsurface depth at which an object can be imaged.
Thin cell sections of HT29 colon carcinoma cells incubated with SiO2 nanoparticles (NPs) of 40 nm size were studied (see [3] for sample preparation). Poststaining was omitted to avoid artifacts. SEM was performed with an FEI Quanta 650 FEG with a low-voltage high-contrast detector (vCD).
Small E0 were selected to limit the escape depth of BSEs because electrons from large sample depths degrade image resolution and contrast. Fig. 1a shows a 2.5 keV BSE image of a 200 nm section. The SiO2 NPs, typically contained in vesicles, can be easily detected due to their bright contrast. Cell organelles display high contrast despite the lack of poststaining in Fig. 1b.
Fig. 2a shows a 1 µm section where E0 up to 7.5 keV can be applied without sample charging. In addition to the incubated SiO2 NPs, Au NPs with a size of 40 nm are present on the surface and can be distinguished due to their higher BSE intensity. BSE images were taken at different E0 between 1.5 and 7.5 keV for depth-dependent detection of SiO2 NPs. With increasing E0 more NPs become visible corresponding to the increasing ID. The depth of NPs from the surface was determined by tilting the sample and applying a triangulation method. In Fig. 2b the experimentally determined particle depths (dots) are plotted as a function of E0 and are compared with calculated ID values obtained by Monte-Carlo simulations (triangles). Based on the escape depth T = f·A·E01.67/(ρ·Z0.89) (Z: average atomic number, A: average atomic weight, ρ: density) proposed in [4], an analytical expression for the ID was obtained by fitting the experimental data with a modified factor f. This expression allows the determination of the ID of BSEs in biological samples. Experiments with entire cells grown on ITO-coated glass are promising with respect to NP detection and are subject of further work.

References
[1] H Niedrig, J. Appl. Phys. 53 (1982), p. 15.
[2] H Pluk et al., Journal of Microscopy 233 (2009), p. 353.
[3] J Seiter et al., J. Microscopy, accepted.
[4] K Kanaya and S Okayama, J. Phys. D: Appl. Phys. 5 (1972), p. 43.


Fig. 1: BSE SEM images of a 200 nm section of an HT29 cell deposited on an ITO-covered glass substrate. (a) Overview image taken at E0 = 2.5 keV. SiO2 NPs and organelles are visible in the cell. (b) High-magnification image taken at E0 = 3.5 keV. NPs and membranes can be well resolved.

Fig. 2: (a) 7 keV BSE SEM image of a 1000 nm section of an HT29 cell. SiO2 and Au NPs show different contrast compared to the cell matrix. (b) Plot of the information depth as a function of E0 with experimentally determined values (dots) and Monte Carlo simulations (triangles). Fit curves are based on the modified Kanaya-Okayama equation [4].

Type of presentation: Poster

IT-4-P-1607 Application for Low Energy STEM with the In-lens Cold FE-SEM

Sunaoshi T.1, Orai Y.1, Ito H.1, Okada S.1, Ogashiwa T.1, Konno M.1
1Hitachi High-Technologies Corporation, Ibaraki, Japan
sunaoshi-takeshi@naka.hitachi-hitec.com

Inorganic carbon materials (primary carbon nanotubes and graphene) and organic polymeric materials are being developed more actively. The demands for fine structural, elemental, and chemical characterization of these materials by electron microscopes are rapidly increasing. These requirements have increased the demand to achieve high resolution STEM imaging at low accelerating voltages. It is necessary to determine methods to improve the contrast intensity at low accelerating voltage operation without loss of resolution. In order to respond to such demands, we have developed the Hitachi SU9000 (Figure 1), a cold FE-SEM (CFE-SEM) with an in-lens type of objective lens, capable of high resolution phase contrast STEM imaging. Through using this technique on this microscope, it is possible to routinely achieve lattice resolution of the graphite {002} planes with a spacing of 0.34 nm. In this study, we improve the observation conditions for obtaining enhanced lattice resolution in STEM imaging at an accelerating voltage of 30 kV. Additionally, we have shown the effectiveness of this method for imaging inorganic carbon based materials. Figure 2 shows a simplified lens diagram of the SU9000. By using newly optimized lens parameters and a specialized sample stage which reduces the distance between the objective lens and sample, the Cs was lowered from approximately 2 mm to 1 mm. Figure 3 shows a high resolution BF-STEM image with its inset Fourier transform (FFT) image, observed along the Si<110> zone axis at an accelerating voltage of 30 kV. The sample was prepared using the NB5000 FIB-SEM equipped with a unique micro-sampling system. The specimen was thinned down to approximately 30 nm thickness. (a) is the standard condition (WD: 3 mm,Cs:2 mm) and (b) is the optimized condition (WD: 1.8 mm, Cs: 1 mm). Both (a) and (b) imaged the Si {111} plane, which has a spacing of 0.314 nm, and reflection spots corresponding to 0.314nm were confirmed from FFT images. However, in the optimized condition, not only Si {111} planes (corresponding to 0.314 nm) but also the {002} planes (corresponding to 0.272 nm) are detected from FFT. This confirms that the image resolution is improved by reduction of Cs. Next we applied the optimized condition to a graphene sample. Figure 4 shows a high resolution BF-STEM image with its inset FFT image, the multi-layer graphene membrane was clearly observed at an accelerating voltage of 30 kV. Lattice fringes were easily observed and the reflection spots corresponding to 0.213 nm were successfully confirmed. These results reveal the potential for high contrast visualization without loss of resolution for any carbon-based materials and the latest semiconductor devices with minimal radiation beam damage.


The multi-layer graphene membrane specimen was kindly supplied by Dr. Tsuyohiko Fujigaya of the Department of Applied Chemistry Graduate School of Engineering, Kyushu University.

Fig. 1: General view of SU9000 In-lens FE-SEM.

Fig. 2: Configuration of SU9000.

Fig. 3: BF-STEM images and FFT images of Si <011> single crystal. (Accelerating voltage is 30 kV )

Fig. 4: BF-STEM image and FFT image of graphene. (Accelerating voltage is 30 kV )

Type of presentation: Poster

IT-4-P-2405 Application of Multi-Tilt Specimen Stage for Advanced Electron Channeling Contrast Technique

Dluhoš J.1, Sedláček L.1, Ižák T.1, Hrnčíř T.1, Jiruše J.1
1TESCAN Brno, s.r.o., Brno, Czech Republic
jiri.dluhos@tescan.cz

New developments for a selected area electron channeling pattern (SACP) acquisition and electron channeling contrast imaging (ECCI) in the scanning electron microscope (SEM) are presented. Novel approaches for electron channeling contrast formation are introduced using multi-axial multi-tilt specimen stage.

The multi axial piezo-driven specimen stage allows high precision movement with all six degrees of freedom. Among other applications, it allows bi-axial tilting around any selected point on the sample. Such a technology enables precise selection of beam to surface angle which is the core principle of electron channeling contrast formation, as shown in Fig. 1.

The well-known "rocking beam" technique for electron channeling pattern (ECP) described e.g. in [1] is based on a special mode of scanning in the SEM. Limitation of this technique is given by spherical aberration of the objective lens, which restricts its use mainly to single crystals. A dedicated Cs corrected rocking beam mode was developed for TESCAN field emission microscopes for acquisition of SACP from a very small area as shown in Fig. 2. The practical use of this correction was demonstrated on polycrystalline samples in [2]. Further extension of the rocking angle by the use of stage tilt was also tested.

Specific properties of the ECCI technique in SEM for the observation of near surface defects are explained. The relation of ECP to the formation of ECCI is crucial for understanding the whole ECCI phenomenon. The oriented ECCI technique for reaching suitable diffraction condition as described in [3] was applied. Advantage of combination of Cs corrected SACP for oriented ECCI technique is shown. Newly, the use of precise bi-axial specimen tilt for oriented ECCI is demonstrated in Fig. 3.

Furthermore, new techniques for ECCI contrast improvement, such as color coded multi-axial specimen tilt, are introduced. The sample tilt angle is coded according to HSV color or RGB model to improve the informational depth of the micrograph (see Fig. 4).

References:

[1] A J Wilkinson et al, Micron 28 No. 4 (1997) p. 279.

[2] J Dluhoš et al, METAL Conference Proceedings (2012) p.453.

[3] B A Simkin et al, Ultramicroscopy 77(1-2) (1999) p. 65.


The research has been supported by the Technological Agency of Czech Republic TE 01020233(AmiSpec)

Fig. 1: Schematic diagram of forming the channeling contrast in relation to deviation from the Bragg condition. Image by Wilkinson et al. [1].

Fig. 2: Comparison of ordinary rocking beam mode without correction of spherical aberration (left) and a Cs corrected ECP mode (right). Images taken on polycrystalline stainless steel with grain size about 20 µm.

Fig. 3: ECCI imaging with the use of SACP a) navigation to diffraction condition on SACP (edge of the band, using a multi axial stage tilt. b) ECCI image of crystal defects

Fig. 4: Composite ECCI micrograph of copper sample with randomly oriented grains using the color coded tilt of the stage. R,G,B – images acquired with specimen tilt from -5° to +5°.

Type of presentation: Poster

IT-4-P-1688 High Sensitivity and Minimum Acquisition Time with the Annular EDX Pole Piece Silicon Drift Detector “Rococo2”

Liebel A.1, Eckhardt R.1, Bornschlegl M.1, Bechteler A.1, Niculae A.1, Soltau H.1
1PNDetector GmbH, Munich, Germany
andreas.liebel@pndetector.de

The operation conditions of Scanning Electron Microscopes (SEM) have changed a lot over the last years and many applications have to deal with very low primary beam energies and currents. Modern Energy Dispersive X-ray (EDX) detectors have to accomplish these tasks i.e. they need to support a large geometric collection efficiency (solid angle) in order to enable fast measurements even at weak X-ray intensities. This demands not only for large area detectors but also for intelligent detector designs.
The annular Silicon Drift Detector (SDD) “Rococo2” uses a highly optimized geometry which covers a very large solid angle. It consists of 4 cloverleaf shaped SDD cells combined on one monolithic chip with a total sensitive area of 60 mm² and a center hole. The detector is shown in Figure 1a. It can be positioned right underneath the pole piece extremely close to the sample which results in a very large solid angle up to 1.4 sr. [1] Comparing this number with the solid angle of a conventional 10 mm² SDD detector of typically 0.01 sr it is obvious that the Rococo2 detector can deliver 100 times larger signal intensities at the same measurement time and conditions. Figure 2 shows EDX mappings of a duplex brass sample which illustrate this benefit.
High energetic electrons which are backscattered from the sample are typically filtered by using magnetic electron traps in front of the EDX detector. In case of the Rococo2 detector this is not possible because the magnetic field would disturb the electron beam. In this case the Backscattered Electrons (BSE) are filtered by hardware filter foils which stop the electrons while transmitting the X-ray photons. We will present measurements with a combination of different filter foils made of 2 µm thick Beryllium and 2 µm Mylar for each two detector cells (see Figure 1b). With the used combination of foils a continuous undisturbed X-ray sensitivity down to carbon and boron can be achieved.
We will further present concepts for a combined annular detector for measuring backscattered electrons and X-Rays simultaneously. By increasing the central hole of the EDX detector it is possible to detect backscattered electrons at high take off angles in the central part of the detector. This enables the detection of X-ray and BSE signals at the same time with relatively high collection efficiency and just one single detector head. Figure 3 shows two 2 concepts for such a detector with the BSE detector positioned either above or at the same level as the annular EDX detector. We will show calculations of the solid angle and the collection efficiency of different EDX and BSE detector combinations and evaluate the results by comparing images or spectra.

[1] A Niculae et al, Microscopy & Microanalysis, vol. 18 S2 (2012) p. 1202-1203


Fig. 1: a) The annular pole piece EDX detector “Rococo2” inside the SEM and b) a view at the top of the detector showing the collimator with a filter combination of 2 µm Be and 2 µm Mylar.

Fig. 2: EDX Mappings of a duplex brass sample showing α and β phases with different concentration of copper and zinc.The left image was obtained with a 10 mm² single cell SDD with approx. 0.01 sr solid angle, the right one with the Rococo2 detector with a solid angle of more than 1 sr at the same conditions and acquisition time.

Fig. 3: Schematic drawings of a combined EDX and BSE detector setup with a) the BSE detector positioned above the annular EDX detector and b) the two different detectors positioned on the same level.

Type of presentation: Poster

IT-4-P-1795 SEM images using an energy/angle selective electron detector.

Otsuka T.1, Nakamura M.1, Yamashita K.1, Hara M.1, Timischl F.2, Honda K.1, Kudo M.2, Kitamura S.1
1JEOL Ltd., Tokyo, Japan, 2JEOL Technics Ltd., Tokyo, Japan
tootsuka@jeol.co.jp

Scanning electron microscopes (SEMs) are usually equipped with two types of detectors: secondary and backscattered electron detectors. The former produces secondary electron images (SEI) rich in topographic information[1, 2], whereas the latter produces backscattered electron images (BEI) rich in composition information[3]. Recently, however, a few other detectors have been installed in addition to these two types of conventional detectors[4]. In these practical detectors, however, it is difficult to see directly the effect of energy and take-off angle of the emitted electrons on the image contrast. In this study, an electron detector was designed and experimentally manufactured to detect electrons emitted in a definite, variable range of energy and take-off angle.


Figure 1 shows a schematic diagram of the newly designed electron detector, the E-θ detector, which can detect electrons emitted from a sample with a selected range of energy and take-off angle. The E-θ detector consists of a slit plate, inner and outer electrodes in a cage, and electron detectors. The slit plate is placed at a lower part of the E-θ detector. It serves as a selector of take-off angle. The range of take-off angle is selected mechanically by sliding the slit plate as shown in Fig. 1 and is measured from the horizontal direction parallel to the sample surface. The selection of electron energy is made by applying voltage to the inner and outer electrodes in the cage in accordance with electron take-off angles. In the case of the low angle range, positive and negative voltages are applied to the inner and outer electrode, respectively, as shown in Fig. 1(A), so that electrons are deflected towards the inside with the increasing amount of deflection with decreasing energy. In the case of the high angle range, the polarity of the applied voltage is reversed as shown in Fig. 1(B), so that electrons are deflected towards the outside. To enhance the electron detection, a metal mesh is placed in front of the electron detectors and a voltage of 2 keV is applied between the detectors and the metal mesh. When a series of concentric ring electron detectors with different diameters are placed at the upper part of the cage, each ring detector can collect electrons with a specific range of energy determined by the voltage applied to the two electrodes. As a preliminary study, we acquired images using commercial Si-photodiode (SiPD) as the electron detectors, as shown in Fig. 2.

[1] M. Kotera et al., Scanning Microscopy Supplement 4 (1990) p. 111.
[2] Y. Lin, D. C. Joy, Surface and Interface Analysis 37 (2005) p. 895.
[3] M. D. Ball, D. G. McCartney, Journal of Microscopy 124 (1981) p. 57.
[4] S. Asahina et al., Microscopy and Analysis, Nanotechnology supplement November (2012)


Fig. 1: Schematic diagram of the E-θ detector.

Fig. 2: SEM images acquired with the E-θ detector. The sample is a spherical single crystal of tungsten.

Type of presentation: Poster

IT-4-P-1825 The study of extreme low landing voltage scanning electron microscopy

Sakuda Y.1, Asahina S.1, Kazumori H.1, Kawauchi K.1, Nokuo T.1, Charles F.2
1JEOL.ltd, 3-1-2 Musashino, Akisima, Tokyo 196-8558 JAPAN, 2JEOL SAS, Espace Claude Monet-1, allee de Giverny, Croissy-Sur-Seine 78290 FRANCE
ysakuda@jeol.co.jp

The development of low voltage (LV) FE-SEMs have been in progress, and spatial resolution better than 1 nm can now be achieved even at 1 kV. General LV FE-SEMs are available for a sample surface observation and low voltage EDS analysis. Furthermore, by choosing appropriate observation conditions, we can selectively obtain different information such as material topography and composition. In this report, authors focused on observations at extreme low impact electron-energy in order to obtain information from surface. For example, the length of electron mean free path at 100 eV in a solid sample is smaller than 1 nm [1]. So that it is expected to observe clearly surface morphology. Moreover, we can expect what the low impact electron energy show less electron beam damages in general [2].

Recently, we can use combined lens both electrostatic and magnetic to minimize Cc as well as field emission type emitter with high brightness [3]. So-called Super Hybrid Lens (SHL) is equipped on JSM-7800F. In addition, negative surface potential can be applied on the specimen surface by so-called Gentle Beam mode (GB). Therefore, the impact electron energy to the specimen surface can be reduced down on 10 eV keeping the probe size small with high coherency.

Figs.1 (a) and (b) show carbon nanotubes observed at 80 eV and 500 eV. Those images clearly show their shapes even at 80 eV. Fig.1 (a) shows less edge effect compared with (b) due to small penetration depth of electrons at lower impact electron energy. The intensity profiles along the lines shown in Figs.1 (a) and (b) are shown in Fig.1 (c). The 80 eV image shows higher contrast than the 500 eV one. We assume that is due to higher efficiency of interaction with carbon materials at lower impact electron energy.

Fig.2 (a) shows a low magnification image of meso-zeolite (LTA) at 80 eV. Figs.2 (b) and (c) show high magnification images at 80 eV and 500 eV. All images clearly show topological information at meso-LTA. Especially, the image of 80 eV shows less edge effect due to small interaction volume. That is a useful feature of LV FE-SEM because it can reveal fine edges and give high accurate measurement of nano porous materials. One other useful feature is the reduction of specimen damage due to electron irradiation. The gap shown between two arrows in a circle in Figs. 2 (b) and (c) is observed to be wider in the latter than in the former. The observation indicates less electron damage at 80 eV.

References:

[1] C. R. Brundle, J. Vac. Sci. Technol., 11, 212 (1974)

[2] L. Reimer, Scanning Electron Microscopy: Physics of Image Formation and Microanalysis, 2nd ed., Springer,

Berlin, New York, (1998)

[3] J. Frosien, J. Vac. Sci. Technol. B 7 (6), Nov/Dec (1989)

[4] O. Terasaki, JEOL News, (2013)


Fig. 1: Low impact electron energy on Carbon nanotubes

Fig. 2: Images of meso-zeolite (LTA) at two impact electron energy of 80 eV (a,b) and 500 eV (c).

Type of presentation: Poster

IT-4-P-1829 Determination of the second critical energy of primary electrons in relation to dielectric thickness and angle of incidence

Evstaf'eva E. N.1, Rau E. I.1, 2, Tatarintsev A. A.2
1Faculty of Physics, M. V. Lomonosov Moscow State University , 2Institute of microelectronics technology and high purity materials RAS
rau@phys.msu.ru

At the second crossover energy E2C of incident electrons the equilibrium is maintained between the electron probe currents I0, the second electron emission (SEE) currents I0δ, the backscattered electron (BSE) currents I0η, as well as the leakage currents IL and the displacement currents Id that are responsible for the accumulated charge Q. At the equilibrium state the equality I0=I0(δ+η)+IL+Id is fulfilled, while for the target remaining uncharged the condition δ+η=1, VS=0 is valid, where δ and η are the emission coefficients of SE and BSE.

Experimental results for dielectrics, in the case when incident electrons impinge on the sample surface at the angle α, can be described by the following semi-empirical expression: E2C=E2C(0)exp[(ln(R2C/2λ))(1-cosα)], (1)

where λ is the effective emission depth of SE, R2C=76E01,67 ρ is the depth of penetration of primary electrons with the energy E0, ρ is the specific density of the dielectric material. As an example, fig.1(a) shows the experimental dependence and the dependence calculated by formula (1) of the second critical electron energy E0=E2C on the angle of incidence α, for the target potential VS=0, i.e. when the target remains uncharged. Fig.1(b) shows the dependence of the energy E2C on the angle of incidence α for ungrounded metals.

Consider the dependences of the VS of PММА films with the thickness d on a silicon substrate at the electron energy E0.

The experimental results are in qualitative agreement with the calculated results as shown in fig.2(a), presenting the dependences VS(d) for MICA plates 2 to 30 μm thick and for PММА films 0.4, 1.4, 2.7, 4 μm thick on the Si-substrate.

At the radiation energies E0 in the range of 0.5–1.0 keV the negative charging begins (note that according to previous views, positive charging was expected because E0<E2C). At this energy the sign polarity of VS changes, i.e. at the point where VS=0 V. For PММА this value lies in the range E0=0.4–0.6 keV, with the thicknesses of the layers of positive and negative charges and the values of these charges are approximately equal (λ≈R0, Q+=Q-), which is responsible for the total absence of charging. In the region of 1 keV <E0<Ecr2 one can clearly observe negative charging, and the greater d, the higher the value of -VS. This range corresponds to the condition λ<R0<d. The value of -VS at first increases and reaches the maximum, then as E0 and R0 grow, it starts decreasing slowly in the absolute value and reaches zero at the points of Ecr2, that are different for each film thickness d. These points can be used in high-voltage lithography, because it is at these values of R0≥2d that the conduction current IT is generated and it carries excessive negative charges (electrons) onto the substrate, hence VS=0 V.


Fig. 1: Characteristics of the value of the second critical electron energy E2C as a function of angle of incidence α for dielectrics (a) and ungrounded metals (b).

Fig. 2: (a) - Dependence of surface potential of dielectric films on their thicknesses: (1) VS(d) for mica (plot 1) and for PММА (plot 2). (b) - Dependences of surface potential VS on incident electron energy E0 for PММА films of different thickness d on Si-substrate.

Type of presentation: Poster

IT-4-P-1852 New possibilities of SEM for two-channel detection of energetically filtered secondary and backscattered electrons

Rau E. I.1, Kupreenko S. U.1, Tatarintsev A. A.1, Zaitsev S. V.1
1Faculty of Physics , M. V. Lomonosov Moscow State University
rau@phys.msu.ru

In this paper we present a preliminary study of new potentialities of SEM – microtomograph equipped with a toroidal spectrometer of electrons and two detector systems based on microchannel plates (MCP). A new modification of the instrument is shown in fig.1. Electron probe 1 scans across the surface of the sample under investigation 3. Toroidal spectrometer in a case 4 is mounted under SEM objective lens 2. SE an BSE emitted from specimen 3 pass through annular inlet slit 6 and are energy-separated in toroidal capacitor 5 and pass through outlet annular apertures 7 and 8. The energy filtered SE and BSE are detected by two MCP 9 (A and B) placed opposite each other. The signals A and B from these detectors can be sent either to PC 11 to record the spectra or to the SEM display. Using block 10 one can do the operations of addition (A+B) and subtraction (A-B) of signals. It is known that such operations allow us to obtain contrast from either the chemical composition of the specimen (Z-contrast) or the surface topography. In our case the contrast is enhanced and allows unique interpretation owing to filtration of detected electrons in a narrow energy window. The electrons with small energy losses escape mostly from the subsurface region and are modulated in escape angles, which favors domination of topographic contrast. The electrons, which lost considerable amounts of their energy, are emitted from much deeper regions and therefore mostly produce Z-contrast. Addition and subtraction of signals from the two oppositely oriented detectors enhances this effect considerably.

The examples presented in fig.2 and fig.3 show fragments of the sample having heterogeneous composition, consisting of the alloy of different materials, in particular, Cr, Si, Cu, W.

The images shown in fig.2 are obtained at the primary electron beam energy E0=10 keV, the current I0=1nА, the energy of filtered BSE forming the image EBSE=8 keV ((a) and (b)), and at the SE energy ESE=4 eV ((c) and (d)). One can see that the image contrast (obtained from signal addition) differs greatly in element composition and in surface topography (obtained from signal subtraction) for both BSE and SE regimes.

Fig.3 presents the images of another region of the sample taken at different energies of E0. The general image of this region demonstrated in fig.3a is taken in the standard SE mode in SEM at E0=5 keV. Fig.3b shows the image taken in the BSE added signal at E0=5 keV, in fig.3c – at E0=15 keV, and in fig.3d – in the subtracted signal. The fact that, contrast in filtered BSE and SE is higher and more informative than that obtained with standard signals and standard detectors in SEM makes it possible to more accurately visualize and reconstruct the sample 3D surface profile both in BSE and SE.


Fig. 1: Scheme of the spectrometer-microtomograph in SEM: 1–electron probe, 2–objective lens in SEM, 3–sample, 4–shielding case, 5–toroidal electrodes, 6,7,8–input and output annular slit, 9–MCP, 10–signal addition/subtraction block, 11–PC or SEM monitor, 12 – high-voltage power unit of the spectrometer, 13 – semispherical grid for SE potential contrast.

Fig. 2: Images of the sample of complex composition taken in the BSE mode with signal addition A+B (a) and subtraction A-B (b). Images in SE with signal addition A+B (c) and subtraction A-B (d).

Fig. 3: Images of complex sample in standard SE-mode (a) and BSE - filtred mode (b, c, d).

Type of presentation: Poster

IT-4-P-1865 Low-Damage SEM Imaging of Organosilicate Glass Thin Films in Semiconductor Industry

Garitagoitia Cid A.1,2,3, Muehle U.2, Rosenkranz R.2, Zschech E.1,2,3
1Technische Universität Dresden, Dresden, Germany , 2Fraunhofer Institute for Ceramic Technologies and Systems IKTS, Dresden, Germany, 3Dresden Center for Nanoanalysis (DCN), Dresden, Germany
aranzazu.garitagoitia@ikts-md.fraunhofer.de

On-chip interconnect stacks of high-performance microelectronic products like microprocessors consist of Cu interconnects and insulating organosilicate glass (OSG). Dielectrics with extremely low dielectric permittivity (k value) are needed to reduce signal delay time and cross-talk in on-chip interconnects systems. The OSG thin films are either dense (so-called low-k materials) or porous (so-called ultra-low-k materials).

Imaging of dense and porous OSG thin films with Scanning Electron Microscopy (SEM) is necessary in semiconductor industry for process monitoring and physical failure analysis. Due to weak chemical bonding in the glass network, these materials show strong degradation effects when observed in SEM, caused by electron-material interaction. Particularly, the glass network is densified during the electron beam application to the sample, which phenomenologically causes a significant shrinkage of the material. This shrinkage avoids e. g. a quantitative determination of geometric features in semiconductor structures, which is required for process monitoring. Imaging with reduced primary beam energy mitigates the materials damage; however, the spatial resolution is usually reduced at lower accelerating voltages. In this study, spatial resolution and OSG thin film degradation during SEM imaging are studied systematically as a function of the primary beam energy. A Carl Zeiss SEM/FIB system NVision40 with Gemini column and three types of detectors is used, the conventional Everhart-Thornley detector, the inlense detector and the energy selective backscattering (EsB) detector. The optimum parameters for SEM imaging of OSG thin films are provided for several types of materials.


We kindly thank Carl Zeiss Microscopy GmbH for funding the investigations in the framework of the project “Untersuchungen zur Charakteristik und Applikation des EsB Detektors”.

Fig. 1: Schematic of the Zeiss Gemini Column with three different electron detectors

Fig. 2: Crack in low-k dielectric after one single scan in SEM

Type of presentation: Poster

IT-4-P-1912 Ionic Liquid Preparation for SEM Observation of Minute Crustacean

Shiono M.1, Sakaue M.1, Konomi M.1, Tomizawa J.2, Nakazawa E.1, Kawai K.3, Kuwabata S.4
1Tokyo Solution Lab., Hitachi High-Technologies Corp., Kawasaki, Kanagawa, Japan, 2Hitachi High-Technologies Corp.,hitachinaka, Ibaraki, Japan , 3Miyoshi Oil & Fat CO.,LTD. Katsushika, Tokyo, Japan , 4Graduate School of Engineering, Osaka University, Yamadaoka, Suita, Osaka, Japan
sakaue-mari@naka.hitachi-hitec.com

     Electron microscopic observation is often possible along the surface of an arthropod, however applying common fixative media present difficulties when penetrating beyond the exoskeleton. Ionic liquids (ILs) are unique in that they are incombustible, non-volatile, and have high ionic conductivity. The application of ILs as part of sample preparation for EM were conducted with some particles dispersed in ILs(1) and observed by TEM, including IL wetted seaweed with observation by SEM(2). Recently, we have developed a new ionic liquid (IL) HILEM, IL1000 for EM observation. The ionic liquid, IL1000 has been designed for EM preparation, having both a high level of safety and high solubility, resulting in a high suitability for biological tissue preparation. In this experiment, minute crustaceans were immersed in 10 % IL1000 diluted with distilled water for a period of 60 minutes to 3 hours. For surface observation retention, any extra IL covering samples were removed by an absorbent cloth. The samples were observed by the Hitachi SU3500 at an acceleration voltage of 5 kV in high vacuum condition.
     Figure 1 shows the secondary electron images of the minute crustaceans Gammaridea. Figure 1(a) is an image of the entire body of the Gammaridea, whose individual appendages are not easily distinguishable due to the overlapping of appendages. The imaged sample was removed from the SEM and tweezers were used with the aid of a binocular to separate its appendages from the body. Figure 1(b) is the separated appendage (the second thoracic appendage) of the Gammaridea and the attached organ shown by arrow in Figure 1(b). The organ shown functioned in protecting the egg, therefore it was determined that this specimen is female. The results show that the IL can deeply penetrate the specimen which aids electron conductivity inside the specimen. Figure 2 is the secondary electron images of the Tanaidacea. Figure 2(a) is the whole image of the Tanaidacea orientated to observe the ventral side. The female of the Tanaidacea has the brood chamber in the thorax (arrow) region. As aforementioned above, the observed specimen was removed from the SEM to separate its brood chamber. Figure 2(b) is a SEM image of the eggs that were removed from the brood chamber. Since the sample was soaked by IL, the sample is resistant to rapid dehydration while under vacuum, including soft materialsuch as eggs can be preserved by this technique. We emphasize that the IL1000 is a useful media for SEM observation of some soft and delicate biological material.

References
[1] E. Nakazawa. et al, Proceedings of the sixty-fourth Annual Meeting of The Japan Society of Microscopy. p 136 (2008)
[2] S. Kuwabata. et al, Kenbikyo, 44: p 61-63. (2009)

 


Fig. 1: Figure 1. Secondary electron image of Gammaridea treated by the ionic liquid. (a) the whole image of the Gammaridea, where a lot of appendages overlap. (b) The separated second thoracic appendage (arrow) functioned to protect the egg. Instrument: SU3500, Acc. Volt. 5 kV, Magnification: x 32(a), x75(b).

Fig. 2: Figure 2. Secondary electron images of Tanaidacea treated by the ionic liquid. (a) the inclined whole image holding the brood chamber (arrow). (b) the eggs scraped out of the brood chamber. Instrument: S-3400N, Acc. Volt. 5 kV, Magnification: x 35(a), x200(b).

Type of presentation: Poster

IT-4-P-1917 New Preparation Method using Ionic Liquid for Fast and Reliable SEM Observation of Biological Specimens

Sakaue M.1, Shiono M.1, Tomizawa J.2, Nakazawa E.1, Kawai K.3, Kuwabata S.4
1Tokyo Solution Lab., Hitachi High-Technologies Corp., Kawasaki, Kanagawa, Japan , 2Hitachi High-Technologies Corp., Hitachinaka, Ibaraki, Japan, 3Miyoshi Oil & Fat CO., LTD., Tokyo, Japan, 4Graduate School of Engineering, Osaka University, Osaka, Japan
sakaue-mari@naka.hitachi-hitec.com

    For SEM observation, it is necessary for biological specimens to be treated with several types of preparation media to preserve their shape under vacuum. Ionic liquids are unique materials because of their natural incombustibility, non-volatility, and high ionic conductivity. Here, we used an ionic liquid to prepare samples for EM observation [1] [2]. The ionic liquid, IL1000 has been designed for use in EM sample preparation with a high level of safety and high solubility [3].
     Figure 1 shows the SEM images of a Helicobactor Bilis sample prepared by conventional procedures. To preserve the flagella structure, the sample was immobilized on the cover slip coated with poly-L-lysine, and freeze-dried after fixation with 2 % glutaraldehyde (GA) in 0.1 M phosphate buffer and dehydrated with acetone in descending concentrations. The conventional sample preparation method takes approximately 8 hours. The surface of the cell body and some flagella are clearly observed (fig.1). On the other hand, in the ionic liquid (IL) method, the sample fixed by the buffered 2 % GA is immersed in 10 % IL1000 solution for 15 minutes and dropped onto filter paper. The sample is then directly transferred into the SEM without further drying. The resulting SEM image of this sample clearly shows the helical shape of the bacteria and flagella (fig. 2). Figure 3 shows the SEM images of mold growing on a rice cake. The square cut sample x5mm2, was immersed in the 10 % IL1000 solution for approximately 4 hours, and then directly transferred into the SEM. Fine threads protruding from the rice cake are clearly observed (fig.3) and the higher magnification image shows the clear and smooth surface of the spores.
    These results indicate that the IL method for biological sample preparation greatly reduces preparation time, and is additionally better at preserving the sample’s original shape in the SEM.

References
[1] S. Kuwabata. et al, Chem. Lett., 35, p600-601. (2006)
[2] E.Nakazawa.et al, Proceeding of the Fifty-sixth Symposium of the Japanese Society of Microscope., 47-2, p92-95. (2013)
[3] K.Nimura. et al, Hitachi Hyoron., Vol.95, 9, p26-31. (2013)

 


Fig. 1: SEM images of Helicobactor Bilis prepared by conventional procedures  (Frozen dried sample)                      Instrument: SU6600, Acc. Volt. 1 kV, Magnification: x 30,000    Sample: Courtesy of Prof. Yoshiki Kawamura, Aichigakuin University

Fig. 2: SEM images of Helicobactor Bilis treated by the ionic liquid IL1000                                                               Instrument: SU6600, Acc. Volt. 1 kV, Magnification: x 25,000    Sample: Courtesy of Prof. Yoshiki Kawamura, Aichigakuin University

Fig. 3: SEM images of mold growing in the rice cake treated by the 10 % ionic liquid IL1000                                   Instrument: SU3500, Acc. Volt. 3 kV, Magnification: x 3,000

Type of presentation: Poster

IT-4-P-2707 Low-Voltage Imaging of Non-Conducting Samples

Beránek J.1, Havelka M.1, Jiruše J.1
1TESCAN Brno, s.r.o., Libušina třída 1, Brno, Czech Republic
miloslav.havelka@tescan.cz

In the past years, considerable attention has been drawn to imaging of non-conducting samples without prior application of conductive coating. Conditions of low voltage microscopy allow such observation with its main benefits: sensitivity to the surface details and possibility to reach charge balance under which the charging of the sample is diminished [1, 2].

Uncoated non-conducting samples often exhibit undesirable charging that prevents the observation of finer details. This effect can be suppressed by using specific landing energy for which the total flow of electrons from the sample equals the charge coming into the sample. An example of the charge balance for nylon fibers is shown in Figure 1: a) exhibits positive charging effects, b) is an illustration of charge balance at 1200 V in agreement with [2], whereas c) has visible signs of accumulation of negative charge.

Conditions for charge flow equilibrium for the non-conducting materials generally lie in low voltage region [2]. To maintain the quality of imaging, preserving high resolution at low acceleration voltages is crucial. In Figure 2, we present images taken in low voltage regime. In Figure 2 a) uncoated polystyrene balls are shown. At 4.2 kV we can see fine details of their surface roughened by etching. The resolution at low voltages can be enhanced in the Beam Deceleration Mode (BDM) [3]. Figure 2 b) shows the structure of TiO2 imaged with the BDM at 800 V. In this mode, the electrons are maintained at higher energy during their path through the column and they are decelerated just after they leave the objective lens. BDM supports further lowering of landing energy, automatically to 50 eV and manually to 0 eV. Figure 2 c) shows para-hexaphenyl imaged at 20 eV. These images were taken by an ultra-high resolution microscope MAIA [4] by TESCAN, which has guaranteed resolution 1.4 nm at 1 kV.

Secondary electrons reveal sometimes surprising amount of details when primary beam interacts only with surface layers of material [1]. In Figure 3, the comparison of cracked oxidized copper imaged at acceleration voltages a) 20 kV, b) 10 kV and c) 2 kV is given. As can be seen in Figures 3 a) and 3 b), the shapes and edges of larger structures are well distinguishable and coarse surface is visible. Figure 3 c), taken at 2 kV, shows detailed structure of the studied object. In comparison with Figures 3 a) and 3 b), in Figure 3 c) the edges lose brightness and the contrast of surface cracks and contours is predominant.

References:

[1] I Müllerová et al, Adv. in imaging and electron physics 128 (2003) p. 310.

[2] D C Joy, Micron 27.3 (1996) p. 247.

[3] J Jiruše et al, 15th Eur. Microsc. Congress Proceedings (2012) p. 165.

[4] J Jiruše et al, Microsc. Microanal. 19 (Suppl 2) (2013) p. 1302.


The support from FR-TI2/736 (MOREMIT) funded by the Ministry of Industry and Trade of the Czech Republic is acknowledged.

Fig. 1: Charging artifacts of nylon fibers at acceleration voltages a) 900 V, b) 1200 V and c) 1500 V. Dark areas and lines in a) are an evidence of positive charging while localized brighter areas in c) are due to negative charging. Image b) at critical voltage shows least charging artifacts.

Fig. 2: a) Uncoated polystyrene balls at 4200 V, b) TiO2 with BDM at 800 V and c) Fiber-like structure of Para-hexaphenyl imaged at 20 V with BDM. Lowering the acceleration voltage makes fine surface details of the presented non-conducting samples clearly visible. Such details are frequently obscured when high acceleration voltages are used.

Fig. 3: Oxidized surface of copper imaged at different acceleration voltages, a) 20 kV, b) 10 kV and c) 2 kV, thus shrinking the interaction volume. The contrast is gradually changed, especially at the edges of surface cracks and contours.

Type of presentation: Poster

IT-4-P-2152 Purification of FEBID gold nanostructures using oxygen plasma

Shawrav M. M.1, Wanzenboeck H. D.1, Belić D.1, Gavagnin M.1, Wachter S.1, Bertagnolli E.1
1Institute of Solid State Electronics, Vienna University of Technology
mostafa.shawrav@tuwien.ac.at

The full text of the abstract is not available. Please contact the presenting author.


Type of presentation: Poster

IT-4-P-2277 Physical background of multiple detection possibilities of the Hitachi SU8030 SEM

Gernert U.1, Berger D.1
1Technical University Berlin, Center for Electron Microscopy (ZELMI), Straße des 17. Juni 135, 10623 Berlin, Germany
ulrich.gernert@tu-berlin.de

The spatial resolution of state-of-the-art high resolution scanning electron microscopes has nearly reached the physical limit of around 1 nm. On the one hand, this has been achieved by optimised electron sources and improved low kV electron optics. On the other hand, advanced and multifunctional detectors have been developed to detect separately energy- and angular-selective secondary electron (SE) as well as backscattered electron (BSE) signals with reduced interaction volumes. While early SEMs used one SE- and one BSE-signal only, a fully equipped modern HRSEM offers the possibility to record up to 20 different signals. Therefore, the skilled operator will find the most suitable detection method to optimise the contrast of the nm-structures under investigations, no matter if they show up due to potential-, edge-, binding energy-, working function-, Z-contrast or anything else.
The aim of this work is the quantitative analysis of nm-structured samples, i.e. all possible signals should be linked to the analysed sample properties and to the physical background of the detection process. Due to the large variety of parameters, this was performed by analysing known samples with the different detection methods of a modern Hitachi SU8030 SEM.
As an example, fig.1 shows a Si-chip with conductive Si-tracks and isolating SiO2-areas around. On top, there is a PMMA polymer layer with cracks. On sample areas, consisting of polymer coating on Si substrate, both detectors, the lower one (mounted at the specimen chamber) and the upper detector (in-lens), show the same image contrast, which in this case is a pure element contrast. On SiO2 substrate, however, using the lower detector to display the cracked polymer leads to the same image contrast as before, whereas the upper detector additionally shows the potential contrast since it records a pure SE1-signal.
Fig.2 shows a Si wafer with square Au pads with 130 nm height. The images are taken with the top detector, which is located well above the upper detector. In non-deceleration mode, a high angle BSE signal is detected with orientation- and Z-contrast (fig. 2b). In deceleration mode, a negative voltage is supplied to the sample decelerating primary electrons as well as accelerating escaping secondary electrons, which subsequently will then be detected with the top detector. In the latter case, the pure SE-signal shows a fine topographic contrast (edges of Au-pads and grain boundaries) and it is sensitive to the surface potential (bright lines in the grooves, fig. 2a).
Further sample properties might be analysed, if primary electron energy, deceleration voltage, detected energy and WD are varied in addition. Consequently, a state-of-the-art SEM provides many complementary imaging modes utilizing its high flexibility.


We kindly acknowledge financial support from the Deutsche Forschungsgemeinschaft (DFG).

Fig. 1: Si-chip analysed with Hitachi SU8030 SEM.  a: Lower detector (SE[L]); accelerating voltage 900 V; working distance 10 mm  b: Upper detector (SE[U]); accelerating voltage 900 V; working distance 10 mm

Fig. 2: Au-pads on silicon analysed with Hitachi SU8030 SEM.  a: Top detector (SE[T]); landing energy 2 keV (accelerating voltage 3 kV; deceleration voltage 1 kV); working distance 1.9 mm  b: Top detector (HA100[T]); landing energy 2 keV (no deceleration); working distance 1.9 mm

Type of presentation: Poster

IT-4-P-3088 STEM in SEM imaging of gold nanoparticles in tissular ecotoxicity experiments

García-Negrete C. A.1, Jiménez de Haro M. C.1, Blasco J.2, Soto M.3, Fernández A.1
1Materials Science Institute of Seville (CSIC - Univ. Seville), Seville, Spain, 2Institute of Marine Sciences of Andalusia (ICMAN-CSIC), Puerto Real (Cadiz), Spain, 3Zoology and Cell Biology Dept., University of the Basque Country, Leioa (Bizkaia), Spain
cjimenez@icmse.csic.es

Because of their prospective widespread use, gold nanoparticles (AuNPs) will certainly account for a considerable and persistent nanomaterial input to environmental systems. Therefore ecotoxicological risks in non target organisms associated with AuNPs are showing increasing consideration. Location of AuNPs has been previously studied in our laboratory analyzing slices of gills and digestive gland tissues of the bivalbe Ruditapes philippinarum after “in vivo” exposure experiments. Analysis was carried out by TEM of ultrathin tissue’s slices (80 nm) operating at 80 kV [1].
In this communication we present the results of investigating the use of an “in vitro” methodology associated to the optimization of the STEM-in-SEM technique for the use of a scanning electron microscope (SEM-FEG) in transmission mode and operated at 20-30 kV.
The advantages of STEM-in-SEM over TEM are discussed [2, 3]. The localization of high Z nanoparticles in low Z tissue matrices is presented here by using the STEM-in-SEM coupled to EDX analysis as a powerful technique. In addition we have optimized the measurements with the goal of working with thicker slices. The work with thick samples also avoid the NPs displacement during cutting and increase the possibility of finding NPs when working with low NPs doses (environmental relevant concentrations).
For the optimization of measurements conditions, the resolution in our SEM-FEG has been estimated using Fast Fourier Transform (FFT) algorithms on specific images of our tissue slices. We have used the SMART macro running inside the “SCION Image” program under windows [4, 5]. Working at magnifications over 100 kx, for slices thicknesses of 200-300 nm and operating voltages of 20-30 kV, leads to resolutions below 10 nm (an adequate value for analyzing AuNPs of 23 nm average diameter).
Figure 1 shows a representative image of AuNPs accumulated into the gill tissues after “in vitro” exposures. From the obtained images it was possible to localize AuNPs (see also Figure 2) associated with vesicles (it can be a large phagosome or also exocytosis). Nanoparticles were also found in residual bodies (exocytosis).
In summary this communications presents new results for “in vitro” fast testing and STEM-in-SEM imaging of engineered AuNPs in a tissular ecotoxicity model.

References
[1] CA García-Negrete, J Blasco, M Volland, TC Rojas, M Hampel, A Lapresta-Fernández, MC Jiménez de Haro, M Soto, A Fernández. Environmental Pollution 174 (2013), 134-141.
[2] O Guise, C Strom, N Preschilla. Microsc Microanal 14 (Suppl2) (2008), 678.
[3] A Bogner, PH Jouneau, G Thollet, D Basset, C Gauthier. Micron 38 (2007), 390-401.
[4] C Probst, R Gauvin, RAL Drew. Micron 38 (2007), 402-408.
[5] DC Joy. J Microsc 208 (2002), 24-34.


The authors gratefully acknowledge financial support from the Junta de Andalucia and EU FEDER (project PE2009-FQM-4554 and TEP-217) and the EU FP7 AL-NANOFUNC project (CT-REGPOT2011-1-285895).

Fig. 1: SEM-FEG image (transmission mode) of a 200 nm slice of gills tissue

Fig. 2: EDX spectrum from the area containing the AuNPs

Type of presentation: Poster

IT-4-P-2413 Hair Cuticular Characteristics: Potential of Animal Identification in Primates Order

HING L. H.1, TEO H. C.1, FOONG M. J.1, HUKIL S.1, SAHALAN A. Z.1, WAN NUR SYAWANI W.1, KASWANDI M. A.2, NORMALAWATI S.3
1School of Diagnostic & Applied Health Sciences, Faculty of Health Sciences, Universiti Kebangsaan Malaysia, 2Institute of Medical Science Technology, Universiti Kuala Lumpur, Taman Kajang Sentral, 43000 Kajang, Selangor., 3Electron Microscope Unit, Faculty of Science & Technology, Universiti Kebangsaan Malaysiaa
hing61@yahoo.com

Hair is one of the common trace evidences in a crime investigation due to its easy shedding nature and could be easily transferred between surfaces or left behind at the crime scenes (1,2). Therefore forensic examination of hair sample plays a significant role in investigation of illegal wildlife-related crime cases (3,4). Four species of animals from Primates order were selected with their hair samples examined under scanning electron microscope. Species of Cercopithecidae family include blue monkey (Cercopithecus mitis), banded leaf monkey (Presbytis femoralis) and vervet monkey (Chlorocebus pygerythrus) (Fig 1) while chimpanzee (Pan troglodytes)(Fig 2) is from Homonidae family. Nol chemical or mechanical cleaning of hair examination was done. In Primates order, all species have  same regular wave cuticular pattern but variation was seen in other features. Blue monkey and vervet monkey have smooth cuticular dorsal margin whereas banded leaf monkey and chimpanzee have rippled structure. Banded leaf monkey  showed intermediate hair cuticular orientation and  is the only species having this characteristic. Statistics analysis proved that average scale layer difference could  be one of the criteria in examination of hair samples.  Comparison showed average scale layer difference of chimpanzee is significant lower than blue monkey and vervet monkey. In Primates order, cuticular dorsal margin, scale position and scale layer difference could be employed to differentiate all four species of animal successfully (4). This study also proved that analysis of cuticular scales pattern and other related characteristics was not affected when conventional cleaning procedure was not employed. Present study indicates that it is possible that positive identification of animal species through hair samples examination using various measurements and examination of hair cuticular characteristics.

References

B. J. Teerink, Atlas and Identification Key: Hair of West-European Mammals, Cambridge: Cambridge University Press (1991).

M.S. Dahiya, S.K. Yadav, Elemental Composition of Hair and its Role in Forensic Identification, Open Access Scientific Reports, 2(4): 10.4172, 721 (2013).

M.S. Dahiya, S.K. Yadav, Scanning electron microscope characterization and elemental analysis of hair: a tool in identification of felidae animal, J Forensic Research, 4(1), 178, 4172/2157-7145.1000178 (2013).

H. Brunner, B.J. Coman, The Identification of Mammalian Hair, Melbourne: Inkata Press (1974).


The author acknowledged the contribution of hairs by the National Zoo of Malaysia

Fig. 1: Hair of vervet monkey

Fig. 2: Hair of chimpanzee

Type of presentation: Poster

IT-4-P-2468 Recent achievements of elemental analysis with low acceleration voltage FE-SEM

Kikuchi N.1, Morita H.2, Ikarashi M.2, Kawauchi K.1, Nokuo T.1, Charles F.3
1JEOL Ltd., 3-1-2 Musashino, Akisima, Tokyo 196-8558 Japan, 2Oxford Instruments KK, IS Building, 3-32-42 Higashi-Shinagawa, Shinagawa-Ku, Tokyo 140-0002 Japan, 3JEOL SAS, Espace Claude Monet-1, allee de Giverny, Croissy-Sur-Seine 78290 France
nkikuchi@jeol.co.jp

Field Emission Scanning Electron Microscopes (FE-SEMs) are widely used for high spatial resolution imaging and analysis. Energy dispersive X-ray spectroscopy (EDS) is a technique used for elemental analysis of a sample, through the detection of characteristic X-rays generated from the sample irradiated by an electron beam.

To obtain the best spatial resolution of imaging, as a matter of course, the accelerating voltage and the beam current of the electron probe should be optimized. In the case of JSM-7800F, the beam current under the optimum condition, especially in a lower accelerating voltage range, was too small to detect characteristic X-rays for the elemental analysis with a reasonably good signal to noise (S/N) ratio. As to the accelerating voltage, landing energy of the electron beam, in fact, has to be changed to be optimized, depending on the excitation potential of the characteristic X-rays of specific elements to be analyzed. In the case of analysis, we have developed a new gun, an inlens Schottky plus FE gun, which produces about two orders of magnitude larger probe current in that range of the landing energy with almost the same spatial resolution of imaging as the one under the optimum condition of JSM-7800F. JSM-7800F equipped with this new gun is named JSM-7800F Prime, which has one other key feature installed; a beam deceleration mode, named gentle beam for super high resolution (GBSH), in which negative potential can be applied to the sample surface to decelerate incoming probe electrons on to the sample. GBSH facilitates high spatial resolution of imaging in low landing energy [1]. Fig.1 shows gold particles on carbon images taken with JSM-7800F and JSM-7800F Prime respectively at the same accelerating voltage of 1kV under the same probe current. The amount of the current was almost two ordered of magnitude larger than the one under the optimum condition in JSM-7800F. The spatial resolution of the latter image is far better than that of the former, of the order of nanometers.

As to the detector of characteristic X-rays, their counting rate is desired to be as large as possible to obtain a better S/N ratio [2]. For this purpose, an improvement has been made to establish an ultrahigh solid angle double EDS detector system from Oxford Instruments.

Fig.2 shows image and their corresponding elemental map of a nanometer size Pt particles on carbon substrates. They were taken with JSM-7800F Prime. A particle with a size of 7nm is clearly observed in the elemental map in Fig.2 taken for 15min in the GBSH mode in 5keV landing energy.

References:

1. L. Reimer, Scanning Electron Microscopy: Physics of Image Formation and Microanalysis, 2nd ed., Springer, Berlin, New York, (1998)

2. Y. Nakajima et.al., Submitted to M&M2014


Fig. 1: Gold particles on carbon images taken with JSM-7800F and JSM-7800F Prime at 1kV under the same probe current.

Fig. 2: Pt particles on carbon substrates. Imaging and EDS analysis (The two EDS detectors of Oxford Instruments, each of which has the area of 150mm2, are installed to form a double detector system with the area of 300mm2 in total ) were made with JSM-7800F Prime in the GBSH mode in 5keV landing energy.

Type of presentation: Poster

IT-4-P-2766 A novel SEM triple in-lens detection system

Wall D. C.1, Vystavel T.2, Tuma L.2, Skalicky J.2, Sasam F.1, Wandrol P.2
1FEI Company, Eindhoven, The Netherlands, 2FEI Company, Brno, Czech Republic
David.Wall@fei.com

The advent of new materials and new techniques in SEM and DualBeam has driven the need for better detection in recent years. Traditionally, in-lens detection systems have focussed on energy selection of signal, due to the small opening angle of signal that can be detected, while modern, below-lens detectors have the benefit of separating angular differences in signal. This has typically meant that the in-lens detection system has had strong benefits for resolving materials with fine difference in the composition, while below-the-lens has been more suited to channelling contrast.
In this abstract, a new type of electron column design is introduced which broadens the spectrum of BSE’s and SE’s that can be detected with the in-lens detectors. The newly introduced NICol SEM column positions the in-lens BSE detector at the lowest point of the column, so that the opening angle of BSE that can reach the detector is far higher than those typically positioned higher up. The benefit of this can be seen in the images of figure 1. where strong channelling contrast is now possible with in-lens detection. This enables the collection of strong grain orientation images even while tilted or in 3D data collection in DualBeam configurations, where previously below-lens detectors is more difficult to use due to possible collisions. Additionally, by segmenting the annular design of this BSE detector into left and right segments, two separate signals can be detected and processed. Adding these these segments delivers material or orientation contrast, while a differential image generates strong topographical contrast. Where topographical images are necessary on charging material, this technique can avoid the charge.

By fully utilizing the experiment geometry, clear separation between high and low energy secondary electrons can also be enabled with the further two in-lens detectors in the SEM. Very low energy, surface sensitive signal will be affected most by the electrostatic field and travels closest to the beam axis. Higher energy secondary electrons less affected by the electrostatic lens are projected onto the middle detector. These effects can be seen in Fig. 2 (a, b) where the lower energy SE image shows excellent surface information while the higher energy SE image shows the best edge contrast. This signal can then be detected on the upper detector. Simultaneous collection of all three of these signals enables the collection of all information in a single scan, reducing charge up effects, and preventing beam damage or contamination to the sample.

 


Fig. 1: SEM image of FIB cross-section through steel sample acquired at 1.8 keV exhibiting strong channeling contrast. This enables clear identification of austenite and ferrite regions.

Fig. 2: Simultaneous SEM images acquired using high energy (left) and low energy (right) secondary electrons using a primary beam energy of 2kV revealing edge and surface sensitivity.

Fig. 3: Simultaneous SEM images acquired using high energy (left) and low energy (right) secondary electrons using a primary beam energy of 2kV revealing edge and surface sensitivity.

Type of presentation: Poster

IT-4-P-2553 Nanostructure Imaging with Topographic and Compositional Contrasts in a Cold-Field Emission Scanning Electron Microscope at Low Accelerating Voltage and with Energy-Filtration

Gauvin R.1, Brodusch N.1, Demers H.1, Woo P.2
1Department of Mining and Materials Engineering, McGill University, Montreal, Quebec, Canada, 2Hitachi High-Technologies Canada Inc., Toronto, Canada
raynald.gauvin@mcgill.ca

For developing new technologies, it is important to characterize the microstructure of materials with high spatial resolution at the nanoscale. To achieve high resolution, field emission scanning electron microscopes (FE-SEM) were developed. These microscopes allow working at low accelerating voltage, below 5 kV, to take advantage of the reduction of the interaction volume with accelerating voltage (from 1 μm in Al at 10 kV to 10 nm at 1 kV). Furthermore, their higher gun brightness compared to conventional thermo-electronic emitters, allow a probe size at the nanoscale. However, technical problems arise when SEM operates at low kV, i.e., the source brightness decreases and the chromatic aberration increases, all SEM parameters being equal. Using deceleration mode minimizes these problems and further improvement is achieved by using a cold-field emitter, which has a smaller energy spread and providing the highest brightness and the smallest source size of a FE-SEM. At low accelerating voltage, the emission volume of backscatter (BSE) and secondary (SEII, emitted by BSEs) electrons signals approach that of SEI (emitted by the primary electrons) signals. However it is not enough to reach the highest resolution. A magnetic field above the sample improves the spatial resolution by collecting mostly high-resolution signals. In addition, the energy-filtration of the electron signals allows selecting the type of contrast detected: topographic, compositional, or crystallographic.

Examples of high spatial resolution imaging are shown in Figure 1. Topographic imaging at a very low accelerating voltage of 50 V is possible with the deceleration mode with still an excellent spatial resolution of 2.8 nm as calculated with SMART-J (Figure 1A). The energy-filtration allows the observation of small compositional contrast as shown in Figure 1B where Al3Li precipitates (δ’) were observed in an AA2099 Al-Li-Cu alloy. A resolution of 2.2 nm was obtained for a combination of SE and BSE signals with a mix of topographic and compositional contrasts (Figure 2A). Simultaneously, an energy-filtered BSE signal was acquired with a resolution of 2.7 nm and a compositional contrast was observed (Figure 2B). Furthermore, Monte Carlo simulations were used to understand and to optimize the SEM parameters of these different imaging modes.

The HITACHI SU-8230 CFE-SEM provides low accelerating voltage, deceleration mode and energy-filtration of the electron signals and thus allows the characterization of the microstructure of materials with high spatial resolution at the nanoscale with various types of contrasts. The development of these new technologies permits to extend the imaging capabilities of the SEM towards new nanoscale applications.


Fig. 1: High resolution micrograph obtained with a CFE-SEM. (A) SE micrograph of a CNT decorated with Pt nanoparticles was acquired at 50 V in deceleration mode with the top detector. (B) Energy-filtered BSE micrograph of an AA2099 Al-Li alloy acquired at low energy with the upper detector.

Fig. 2: High resolution micrograph obtained with a CFE-SEM. CNTs decorated with Pt nanoparticles micrographs were acquired at 1 kV with: (A) combination of secondary and backscattered electron (SE+BSE) signals by the upper detector; (B) energy-filtered BSE signal by the top detector.

Type of presentation: Poster

IT-4-P-2654 POROSITY DETERMINATION ON IRON ORE PELLETS USING OPTICAL MICROSCOPE AND ELECTRON MICROSCOPE

Graça L. M.1, Lagoeiro L. E.2, Vicente T. M.3
1Geology Department, School of Mines, Federal University of Ouro Preto, 2Geology Department, School of Mines, Federal University of Ouro Preto, 3Mining Department, School of Mines, Federal University of Ouro Preto.
vicente_thais@hotmail.com

New procedures have been developed with the aim of improving the iron ore characterization and its agglomerated product, the pellets. The pelletizing plants have a considerable importance worldwide, as it becomes feasibly economic to use the fine particles (P90 of 0.045 mm). One of the main physical characteristics of pellets is its porosity, which directly interferes on its geometallurgical quality. Facing the lack of specific equipment to determine the porosity in indurated pellet; two methodologies were evaluated in order to determine the values for this variable: the reflected light optical microscope (OM) with and the scanning electron microscopy with the electron backscatter diffraction (SEM-EBSD). The same areas of the pellets (edge and center) were analyzed in both techniques, and with equal magnification (100X), in order to compare the results. Considering the MO results, mosaics of each examined area were created for subsequent imaging treatment that consisted of manual outline of regions that corresponds to the pore areas. The marked areas were subsequently quantified by specific software and it represented the pore percentages for each evaluated area. Considering the SEM-EBSD results, indexing maps of the crystal lattices of hematite, magnetite, wusthite and quartz phases were produced. In addition to the percentage of each mineral phase, the generated map determines the zero solution, which represents the regions of no indexing and therefore with no mineral phase. The percentage of zero solution in this technique represents the existing pores on the investigated area. It was collected three pellet samples, in each one the center and the edge were investigated. The percentages determined by EBSD were higher than those found by MO, both on border and center areas. The outline of the pores from MO was manual, which depends on the observer judgment and that may influence the final results. On the other hand, the SEM-EBSD does not index the amorphous phases, which is generated by the induration process. Although the amorphous material occurs in low percentages, it increases the zero solution. The determined porosity in both methods was higher on the central region of the pellet, as it was expected, since this region settles the nucleating particles. As in the MO, the average of the results obtained in edge regions of the pellet was 48.8% and in the central region of the pellet 51.8%. From the SEM, the average results obtained at the edge of the pellets was 53.66% and at the center 55.63%. Regions with higher porosity showed more differences in results between the methodologies. The analyzes made by the OM and EBSD showed coherent and consistent data when two methods were compared.


This study was partially supported by the FAPEMIG and the CNPq. All the analyses were performed in the MICROLAB at the Federal University of Ouro Preto.

Type of presentation: Poster

IT-4-P-2700 Volume reconstruction of biological samples by alternate physical and virtual slicing

Hovorka M.1, Boughorbel F.2, Potocek P.2, Cernohorsky P.1, van den Boogaard R.2, Hekking L.2, Korkmaz E.2, Langhorst M.2, Lich B.2
1FEI, Podnikatelska 6, 635 00 Brno, Czech Republic, 2FEI, Achtseweg Noord 5, 5651 GG Eindhoven, Netherlands
milos.hovorka@fei.com

The ability of scanning electron microscopy (SEM) to image large volumes with high spatial resolution, throughput and reliability goes hand in hand with the development of new approaches for data acquisition. Resin embedded tissues stained with heavy metals pose typical challenges with generally low electrical conductivity and charging. One way of imaging them in SEM is based on alternate slicing and imaging of the block-face. SBFSEM (Serial Block-Face SEM) utilizes an ultramicrotome inside the SEM chamber to cut slices of defined thickness. The revealed block-face is scanned and backscattered electrons are collected [1]. The depth resolution is determined by the achievable slice thickness and imposes a limit on voxel isotropy and on the quality of the reconstructed 3D information. The acquisition parameters and hence data throughput depend not only on the sample properties but significantly on the detection part of the microscope as well. Sample charging can be suppressed by working in low vacuum mode, in-situ coating of the surface with a very thin metal layer, by the usage of the accelerating voltage in the low kV range or introducing more metals during sample preparation.

 

We introduce the new integrated solution for SEM volume data acquisition based on a refined SBFSEM technique. It combines physical and virtual slicing which allows for extending the current resolution limit. Virtual slicing is enabled by using the MED-SEM (multi-energy deconvolution SEM) which is a non-destructive technique capable of high-resolution reconstruction of the top layers of the sample [2]. Following each cut, the exposed block-face is imaged and not only one image but a series of images is acquired using different accelerating voltages. Collected images serve as the input for a deconvolution algorithm that computes several subsurface layers. Subsequently, a given thickness of the tissue is removed mechanically using a diamond knife, a fresh block-face is exposed and the whole process is repeated for the needed number of iterations. While in the case of physical slicing the minimal slice thickness, and thus the depth resolution, is limited; virtual slicing is capable of extending it towards nanometer range and hence high-resolution isotropic datasets can be generated. To allow automatic data acquisition the whole workflow was integrated into a hardware and software solution that combines an SEM, an in-situ microtome and a reconstruction software. Increased ease of use is further facilitated through newly developed advanced auto-functions for electron column alignment.

 

References

[1] B. Titze & W. Denk, Journal of Microscopy, vol. 250(2), pp. 101–110 (2013).

[2] F. Boughorbel et al., SEM Imaging Method, Patent US 8,232,523 B2, 31st July 2012.


Type of presentation: Poster

IT-4-P-2793 Monte Carlo simulations of electron trajectories for the study of betavoltaic battery configurations

Napchan E.1
1DLM Enterprises, London NW6 1QH, UK
eli@napchan.com

Battery development goals are to produce small, light, safe, high power and very long lasting batteries. Betavoltaics batteries use semiconductors to convert beta particles (electrons) emitted from a radioactive source, much like photovoltaic panels convert sunlight to electricity. For betavoltaic devices the source can be within the devices themselves, while the radiant sun energy comes from outside the photovoltaic devices. A further difference is that betavoltaic cells can be stacked up.
The simplest structure for a betavoltaic battery consists of the beta layer on top of a pn junction producing electron-hole pairs, which are collected on both sides of the junction. Beta emission in the layer is isotropic within the layer, with randomization of the emission location and the emission characteristics. Each electron emission is isotropic in a sphere, calculated using direction cosines from the random localized emission point.
The Monte Carlo simulation program used is called MC–SET and deals with deposited beam energy calculations and with multi-layers. The simulation tracks each electron in its trajectory inside the specimen, and at each step calculates the energy lost by the electron. The energy deposited from all the electrons in the simulation is stored in a 3-D energy matrix. Other parts of the electrons energy, such as backscattered, transmitted and out to the device electrons are also recorded during the simulation.
The purpose of this investigation is to describe a methodology for simulating beta voltaic batteries, with different geometric configurations. The relationship between the nuclear radiation emission and the energy obtainable is evaluated.
Figure 1 presents the electron depth dose for a bulk Ni specimen, with a normal beam direction. The two selected energies correspond to the average beta emission energy and the maximum beta energy for the Ni-63 isotope. For the high value absorption in the Ni layer occurs at depths of up to about 10 um, while the curve for 17 keV indicates that all the average beta particle energy is absorbed within 1 um of Ni-63. Figure 1 inset shows the depth dose for a layered structure of Si-Ni-63-Si, for 2 um Ni-63 layer. This curve shows the relative amounts of energy deposited in the Ni-63 and the Si layers, the latter being the effective maximum energy available for conversion.
Figure 2 gives the energy deposited in one Si layer, for increasing values of Ni-63 thickness. The left hand curve corresponds to same activity for all layers, i.e. same number of beta emissions, while the right hand curve corresponds to all layers having the same specific activity (beta emissions per gram), corresponding to a typical Ni-63 isotope specific activity of 15 Ci/gr.


Fig. 1: Ni electron depth dose for 2 energies based on Ni-63 emission data and (inset) electron depth dose for Si-Ni-63-Si device

Fig. 2: Relative amount of beta energy emission from Ni-63 layer deposited in Si layer

Type of presentation: Poster

IT-4-P-2865 Estimation of the resolution in 2D Wet-STEM and Wet-STEM tomography by Monte Carlo simulations

Xiao J.1, Perret A.1, Foray G.1, Masenelli-Varlot K.1
1Université de Lyon, INSA-Lyon, Villeurbanne, France
juan.xiao@insa-lyon.fr

The microstructural characterization of water-containing materials in conditions closer to their native state is possible through Environmental scanning electron microscopy (ESEM) experiments. Among the possible ESEM imaging modes, Wet-STEM permits to observe nano-objects in suspension in a liquid with a nanometer resolution [1]. This technique is based on STEM (Scanning Transmission Electron Microscopy) configuration in ESEM. In parallel, a device has been developed for the characterization of the 3D structure of non-conductive and low-contrast materials, and it gives a compromise between the resolution level of a few tens of nm and the large tomogram size due to the large thickness of transparency [2]. Very recently, the implementation of a Peltier stage in the tomographic sample holder has enabled the acquisition of image series in wet samples (wet-STEM tomography) [3].
During Wet-STEM experiments, the contrast is influenced by water thickness and the particle size and composition. Furthermore, the thickness of water varies with the tilt angles, which can lead to contrast inversions. When performing Wet-STEM tomography, contrast inversions have to be avoided when tilting the sample since they may lead to reconstruction artifacts.
In the first part of this study, Monte Carlo simulations will be used to calculate the contrast which can be obtained when observing nanoparticles in suspension in water. We will present how the contrast is affected by the position of a Carbon particle, and its dimension compared to the thickness of the water film (see Figure 1). Then, the contrasts in an experimental Wet-STEM image (see Figure 2) and those calculated from Monte Carlo simulations will be compared.
In the second part, the Monte Carlo simulations will be used to define the best suited sample geometry for Wet-STEM tomography experiments. In particular, the conditions to avoid contrast inversion will be defined, and the resolution will be discussed in function of the nanoparticle composition.

[1] A. Bogner et al., Ultramicroscopy, 104 (2005), 290-301.
[2] Russias J, J. Am. Ceram. Soc., 91,(2008), 2337-2342. P. Jornsanoh et al., Ultramicroscopy, 111 (2011), 1247-1254.
[3] K. Masenelli-Varlot et al., Microscopy and Microanalysis, 2014. doi:10.1017/S1431927614000105


The authors acknowledge the Centre Lyonnais de Microscopie (CLYM) for the access to the microscope, the CSC and Institut Universitaire de France for financial support.

Fig. 1: Numbers of collected electrons for several Carbon particles with different thicknesses of water

Fig. 2: Contrast variation for several Carbon particles with different thicknesses of water

Fig. 3: Experimental Wet-STEM image of a SBA latex suspension – scale bar 500 nm

Type of presentation: Poster

IT-4-P-3005 Microstructural characterization of metallic materials using advanced SEM techniques

Piňos J.1, Konvalina I.1, Kasl J.2, Jandová D.2, Mikmeková Š.1
11. Institute of Scientific Instruments of the ASCR, v.v.i., Czech Republic, 22. VZU Plzen- research and Testing Institute Plzen, Czech Republic
pinos@isibrno.cz

The development of advanced materials is inseparably connected with detailed knowledge of the relationship between microstructure and mechanical properties. Traditional high-voltage scanning electron microscopy (SEM) is one of the most commonly used techniques for microstructure analysis, though it may be insufficient particularly for the characterization of advanced materials exhibiting a complex microstructure.

The benefits of using slow electrons have been described in several articles [e.g. 1,2]. Experiments have been performed with a XHR SEM Magellan 400L (FEI Company) equipped with two detectors for secondary electrons (SE), an Everhart Thornley detector and an in-lens TLD detector, and solid-state BSE detector (CBS) located below the pole piece. This microscope can also be operated in the beam deceleration (BD) mode [3]. The field of the BD not only decelerates the primary electrons, but also accelerates the emitted (signal) electrons towards the detector. Furthermore, high-angle backscattered electrons (BSE) are also collimated towards the optical axis and are detected. These electrons carry, first and foremost, crystal orientation contrast. SE and low-angle BSE can be detected by the TLD detector located inside the objective lens. Angle-resolved detection of BSE is performed using a CBS detector divided into four concentric segments.

Fig. 1 shows the dependence of material contrast between BN precipitates, Laves phase and matrix on the landing energy and increasing of contrast between differently oriented areas in advanced creep resistant steel COST CB2. The presence of BN precipitates and Laves phase has been verified by EDX analysis. The prospect of angular separation of BSE is shown in Fig. 2. The trajectories of signal electrons were simulated in EOD software [4]. It is clearly visible that the low-angle BSE detected by means of the segments closest to the optical axis (S1 and S2) provide information about the chemical composition of the specimen. Segment S3 offers high crystallographic contrast, and material contrast between Laves phase and matrix is entirely suppressed. The final segment S4 exhibits topographic contrast due to the detection of electrons emitted under very high angles from the optical axis, which are products of the interaction of the primary electrons with surface irregularities.

[1] L. Raimer: Image formation in low-voltage SEM. SPIE Press (1993)

[2] I. Mullerova and L. Frank: Scanning low-energy electron microscopy. Adv. Imag. Elect. Phys., Vol. 128 (2003)

[3] Product specification. XHR SEM Magellan 400L. FEI Company

[4] J. Zlamal and B. Lencova: Nucl. Instr. Meth. Phys. Res. A, Vol. 645 (2011)


This work was supported by project no. TE0120118 (Competence Center: Electron Microscopy). The author Šárka Mikmeková is sponsored by an FEI Company Scholarship.

Fig. 1: Fig. 1 The same field of view imaged at 4 keV, 1 keV and 0.44 keV landing energy using the DB mode (specimen bias – 4 kV in all cases), together with corresponding EDX maps of N and Mo distribution.

Fig. 2: Fig. 2 Micrographs obtained at 440 eV landing energy (specimen bias – 4 kV) by means of separated parts of the CBS detector, together with information about detected angles by each segment.

Type of presentation: Poster

IT-4-P-3064 An approach to study antigenotoxicity assay in plants using Confocal Laser Scannig Microscopy and Scanning Electron Microscopy.

Walia A.1, Sharma M.2, Kumar K.1, Bhardwaj R.2, Thukral A. K.2
1Centre for Emerging Life Sciences, Guru Nanak Dev University, Amritsar, Punjab, India, 2Department of Botanical & Environmental Sciences, Guru Nanak Dev University, Amritsar, Punjab, India
adwalia@gmail.com

A new and rapid procedure has been followed using Confocal Laser Scannig Microscopy and Scanning Electron Microscopy for use in the determination of genotoxicity and antigenotoxicity of compounds in plants.

In the present study the above two techniques were used to analyse the genotoxicity and antigenotoxicity effects of compounds on plant system. Certain fluorescent dyes are more reliable indicators of cell viability than the commonly used colored dyes. DNA intercalating dyes like propidium iodide are known to pass only through the membranes of dead or dying cells. Staining with propidium iodide (PI) can be used for the determination of non viable cells. In this study we have evaluated the genotoxic effect of Cr(VI) at different concentrations along with ascorbic acid as a reducing agent in the plant roots. The results showed that the metal ions have a significant effect on the viability of root cells in a dose dependent manner. Also the reducing agent has its effect on reversing the negative effect of these metal ions.

The metal ions are not only genotoxic to plants but they also affect their root growth. To study the pattern of root growth using the same compounds we have scanned the roots of these plants using Scanning electron microscope. The results have shown significant changes in the features of the root tips in different binary combinations of Cr(VI) and ascorbic acid. The study suggests that these techniques can be effectively used for the study of physiological toxicity and antigenotoxicity assays in plants.


We are thankful to University Grants Commission for providing financial assistance to Dr. A.K. Thukral to conduct this work .

Type of presentation: Poster

IT-4-P-3090 Low-voltage STEM tomography: an alternative for soft polymers and hydrated samples

Masenelli-Varlot K.1, Roiban L.1, Malchère A.1, Dhungana D. S.1, Xiao J.1, Jomaa M. H.1, 2, Ferreira J.1, Cavaillé J. Y.1, Seveyrat L.2, Lebrun L.2
1Université de Lyon, INSA-Lyon, CNRS, MATEIS, 7 avenue J. Capelle, 69621 Villeurbanne cedex, France., 2Université de Lyon, INSA-Lyon, LGEF, 8 rue de la physique, 69621 Villeurbanne cedex, France.
Karine.Masenelli-Varlot@insa-lyon.fr

Tomography has become a key characterization tool in materials science as well as in biology. The principle of tomography is based on the acquisition of a series of projections images at different tilt angles, computation of the volume using dedicated algorithms and data segmentation and three-dimensional (3D) quantification. Several tomography techniques are available, using different types of radiations, depending on the observation scale. X-rays are currently used for the 0.5 µm–1 mm resolution level and the three-dimensional characterization of nanoscaled structures requires transmission electron microscopy (TEM) tomography or an atom-probe approach. At the mesoscopic scale, corresponding to a resolution level between 10 nm and 500 nm, Scanning Electron Microscopy (SEM)-based techniques – such as Focused ion Beam (FIB) or serial block face SEM – use a slice-and-view method to directly obtain slices of the materials volume.
Moreover, in Environmental SEM (ESEM), the presence of the gaseous environment and the control of the sample temperature have also permitted the imaging of nanoparticles in liquid with a nanometer resolution, through STEM-in-SEM observations [1]. Its main advantage lays in the fact that water condensation or evaporation can be finely tuned by varying the environmental pressure, which enables in situ hydration / dehydration experiments.

In the first part of this presentation, we will briefly present an alternative tomography technique for the 3D characterization of materials at the mesoscopic scale. This method, called low-voltage STEM tomography, consists in performing tilted tomography in a SEM (in the transmission mode, the so-called STEM-in-SEM mode), see Figure 1 [2]. The potentialities of low-voltage STEM tomography will be compared to that of other 3D techniques through the study of polyurethane films containing two different kinds of carbon nanotubes (see Figure 2).
In the second part of this presentation, we will present the possibility of observing the 3D structure of hydrated materials [3]. In particular, we will discuss the role of different experimental parameters such as the temperature and the electron dose received by the sample. Two examples will be used: a porous material and a latex suspension. Monte Carlo simulations will also be used to estimate the resolution which can be expected in both cases.

[1] A. Bogner et al., Ultramicroscopy 104 (2005), 290-301.
[2] P. Jornsanoh et al., Ultramicroscopy 111 (2011), 1247-1254.
[3] K. Masenelli-Varlot et al., Microscopy and Microanalysis, http://dx.doi.org/ 10.1017/S1431927614000105


The authors acknowledge the CLYM (Centre Lyonnais de Microscopie) for the access to the ESEM XL30FEG microscope, the Agence Nationale de la Recherche and the Institut Universitaire de France for financial support.

Fig. 1: Device for low-voltage STEM tomography, composed of a) and b) piezoelectric elements; c) sample holder and d) STEM detector. The dashed lines represent the position of the Peltier stage.

Fig. 2: Low-voltage STEM tomography on polyurethane thin films containing 2 vol.% of carbon nanotubes (CNT): orthogonal slices extracted from the volume. a) grafted CNTs and b) ungrafted CNTs.

Type of presentation: Poster

IT-4-P-3106 SEM observation of several biological samples using a hydrophilic asymmetrical tetraammonium-type room temperature ionic liquid as a visualizing agent

ABE S.1, KAWAI K.2, YOSHIDA Y.1
1Graduate School of Dental Medicine, Hokkaido University, Sapporo, Japan, 2Miyoshi Oil & Fat Co., Ltd., Tokyo, Japan
sabe@den.hokudai.ac.jp

A room temperature ionic liquid (RTIL) is an organic salt that is liquid at room temperature and has specific physical properties such as noncombustibility, no vapor pressure, high heat resistance, and high ionic conductivity. These unique properties have led many researchers to study the application of ionic liquids in various fields including electronics and chemistry. Kuwabata et al reported that they had succeeded in using RTILs for electroconductive-pretreatment of some samples for scanning electron microscopy (SEM). Because RTILs have electrical conductivity and very low vapor pressure, they can maintain a liquid state even in vacuum such as in an SEM sample chamber. Thus, they can act as visualizing agent for SEM observation. Some types of RTILs, such as imidazolium salts, pyrimidinium salts and ammonium-type salts, have been investigated for the electroconductive pretreatment.

To apply this technique for wet biological samples, we used a novel asymmetrical tetraammonium-type RTIL (HILEM IL1000, Hitachi High-Technologies Corp., Tokyo, Japan). It has chemical structure similar to a choline, which is a bioactive compound. Its properties such as high fluidity, hydrophilicity and biocompatibility can allow using as the agent for SEM observation of biological samples. To elucidate usefullness of RTIL pretreatment, we investigated the conductivity pretreatment for SEM observation of the novel tetraammonium-type RTIL (IL1000). By immersion in an IL1000 solution, clear SEM images of several types of biological samples were successfully observed. We also succeeded in visualization of some bio samples, such as protozoans, red blood cells and bacteria, using IL1000 without dilution. In particular, the size of red blood cells pretreated with IL1000 was in good agreement with that of optical microscopic (OM) observation. When they were treated with traditional method, the obtained SEM images were shrunken compared with those in OM observation. Thus, these results suggested that the tetraammonium-type RTIL used in t his study (IL1000) was suitable for visualizing of biological samples for SEM observation as a "living" morphology. In addition, treatment without the need for dilution can obviate the need for adjusting the RTIL concentration and provide for a rapid and easy conductivity treatment for wet biological samples.


Type of presentation: Poster

IT-4-P-3134 Topotactic transformations of Goethite to Hematite during low metamorphic conditions

Souza D. S.1, Lagoeiro L. E.1, Barbosa P. F.2
1Federal University of Ouro Preto,Departament of Geology,Minas Gerais, Brazil, 2Federal University of Minas Gerais, Microscopy Center,Minas Gerais,Brazil
dansilvasouza@gmail.com

Crystallographic similarities between hematite and goethite allow us to model some topotactic transformation between these two minerals. For natural occurrence of goethite and hematite, such process is scarcely known. The aim of this contribution is to investigate the transformation that occurs in response to a change in deformation and metamorphic conditions of iron-formation rocks deformed at a very low temperature. We applied the EBSD technique to investigate the transformation between ferric oxides and oxyhydroxides combined with the observation of microstructures related to that process.
The samples came from iron-formation rocks in southeast of Brazil, in a region called Iron Quadrangle. Their mineralogy consists basically of quartz, iron oxides and oxyhydroxides. On the optical microscope magnetite is almost completely oxidized to hematite. Several magnetite core grains are filled with goethite and rimmed by hematite grains. The inner goethite and surrounding hematite grains occur in aggregates of irregularly shaped grains of varied sizes.
The EBSD analyzes of the clasts show a close relationship between different crystallographic axes of hematite and goethite crystals. The poles to the basal planes of hematite {001} match those of goethite crystals {001}.
Hematite and goethite, although belonging to different space group symmetries, have similar close packing structures. The structures of hematite and goethite can be described as a slightly distorted hexagonally close-packed of anions (O2- and OH-) stacked along their [c] axes. In these conditions, atom displacements are reduced, so that clear vectorial relations can be established between crystal parameters of the two structures. It is known that the transformation to goethite does not modify significantly the layers of anions in the structure of hematite. Therefore, the expected crystallographic orientation relationship can be described between these two phases.
We proposed that the transformation described in the studied iron formation rocks was performed in two different stages. Initially the original magnetite crystals were hydrated and transformed by oxidation into goethite. This might have been caused by a percolation of low temperature aqueous fluids in the early stages of the deformation. Subsequently, as the deformation proceeds and the temperature increases with the progressive metamorphism of the iron-formations, the newly formed goethite crystals dehydrate and transform into hematite. This can be described as a topotactic transformation because both hematite and goethite show coincidence of orientation in planes {0001} and {001}, and directions <a> and [010], respectively.


Fig. 1: Figure 1 –porphyroclast of magnetite completely transformed to goethite (dark gray) goethite in the magnetite crystal edges is transformed to hematite (light gray); matrix of elongated crystals of goethite (dark gray) occurring as aggregates with quartz ribbons.

Fig. 2: Figure 2- Pole figures for Goethite (a) and Hematite (b). Lower hemisphere, equal area projection. Note a coincidence between the poles to the basal planes of Goethite and Hematite, {001} and {0001}, respectively. The Y0 direction of the microscope system corresponds to the Z0 direction of the sample reference system.

Type of presentation: Poster

IT-4-P-3145 Application of EBSD technique to analyze the microstructure and texture in Peridotites from Archipelago of São Pedro and São Paulo

Pinto S. O.1, Lagoeiro L. E.1, Barbosa P. F.2, Simões L. A.3
1Federal University of Ouro Preto, 2Federal University of Minas Gerais, 3State University of Rio Claro
suellen_olivia@yahoo.com.br

The Archipelago of São Pedro and São Paulo consists in a set of island where rocks from the Upper Mantle outcrops above the sea level [1]. They are a rare example in the world with such this feature. Samples from their rocks were collected for microstructural and crystallographic texture analysis with the aim to get insight into the mechanisms involved in the formation of these rocks, as well as, to infer the deformation mechanisms that developed the observed structures. To achieve that, we use a combination of optical microscopy, to see the whole picture of the microstructure alongside with the powerful of the EBSD analysis. The rocks are Ultramylonite of Peridotite composition. Porphiroclasts of Olivine are deformed by dislocation creep and show subgrains, sweeping undolose extinction, and tails of recrystallized grains of delta type indicate a sinistral sense of shear [2]. The new recrystallized grains adjacent to the clast show crystallographic preferred orientation (CPO) compatible with recrystallization mechanisms of subgrain rotation with some grain boundary migration. In contrast, moving towards the matrix the Olivine grains are much smaller than those close to the clast and there is a weak to random crystallographic texture. A mechanism involving grain boundary sliding assisted by diffusional creep is proposed for the accommodation of the deformation in the matrix. The main challenge is that, it is also not completely ruled out some reaction between minerals in the matrix. Since some grains do not match any minerals loaded in the EBSD acquisition software (CHANNEL 5, in Flamenco mode), and if the nonindexing problem is a matter of phase reaction or the resolution SEM used in the analysis.


References:
[1] www.mar.mil.br/secirm/publicacao/arquipe.pdf
[2] Ron H. Vernon, A Practical Guide to Rock Microstructure, Cambridge University Press, Oct 7, 2004


Fig. 1: Optical image from a characteristic clast from the ASPSP.

Type of presentation: Poster

IT-4-P-3156 Microflow and Thermal Control System Design for Wet Cell in the Scanning Electron Microscope

Lee H. H.1, Lee C. Y.1, Chiang C. L.1, Tsai K. C.1, Lin W. T.1, Wang H. W.1, Fang J. M.2, Huang T. W.3, Liu S. Y.3, Tsai C. Y.3, Chen F. R.3
1Center for Measurement Standards, Industrial Technology Research Institute, Taiwan R.O.C., 2Taiwan Electron Microscope Instrument Corporation, Taiwan R.O.C., 3Engineering and System Science Department, National Tsing Hua University, Taiwan R.O.C.
kylelee@itri.org.tw

Scanning Electron Microscope has been widely used in different scientific areas such as material analysis, biology and life science. SEM integrated with a wet cell was proposed in recent years to satisfy the needs from live cell imaging. In the SEM, it requires a vacuum chamber to allow the operation of the electron beam and to minimize the scattering from other sources. To extend the ability of live cell observation in the SEM, the Si3N4 thin film supported by silicon microchip was developed in order to cultivate biological materials in the wet cell. Therefore, the wet cell can be used to visualize live tissues in fully hydrated conditions and to maintain the culture environment. It would be particularly valuable when applying to the analysis of lipid membranes in cells as they are difficult to preserve during dehydration and washing steps. The processes of sample preparation can be more efficient by using the wet cell. Furthermore, with the continuous flow control and real-time monitoring, the long-term operation can be achieved and also expand the applications of the wet cell.

In Fig. 1, it shows the meshes of simulation model in the wet cell. The pressure is critical because the flowrate needs to be well maintained in order not to break the thin film. To increase the flowrate and reduce the costs of the fluid mechanics, the geometry and flow conditions were optimized based on the skills of Design Of Experiments. The resulting pressure on the 50 nm thin film was simulated by CFD software (ANSYS Fluent v14) and the results are shown in Fig. 2. It demonstrates that the inlet flowrate still could be raised and the size of the wet cell can be reduced. With regard to the thermal control system, a preheating system of buffer liquid was developed by using a thermal control module. In addition to preheating system for the buffer liquid, an embedded micro thermal control element was also designed inside the wet cell with the capability of fine-tuning so as to achieve accurate and rapid temperature control. Furthermore, several design parameters including noiseless, non-vibration and long working life were also considered for the thermal control system. After preliminary experiments, the results are shown in Fig. 3. and the heating rate inside the wet cell can be achieved to 3.7 ℃ per minute. Eventually, the microflow and thermal control systems were integrated and the system architecture is shown as Fig. 4. With the microflow and thermal control modules integrated in the SEM, the flowrate and fluid temperature can be adjusted by users and the flow conditions including temperature, pressure and even the fluid properties can be simultaneously monitored as well.


Thanks for the technical assistance and suggestions from R&D team of Taiwan Electron Microscope Instrument Corporation (TEMIC) during system integration.

Fig. 1: Meshes of the liquid volume inside the wet cell

Fig. 2: The pressure on the thin film inside the wet cell

Fig. 3: The temperature variation in the wet cell during heating process

Fig. 4: System architecture of microflow and thermal control system

Type of presentation: Poster

IT-4-P-3166 Combined EDS and WDS analysis of thin specimens with high spatial and energy resolution in the scanning electron microscope

Mitsche S.1, Poelt P.1
1Graz University of Technology, Institute for Electron Microscopy and Nanoanalysis, Steyrergasse 17, 8010 Graz, Austria
stefan.mitsche@felmi-zfe.at

X-ray analysis of bulk samples by energy-dispersive x-ray spectrometry (EDS) in a scanning electron microscope (SEM) is widely used to gain chemical information of materials. The combination of EDS and bulk samples is limited by the low energy resolution of EDS, the associated high detection limit and the low spatial resolution due to the interaction volume of the electrons which is correlated to the beam energy used. A wavelength-dispersive spectrometer (WDS) can increase the energy resolution dramatically. An improvement of the spatial resolution can be obtained by use of very thin samples (less than 50 nm). The combination of EDS, WDS and thin specimens improves the spatial and energy resolution and detection limits and keeps the analysis time down.

All investigations were performed on a Zeiss Ultra 55 equipped with a 10 mm2 Si(Li)–EDS detector and a parallel beam WDS detector with a multi capillary optics from EDAX. Thin specimens were prepared by ion milling with a FIB.

In order to minimize shadowing effects a special sample holder was designed (see Fig. 1). With this sample holder a linescan across a FIB-lamella from a semiconductor device (see Figure 2a) was performed by EDS and WDS. The resulting intensities of the Ti-K peak recorded with both EDS (grey) and WDS (dark grey) are shown in Figure 2b (dwell time 1000ms for EDS and 2000ms for WDS). This Figure demonstrates that the x-ray intensities for both EDS and WDS are sufficiently high for the x-ray analysis of thin specimens. Figures 2c and 2d demonstrate the benefit of the high energy resolution of WDS. Whereas the signal of the W-M line recorded with EDS runs quite similar to that of the Si-K line, WDS clearly proves that in the region with high Si content no W is present.

A FIB-lamella of a circuit board with a 40 nm Au layer followed by a 14 nm Pd layer was analyzed by EDS to validate the improvement of the spatial resolution of x-ray analysis by investigating thin instead of bulk specimens (see Figure 3a). The linescan across these two layers is plotted in Figure 3b. This Figure proves that both layers can be detected separately by x-ray analysis. As a Lorentz distribution was the best fit to the scan of the Pd intensity its half width was used as a measure for the thickness of the Pd layer. It gave a value of 15 nm which is close to the value obtained by the electron image


I would like to thank the FFG for the financial support (Project number: 825165)

Fig. 1: Modified specimen holder, a) side view b) top view c) bottom view of the optimized holder; 1: optimized holder, 2: original Zeiss STEM holder, 3: FIB-lamella

Fig. 2: X-ray linescans across a semiconductor structure, a) BSE image with linescan marked (image width: 800 nm) , b) EDS (grey) and WDS (dark grey) signal of the Ti-K line c) EDS and WDS signal of the Si-K line, d) EDS and WDS signal of the W-M line; abscissa: distance in microns, ordinate: number of x-rays.

Fig. 3: X-ray linescan across a circuit board, a) BSE image with linescan marked (image width: 380 nm), b) EDS signal of Au-L line (dark grey) and Pd-L line (dashed line grey), light grey the Lorentz fit to the Pd scan; abscissa: distance in nanometre, ordinate: number of x-rays.

Type of presentation: Poster

IT-4-P-3232 PERFORMANCE OF YAG:Ce SCINTILLATORS FOR LOW-ENERGY ELECTRON DETECTORS IN S(T)EM

Lalinsky O.1, Bok J.1, Schauer P.1, Frank L.1
1Institute of Scientific Instruments of the ASCR, Brno, Czech Republic
xodr@isibrno.cz

Cerium activated single crystals of yttrium aluminium garnet (YAG:Ce) Y3-xCexAl5O12 are widely used as scintillators in electron detectors for S(T)EM [1]. Nowadays, it is sometimes necessary to detect low-energy electrons without post-acceleration. In such cases, extremely sensitive detectors are required that are able to detect even electrons with energies of only hundreds of eV while avoiding charging of the scintillator surface. However, commonly used scintillators strongly lose their light yield with the decrease of the incident electron energy [2]. Moreover, a thinner conductive layer on the scintillator surface has to be used to allow low-energy electrons to pass through. Possible charging of the surface negatively affects its cathodoluminescence (CL) light yield. The low-energy electron excitation takes place closer to the scintillator surface where damage can be expected owing to its preparation, which also reduces the CL light yield. The aim was to study the influence of the scintillator and its conductive layer on the low-energy electron detection efficiency.
In general, the following demands are made of the conductive layer: it should have the highest possible optical reflectivity, conductivity and electron transparency. These demands are mostly met for metals with a low atomic number. However, if the layer is very thin, it can form “islands”, i.e. a non-continuous layer of drastically decreased conductivity. We decided to apply scandium as a possible option. The YAG:Ce single crystals were studied using both Monte Carlo simulation and CL measurement. The MC method (Fig. 1) used Mott cross-sections and the Bethe slowing-down approximation. Using the CL apparatus (Fig. 2) [3], incident electron energy can be changed from 0.8 to 10 keV. The detection dynamic range spans 5 orders of magnitude. The experimental results are shown in Fig. 3. The significant decrease of efficiency at lower energies may be caused by the layer which doesn’t allow more electrons to pass through, by the YAG:Ce single crystal that has a lower light yield near the surface, and finally, if the layer isn’t conductive enough, it can be charged and retard incident electrons. Even so, we have found that the YAG:Ce scintillator with a 3 nm scandium layer is applicable for the detection of electrons having an energy as low as 800 eV.

References

[1] P. Schauer, Nucl. Instrum. Methods Phys. Res. B 21 (2011) 2572–2577.
[2] G. F. J. Garlick, Brit. J. Appl. Phys. 13 (1962) 541–547.
[3] J. Bok, P. Schauer, Rev. Sci. Instrum. 82 (2011) 113109.


Thanks are due to CRYTUR, Ltd., for the supply of scintillators, to J. Sobota for the preparation of scandium layers, to the Technology Agency of the Czech Republic (TE01020118), to the European Commission and to the Ministry of Education, Youth and Sports of the Czech Republic (CZ.1.07/2.3.00/20.0103).

Fig. 1: Interaction volumes in the YAG:Ce single crystal simulated by the Monte Carlo method. Simulated without any conductive surface layer.

Fig. 2: Experimental equipment used for cathodoluminescence (CL) property measurements.

Fig. 3: The cathodoluminescence light yield of the YAG:Ce single crystal scintillator coated with a scandium layer of a thickness of 3 and 5 nm, respectively, as a function of the incident electron energy.

Type of presentation: Poster

IT-4-P-3358 Characterization of β-phase in Al-Mg-Si alloys by SLEEM and STLEEM techniques

Ligas A.1, Hida S.2, Matsuda K.3, Mikmeková Š.1
1Institute of Scientific Instruments of the ASCR, v.v.i., Czech Republic, 2Graduate School for Science & Engineering for Education, University of Toyama, Japan, 3Graduate School for Science & Engineering for Research, University of Toyama, Japan
ales.ligas@isibrno.cz

Knowledge of the distribution and morphology of the Mg2Si precipitates (i.e. β-phase) in Al-Mg-Si alloys are very important for many practical reasons [1,2] and the scanning electron microscopy (SEM) technique is widely used for their visualization. Unfortunately, in the standard SEM images these precipitates are barely visible and finding them can be very difficult. Using the cathode lens (CL) mode in the SEM (so called SLEEM [3]) these difficulties have been overcome and a very high contrast between the hexagonal-shaped β-phase and the matrix has been obtained. Moreover, it has been found that the SLEEM images offer the possibility to distinguish between the hexagonal-shaped and the conventional β-phase based on their different brightness, not only on their shape, which can be in some cases difficult or even impossible. Mg2Si precipitates have been also characterized by means of the scanning transmission low energy electron microscopy (STLEEM [4]) method based on the using of a STEM detector in the SEM operated in the CL mode.

[1] Laughlin DE, Miao WF. Automotive alloys II. TMS Warrendale PA. 1998; 63-79.
[2] Edwards GA, Stiller K, Dunlop GL, Couper MJ. The precipitation sequence in Al-Mg-Si alloys. Acta Materialia 1998; 46: 3893-3904
[3] Mullerova I, Frank L. Scanning low energy electron microscopy. Advances in imaging and electron physics 2003; 128: 304-443
[4] Mullerova I, Hovorka M, Hanzlikova R, Frank L. Very low energy scanning electron microscopy of free-standing ultrathin films. Materials Transactions 2010; 51: 265-270


The financial support of the project no. TE0120118 (Competence Centre: Electron Microscopy) from the Technology Agency of the Czech Republic is greatly acknowledged. The author (ŠM) is sponsored by FEI Company scholarship.

Fig. 1: Comparison between the SEM secondary electron image (a), backscattered electron image (b) obtained at 10keV and the SLEEM image (c) obtained at 2 keV landing energy of the hexagonal-shaped β phase.

Fig. 2: The same point of view imaged at 10 keV primary energy in the standard mode: by means of SE electrons (a) and BSE electrons (b) and in the CL mode: at 5 keV (c) and at 0.63 keV lading energy (d).

Fig. 3: The same point of view imaged at 10 keV primary energy in the standard mode: by means of SE electrons (a) and BSE electrons (b) and in the CL mode: at 5 keV (c) and at 0.63 keV lading energy (d).

Type of presentation: Poster

IT-4-P-3491 Application of Ultra Low Voltage Secondary and Backscatter Imaging in FE-SEM for Imaging of Nanomaterials

Erdman N.1, Robertson V.1, Shibata M.1
1JEOL USA Inc
erdman@jeol.com

Significant advances in electron optics and detectors in field emission scanning electron microscope (FE-SEM) in the last decade have allowed the researchers to observe a variety of materials and biological specimens with ultra-high resolution and exceptional surface detail. In particular, low voltage imaging has been successfully employed as a key technique for charge control and reduction. Enhancements in electron column optics towards smaller chromatic and spherical aberration coefficients, with improved ability to deal with charging specimens via precise control of the landing energy of impact electrons and electron signal detection through in-column signal filtering or signal collection angle control have opened new avenues for specimen observation [1]. These new design improvements have significantly advanced the ability to image insulating specimens with previously unattainable nanometer scale resolution [2] at landing voltages as low as 10V (Fig. 1). In this paper we will discuss common approaches and challenges associated with ultra-low voltage imaging. The instrument employed for these studies is JSM-7800F ultra-high resolution FE-SEM that features a hybrid lens design and the ability to bias the specimen stage thus decelerating the primary beam (Gentle Beam). When beam deceleration is employed the accelerating voltage which along with lens aberrations determines the minimum probe size and thus the resolution limit is retarded by a negatively charged bias to a lower landing energy. The landing voltage can be varied to achieve the necessary charge balance as well as high resolution performance at ultra-low voltages. Beam deceleration also serves as a form of aberration correction [3]. The use of Gentle Beam function preserves all the advantages of high kV imaging (gun brightness, small probe size) with added advantages of reduced charging, reduced specimen contamination and improved surface detail. Moreover, through-the-lens detection system features an ability to precisely filter the detected signal, providing the user with an additional degree of control during the imaging. We will demonstrate our experiences with imaging a variety of specimens, such as zeolites, biological nanostructures, oxides, nano-structured metals and more. The advantages of low kV imaging for such techniques as cathodoluminescense imaging and voltage contrast will be highlighted. Additional methods for charge balance, such as reduced probe current and adjustment of scan speed will also be discussed. References: [1] D.C. Bell and N. Erdman. Low Voltage Electron Microscopy: Principles and Applications (2012) [2] S. Asahina et al., Microscopy and Analysis, (2012) p.S12. [3] L. Frank and I. Müllerová, Ultramicroscopy, 106 (2005) p. 28


Fig. 1: Anopore membrane filter imaged uncoated at 10V. Pore walls (14 and 21 nm) are clearly resolved.

Type of presentation: Poster

IT-4-P-5734 Palynomorphology of Dianthus petraeus (Caryophyllaceae)

Mačukanović-Jocić M.1, Jarić S.2
1Faculty of Agriculture, University of Belgrade, 11080 Belgrade-Zemun, Serbia, 2Department of Ecology, Institute for Biological Research ‘Siniša Stanković’, University of Belgrade, 11060 Belgrade, Serbia
marmajo@agrif.bg.ac.rs

The full text of the abstract is not available. Please contact the presenting author.


Type of presentation: Poster

IT-4-P-5765 E-BEAM CROSSLINKING AND THERMAL DEGRADATION OF HYDROGEL UNDER ELECTRON MICROSCOPE

Wong Y. H.1, Kuo C. H.1, Huang T. W.1, Liu S. Y.1, Hsieh H. Y.23, Chen F. R.1, Tseng F. G.14
1Department of Engineering and System Science, National Tsing Hua University, TAIWAN, 2Department of Mechanical Engineering, National Taiwan University, TAIWAN, 3Institute of NanoEngineering and MicroSystems, National Tsing Hua University, TAIWAN, 4Division of Mechanics, Research Center for Applied Science, Academia Sinica, TAIWAN
brian791213@gmail.com

Introduction

In the past decades, hydrogel has been becoming an important medium for incorporating cells together to from 2D or 3D structures for tissue engineering applications.Electron beam can be used to pattern the resulting hydrogels on silicon or glass surfaces with nanoscale and microscale feature sizes by radiation crosslinking[1]. The water content plays a role to form the hydrogen and hydroxyl radicals which initialize the polymerization reaction. To visualize the dynamic evolution of hydrogel, we use a hermetic micro-device (wet-cell) to preserve the hydrated hydrogel in vacuum system under electron illumination. We has reported the innovation of self-aligned micro wet-cell and demonstrated the TEM examination of hydrated Deinococcusradioduransin our previous work[2]. In this paper, we furthermore introduce an advancing micro wet-cell with miniature heater integrated to achieve a temperature manipulating for the rapid recovering of hydrogel.

Chip fabrication and set up

The multiple-electrode wet-cell device is composed of two silicon chips with complementary structure, as shown in Fig.1.The “cover piece“, a 3mm x 3mm square-shaped device made of 250μm-thick Si wafer, has an observing window which is formed by the bulk micromachining and covered by a silicon nitride membrane; the “electrode piece”, a 3mm x 6mm rectangle-shaped device made of 250μm-thick Si wafer, consists of a similar observing window with additional Ni/Cr heater (200nm/50nm, 20.07 kΩ) as well as the extended metal pad for wire connecting.

Experiment Results

To measure the temperature rising with increasing applied voltage, we use an infrared-thermal microscope to visualize the distribution and variation of temperature. Several different applied voltages and their corresponding temperatures are shown in Fig.2 . We can control our heater temperature ranging from 30 to 70 °C, reversibly. We also used the Hitachi TM-1000 and made a circuit on the side wall of the SEM, so that we can directly applied tunable voltage in the vacuum system. We first applied the e-beam radiation onto the specific region of GelMa. In Fig.3a-b, electron radiation posses higher effect on GelMa as we zoom in the field of view.The cross-linked regions, as shown in Fig.3 c-d, can be approximately characterized as 1,973μm2. The degradation of the hydrogel was observed after heating, and the cross-linked regions has been fully dissolved at a temperature of 50.2°C.Finally ,the detail results will be discussed in detail in the conference.

REFERENCES:

1. Tsu-wei Huang and Fu-Rong Chen,Korean Journal of Microscopy, 38, 41 (2008)
2. Tsu-wei Huang and Fan-Gang Tseng, Proc. MicroTAS’09, Jeju (2009)
3. Hsin-Yi Hsieh and Fan-Gang Tseng, Lab on a Chip, 10.1039/c3lc50884f


This work was supported by National Science Council
(NSC102-2321-B-007-007 and NSC 102-2120-M-007-006-CC1).

Fig. 1: Figure 1.(a)Schematic of two parts of our device assembly (b)Schematic of E-beam and chip observing area

Fig. 2: Figure 2.IR microscope image of our device heatingby applied voltage from 50.2℃ to 70.1℃

Fig. 3: Figure 3.Orange frame is the electrode part . Blue frame is the unchanged part. (a)-(b)electron radiation affected GelMa as we zoom in (red frame region). Heating result is acquired after e-beam treatment. (c)-(d) Due to heating, the contrast of a pointed region is changed which can be observed in SEM but not in OM

Type of presentation: Poster

IT-4-P-5785 Time-Resolved Electron Beam Induced Current (TREBIC) Method for Spatiotemporal Analysis of GaN HFETs Structures.

Šatka A.1, Priesol J.1, Donoval D.1, Uherek F.1
1Institute of Electronics and Photonics, Slovak University of Technology, Ilkovičova 3, 812 19 Bratislava, Slovakia, 2International Laser Centre, Ilkovicova 3 841 04 Bratislava 4 Slovak Republic
alexander.satka@stuba.sk

GaN-based Heterostructure FETs are becoming preferable devices for high-speed and high-power applications in harsh environments. However, they suffer from the significantly deteriorated switching properties whose mechanisms are intensively investigated [1]. One of the possibilities to study spatiotemporal behavior of the device structures is to temporarily inject limited amount of charge using focused electron beam and to investigate the device’s time response. The lifetime in GaAs heterostructures was investigated by TREBIC in [2]. Combined time-resolved OBIC and TREBIC were used to map time response of InGaAsP/InP pin photodiode [3]. Dynamic behavior of HEMTs was investigated using backside infrared OBIC technique [4]. Stroboscopic TREBIC system was applied for analysis of dynamically biased power devices in [5].
In this contribution we report on the developed TREBIC system using sampling time-gated boxcar averaging techniques for spatiotemporal analysisof GaN based HFET structures. A field emission gun SEM equipped with beam-blanking system and multi-contact vacuum feedthrough has been used. Unpackaged HFETs soldered in a sample holder were connected to the feedthrough by 50 ohm coaxial cables. Induced current was detected by I/V converter (Figure 1). Alternatively, e-beam induced conductivity changes were detected, when a voltage on the resistor divider formed by resistor Rd and transistor channel Rch was preamplified by voltage amplifier. TREBIC signal in selected point or in selected area on sample was detected by digital oscilloscope with periodic signal used for beam-blanking. This technique offers excellent time resolution but it suffers from the huge amount of transferred data for each position of the e-beam [3]. Therefore, time-gated boxcar integration technique has been used to map the EBIC signal after selected delay Td from the rising edge of the e-beam (Figure 2). The gate width Tg determines the time resolution. Box car averaging of pulses was set as a compromise between the sensitivity and acquisition time. TREBIC signals measured from InAlN/GaN HFET device are shown in Figure 3, revealing slow response of the device to the charge injected in the G-D region of the transistor. From the series of TREBIC maps taken at various Td and Tg (e.g. in Figure 4) formation of a virtual gate in G-D region and inhomogeneous electric field build-up and recovery at the gate edge has been observed.

1. J.G.Tartarin et al., In: Proc. of the IWS 2013, IEEE, 2013
2. L.J. Balk et al., In: Proc. of the SEMAS, Part I, 447-456 (1975)
3. A. Šatka, J. Kováč, Microel. Eng. 24, 195-201 (1994)
4. D. Pogany et al., Microel. Reliab. 42, 1673-1677 (2002)
5. A. Pugatschow et al. In: Proc. of the 18th ISPSDIC, Naples, 2006, 4pp.


This work has been supported by the Slovak Research and Development Agency (contract No. APVV-0367-11) and by Slovak Grant Agency (project VEGA No. 1/0921/13).

Fig. 1: Schematic drawing of the TREBIC experimental set-up using time-gated boxcar integration and oscilloscope techniques.

Fig. 2: Time dependence of induced current and formation of TREBIC maps.

Fig. 3: Normalized TREBIC curves as a function of gate voltage VGS taken at VDS = 9.6V and e-beam energy Ebeam = 1 keV.

Fig. 4: TREBIC map taken from the top of HFET structure (see inset) immediately after the switching e-beam on (Td = 0 s) using gate width Tg = 3 ms. Pulse width Tp = 50 ms, VDS = 4 V, VGS = -4 V, Ebeam = 5 keV.

Type of presentation: Poster

IT-4-P-5814 Development of "Adaptive SEM" Technology for in situ Genome/Proteome Expression Analysis in Single Cell Level

Kim H.1, Terazono H.1,2, Takei H.1,3, Yasuda K.1,2
1Kanagawa Academy of Science and Technology, Kanagawa, Japan, 2Institute of Biomaterials and Bioengineering, Tokyo Medical and Dental University, Tokyo, Japan, 3Department of Life Sciences, Faculty of Life Sciences, Toyo University, Gunma, Japan
ykp-kim@newkast.or.jp

Obtaining expression information in a cellular system is essential for understanding the mechanisms of living systems. One useful way is in situ measurement of expressed biomarkers in single cell level using a lot labels; however, production and identification of such labels are still challenging. We propose a new sensing technology based on the field emission scanning electron microscopy (FE-SEM), which is a comprehensive development of production and identification of nano-particle (NP) labels for simultaneous in situ measurements of expressed biomarkers in single cell (Fig.1).

For the fabrication of NPs, various sizes of polystyrene spheres were used as templates, and metals were deposited on the spheres by thermal evaporation. By using this method, more than 500 types of NPs were fabricated. Metal shell layers were formed by thermal evaporation; therefore, multi-layered NPs can be fabricated with sequential evaporation. We used double-layered NPs; outer is Au for easy immobilization of biomolecules to use these NPs as labeling probes, and inner layer is various to apply label varieties in FE-SEM observation. In this study, probe DNAs were immobilized onto the outer Au layers (referred to as "NP probe" hereafter), and target DNAs on a substrate were reacted with the NP probes as a model. For its detection, FE-SEM observations were performed to identify numbers and elements of hybridized NP probes on the substrate. Spatial distributions and diameters of NP probes were identified by secondary electron (SE) observation, and elements of NP probes were identified by backscattered electron (BE) observation as the difference of intensities in the BE image caused by the difference of atomic number of inner metal layer (Fig.2). In results, six different elements were simultaneously distinguished by BE observation [1], indicated that targets can be simultaneously labeled and identified with high spatial resolution by the combination of NP probe labeling with BE image analysis. In addition, detection sensitivity of target DNA in this method was femto-molar order [2] (Fig.3), which is 1,000 times higher than that in conventional fluorescent labeling and optical detection, indicated that our method is suitable for the detection of a few biomarkers in single cell. We call it "adaptive SEM" technology (i.e., NP identification is "adaptive" for various targets). These results indicate a possibility for quantitative in situ detection of expressed biomarkers in single cell level by the suggested technology based on NP probe labeling and FE-SEM identification.

References

[1] Kim, H., Negishi, T., Kudo, M., Takei, H., Yasuda, K., J. Electron Microsc. 59 (5), 379-385 (2009)

[2] Kim, H., Kira, A., Yasuda, K., Jpn. J. Appl. Phys. 49 (6), 06GK07, 1-7 (2010)


We thank Ms. M. Murakami and Ms. M. Naganuma for their technical assistance. This work was financially supported by the Japan Prize Foundation, JSPS, and KAST.

Fig. 1: Overview of "adaptive SEM" technology. Firstly, target biomarkers in single cell were simultaneously labeled with NP probes on which probe DNAs were coated, and next, the attached NP probes were identified with SE and BE observations of FE-SEM to identify expressed targets in the cell.

Fig. 2: Discrimination of NP probes. In this example, 120 nm of Au, Ag, and Ni NP probes were simultaneously observed with SE (a) and BE (b) detections using FE-SEM. As shown in (b), difference of elements can be distinguished as the difference of intensities in the BE image. Bars, 500 nm.

Fig. 3: Evaluation of detection sensitivity using DNA chip. (a) Overview of the evaluation. Target DNA on the chip was labeled with NP probe, and the result was observed using FE-SEM. Hybridized NP probes were detected as white dots in FE-SEM images. Bars, 1 μm. (b) Relationship between the number density of observed NP probes and target concentration.

Type of presentation: Poster

IT-4-P-5889 Microflow and Thermal Control System Design for Wet Cell in the Scanning Electron Microscope

Lee H. H.1, Lee C. Y.1, Chiang C. L.1, Tsai K. C.1, Lin W. T.1, Wang H. W.1, Fang J. M.2, Huang T. W.3, Liu S. Y.3, Tsai C. Y.3, Chen F. R.3
1Center for Measurement Standards, Industrial Technology Research Institute, Taiwan R.O.C., 2Taiwan Electron Microscope Instrument Corporation, Taiwan R.O.C., 3Engineering and System Science Department, National Tsing Hua University, Taiwan R.O.C.
KyleLee@itri.org.tw

Scanning Electron Microscope has been widely used in different scientific areas such as material analysis, biology and life science. SEM integrated with a wet cell was proposed in recent years to satisfy the needs from live cell imaging. In the SEM, it requires a vacuum chamber to allow the operation of the electron beam and to minimize the scattering from other sources. To extend the ability of live cell observation in the SEM, the Si3N4 thin film supported by silicon microchip was developed in order to cultivate biological materials in the wet cell. Therefore, the wet cell can be used to visualize live tissues in fully hydrated conditions and to maintain the culture environment. It would be particularly valuable when applying to the analysis of lipid membranes in cells as they are difficult to preserve during dehydration and washing steps. The processes of sample preparation can be more efficient by using the wet cell. Furthermore, with the continuous flow control and real-time monitoring, the long-term operation can be achieved and also expand the applications of the wet cell. In Fig. 1, it shows the meshes of simulation model in the wet cell. The pressure is critical because the flowrate needs to be well maintained in order not to break the thin film. To increase the flowrate and reduce the costs of the fluid mechanics, the geometry and flow conditions were optimized based on the skills of Design Of Experiments. The resulting pressure on the 50 nm thin film was simulated by CFD software (ANSYS Fluent v14) and the results are shown in Fig. 2. It demonstrates that the inlet flowrate still could be raised and the size of the wet cell can be reduced. With regard to the thermal control system, a preheating system of buffer liquid was developed by using a thermal control module. In addition to preheating system for the buffer liquid, an embedded micro thermal control element was also designed inside the wet cell with the capability of fine-tuning so as to achieve accurate and rapid temperature control. Furthermore, several design parameters including noiseless, non-vibration and long working life were also considered for the thermal control system. After preliminary experiments, the results are shown in Fig. 3. and the heating rate inside the wet cell can be achieved to 3.7 ℃ per minute. Eventually, the microflow and thermal control systems were integrated and the system architecture is shown as Fig. 4. With the microflow and thermal control modules integrated in the SEM, the flowrate and fluid temperature can be adjusted by users and the flow conditions including temperature, pressure and even the fluid properties can be simultaneously monitored as well.


Thanks for the technical assistance and suggestions from R&D team of Taiwan Electron Microscope Instrument Corporation (TEMIC) during system integration.

Fig. 1: Meshes of the liquid volume inside the wet cell.

Fig. 2: The pressure on the thin film inside the wet cell.

Fig. 3: The temperature variation in the wet cell duringheating process.

Fig. 4: System architecture of microflow and thermalcontrol system.

Type of presentation: Poster

IT-4-P-5937 New scintillation low-energy BSE detector

Kološová J.1, Beránek J.1, Jiruše J.1, Horodyský P.2
1TESCAN Brno, s.r.o., Brno, Czech Republic , 2CRYTUR, spol. s r.o., Turnov, Czech Republic
jolana.kolosova@tescan.cz

In the electron microscopy research of nanomaterials, biomaterials or semiconductors, low energy electron beam imaging is often necessary. Reducing the primary beam energy decreases the depth of specimen radiation damage, enables clear visualization of non-conductive samples and leads to enhanced specimen surface contrast.
Low accelerating backscattered electron (BSE) imaging with sufficiently high signal to noise level can be done with the new generation of solid state detectors. These detectors have good sensitivity in the low energy region and their speed approaches the speed of scintillation-type detectors. However, in dual beam systems the deposition of sputtered material on the detection surface can lead to deterioration of performance. Further drawback is the sensitivity of solid state detectors to light.
Scintillation detectors are fast and versatile, but their sensitivity drops rapidly in the low energy region thanks to the ‘dead layers’ on the detection surface (e.g. conductive coating), which are impenetrable for slow electrons. CRYTUR in cooperation with TESCAN has developed a new scintillation type BSE detector with special surface treatment, which guarantees enhanced sensitivity in the low energy region.
Detection limit of the new detector is less than 1 keV. It’s high performance in the field of energies under 3 keV makes it ideal for example for BSE imaging of surface details and contrast changing (see Figure 1), high resolution imaging of sensitive biological samples, or artifact free imaging of nonconductive samples (see Figure 2).


Fig. 1: Change of contrast in BSE images of CeO2 ceramics taken at 3 kV (left) and 1 kV (right) accelerating voltages. More surface details are resolved with lower primary beam energy.

Fig. 2: BSE image of Vitrina pellucida shell taken at 3 kV (left) and 2 kV (right) acceleration voltages. Charging artifacts are not visible at 2 kV.

Type of presentation: Poster

IT-4-P-5966 Effect of hydroquinone treatment on OTO en bloc stained biological specimens.

Togo A.1, Higashi R.1, Ohta K.2, Nakamura K.2
1Kurume University School of Medicine EM LAB, 2Kurume University School of Medicine Department of Anatomy
togou_akinobu@med.kurume-u.ac.jp

INTRODUCTION
In recent three dimensional (3D) ultrastructural reconstruction techniques such as serial block face scanning electron microscopy (SBFSEM), TEM-like ultrastructural images of biological specimens are directly obtained from block surface of resin embedded specimens. Since this TEM-like block face images (BFI) is usually obtained using backscattered electrons (BSE) as a material contrast image, specimens are stained strongly by heavy metals prior to embedding into resin. In order to enhance the membrane contrast for BFI, we usually stain specimens by the method of Deerinck (2010). As a recent large volume reconstruction requires very long time to obtain image sets, we need a new staining method which provides much higher contrast that enable to acquire images in a shorter time. Takahashi et al. (1986) have reported that hydroquinone (HQ) treatment during the traditional electro-conductive staining increases specimen conductivity and drastically reduces charge problem for SEM observation. They concluded that HQ treatment might increase the efficient secondary electron (SE) generation. As BFI could be obtained not only by BSE but SE, we examined whether HQ treatment in en bloc staining protocol increased the contrast for BFI using SE in this study.

MATERIALS & METHODS
C57BL/6 mouse liver was used. The animals were deeply anesthetized with diethyl ether and sodium pentobarbital, and fixative of 2% paraformaldehyde and 2.5% glutaraldehyde in 0.1M cacodylate buffer (pH 7.4) was transcardially perfused through the left ventricle with heparin containing saline. After perfusion, liver tissues were removed and cut into small cubes about 1 mm3 in the fixation, and were further fixed in the same fixative for 2h at 4℃. After that, en bloc staining was performed as follows: The specimens were treated with reduced-OTO staining method (1.5% potassium ferrocyanide and 2% OsO4, 1% TCH, 2% OsO4). Subsequently specimens were treated with 1% HQ solution. Some specimens were skipped this step as a control. Then, they were further stained by 4% uranyl acetate and Walton’s lead aspartate solution. After staining, specimens were dehydrated in an ethanol series and were embedded in epoxy resin (EPON812, TAAB). The surface of resin block with specimens was observed by SEM (Quanta 3D FEG, FEI).

RESULTS AND DISCUSSION
The contrast of SE images was drastically increased by HQ treatment, although there is no effect for BSE images. This result suggests that HQ treatment effectively enhances SE generation from specimens. This enhancement may accelerate data acquisition speed for SBFSEM 3D reconstruction.


Type of presentation: Poster

IT-4-P-6025 Two customized low-cost systems for STEM in SEM microscopy

Caciagli A.1,2, Tessarolo F.1,3, Piccoli F.2,4, Caola I.2,4, Nollo G.1,3, Caciagli P.4
11) Department of Industrial Engineering, University of Trento, via delle Regole 101, 38123 Mattarello Trento, Italy , 22) Section of Electron Microscopy, Azienda Provinciale per i Servizi Sanitari di Trento, via Degasperi 79, 38123 Trento, Italy, 33) Healthcare Research and Innovation Program, Bruno Kessler Foundation, Via Sommarive 18, 38123 Povo Trento, Italy , 44) Department of Medicine Laboratory, Azienda Provinciale per i Servizi Sanitari di Trento, via Degasperi 79, 38123 Trento, Italy
federico.piccoli@apss.tn.it

Scanning transmission electron microscopy (STEM) has become an established technique in microscopy, combining the transmission mode with the sample scanning (Bogner 2007). Although STEM can be performed in dedicated facilities, it is possible to implement it on existing SEM microscopes (STEM-in-SEM) in a relatively simple and cheap way (Vanderlinde 2004). Two customized low-cost stages for STEM-in-SEM microscopy are presented with some representative images obtained from material and life science.
The systems were optimized for a XL 30 FEI, but they can be easily adapted to other microscopes with minor modifications. STEM Signal is detected by the back-scattered electron detector (BSED) or the secondary electrons detector (SED). TEM grids fit in both systems. The first system (STEM-A) (Fig. 1a) was adapted from the configuration proposed by Merli and Morandi in 2005. An aluminium table is mounted on four columns screwed in the x-y stage. The grid is placed in a 2.5 mm hole at a working distance of 10 mm. An aluminium ring blocks the grid. The BSED is cemented on the original sample older allowing to adjust the BSED-grid distance by regulating the z stage axis. To obtain both dark-field (DF) or bright-field (BF) images, BSED should be centred respectively on or off the beam axis. The second system (STEM-B) (Fig. 1b) was derived from the configuration proposed by Golla at al. in 1994. The beam is focused on the sample via a graphite tunnel that attenuate SE signal (different tunnel heights allow choosing the optimal SE-TE signals ratio). Transmitted electrons are scattered back to the SED by two gold-sputterd coverglasses.
STEM-A proved better results with high beam voltage (>20 KV), guaranteeing a sufficient transmitted signal. A distance from 30 to 40 mm between sample and detector gave best results. Details down to 20 nm can be easily resolved (Fig. 2a ). BF and DF images provide complementary information. STEM-B with tunnels of different heights allowed various advantages. The shortest tunnel provides a better image quality, even though a SE signal was also present (Fig. 2b). Deeper tunnels allowed to collect pure TE signal. A less sensitivity to spot size was found in respect to the STEM-B (spot size of 2.0 or 3.0 can be used for magnification up to 100000x). It proved to work with both low and high voltages providing suitable electron transparent sample (Fig. 3) No DF images are available with this system.
Both STEM in SEM system are cheap and offer different configurations that can be adapted to sample characteristics, allowing to achieve good resolution and contrast with conventional SEM microscopes for sample screening before analysis with TEM or high-resolution dedicated STEM.


Fig. 1: STEM in SEM systems: STEM-A (a) and STEM-B (b) inside the closed (open in the insets) chamber of a XL 30 (FEI). Transmitted electrons signal is collected by BSE or SE detectors, respectively.

Fig. 2: Ag nano-particles on a holey carbon film coated grid imaged with the STEM-A system. Dark field image (a). Ag nano-particles on the surface of a CaCO3 micro-particle imaged by the STEM-B system equipped with the short graphite tunnel. Bright field image (b). Original magn. 204800x (a) and 51200x (b).

Fig. 3: Escherichia coli (a) and Staphylococcus epidermidis (b) cells imaged by the STEM-B system equipped with the short graphite tunnel. Samples were alcohol fixed, dehydrated, unstained. Microstructural details of cell surface (cell pili, black arrows), and inner structures (chromatin, white arrows) are visible. Original magn. 32000x (a) and 81000x (b).

Type of presentation: Poster

IT-4-P-6037 BASE – a web based, open source solution for microscope reservation and accounting

Iwan H.1, Timmermann J.1, Ritter M.1
1Hamburg University of Technology
ritter@tuhh.de

Central electron microscopy facilities often are responsible for all microscopes and peripheral devices such as sample preparation and analysis equipment at a University. Usually, central facilities are also responsible for the distribution of microscope time to other scientists, institutes or departments and they generally charge an hourly fee for microscope usage and for providing other services. Several scheduling systems, which were available at the time this project was initiated, offered only reservation features or were lacking at least one or more important requirements for accounting, such as project management or cost separation. Therefore, a new system for booking and accounting for microscope time has been developed and introduced at Hamburg University of Technology (TUHH) to reduce administrative burden. This new system is called BASE (Booking and Accounting System for Electron Microscopy) and is based on PHP and MySQL. It offers the following main features: Project management, project-based booking of microscope time, redistribution of microscope time to one or more projects after usage, setting caps on expenses, customer accounts with balance history, timelines (versatile categorization and accessibility of microscope time), automatically expiring microscope access (to enforce regular safety briefings), download area, invoice management etc. Therefore, BASE should fit most needs of central electron microscopy facilities in academic environments. It soon will be available as a web-based, open source solution.


Fig. 1: Schematic principle of BASE: BASE is a project-based booking and accounting solution for electron microscopy facilities. Microscope users work for one or more projects that are billed to a customer, i.e. an institute, department or an external company. Within a project, microscopes and services that are provided by the EM facility can be used.

Fig. 2: BASE login screen for TUHH showing the status of the microscopes.

Type of presentation: Poster

IT-4-P-6044 Energy filtered imaging in a scanning electron microscope

Boese M.1
1ZEISS Microscopy, Oberkochen, Germany
markus.boese@zeiss.com

Electron spectroscopy and spectroscopic imaging in a scanning electron microscope is currently used to separate secondary electrons (SE) and backscattered electrons (BSE) in SEMs. Due to limitations of the energy resolution the underlying spectral information of the emitted SE and BSE remains less investigated and mostly unused for imaging. I this study we will present some results of a retarding field detector with an improved energy resolution.

The ZEISS EsB (Energy selective Backscatter) detector was the first commercially available retarding field detector for a SEM. This detector has a very good surface sensitivity since it can work even at low Voltages. With energy filtered imaging it is possible to enhance material contrast even beyond the usual Z-contrast imaging.

Especially the energy filtering properties for low loss BSE imaging conditions were explored by Jaksch [1] and contrasts were shown which can not be explained with the common BSE contrast mechanism. The material dependent difference of the BSE spectral distributions is utilized here for energy filtered imaging to enhance the contrast between different phases.

In general the spectral distribution for BSE emitted from low Z materials are flat, whereas the spectra of high-Z elements have a pronounced maximum near the elastic peak. This behavior is due to the higher penetration depth and the resulting multiple scattering events of the BSEs. By use of energy windows a gain of contrast can be achieved which is superior to the commonly used Z contrast in conventional BSE imaging [2].

As an example for filtering BSEs we can achieve a strong contrast gain between Y2O3 and MoSi2 grains in a SiC matrix, by using only electrons with energy losses up to 100eV from a primary energy of 700V.
Furthermore energy filtered imaging of secondary electrons can enhance contrast for imaging conductors next to less conductive materials. Here the spectra of the secondary electrons are shifted towards each other for the different materials by several eV. As an example the contrast between multilayer graphene and a carbon support film will be presented.
The dependency of the BSE energy loss due to multiple scattering and sample depth can as well be used to gain 3D information from a sample for BSE tomography [3]. The selection of energy windows can be chosen from the elastic peak for pure surface information down to the maximum escape depth of the BSE with maximum energy losses. As an example the different thicknesses of a graphite sample can be imaged by using energy windows of different energy losses.


1. Jaksch, H.: Microsc. Microanal. 17,Suppl. 2 (2011), 902.
2. Cazaux, . J.: Electron Microsc. 61.5 (2012), 261.
3. Niedrig, H. and Rau, E. I.: Nucl. Instrum. Meth. Phys. Res. B 142(1998), 523.


IT-5. Analytical electron microscopy

Type of presentation: Invited

IT-5-IN-1534 Can orbitals be mapped in the TEM?

Löffler S.1,2, Bugnet M.3, Gauquelin N.3, Hambach R.4, Lazar S.3,5, Pardini L.6, Draxl C.6, Kaiser U.4, Botton G. A.3, Schattschneider P.1,2
1Institute of Solid State Physics, Vienna University of Technology, Austria, 2University Service Centre for Transmission Electron Microscopy, Vienna University of Technology, Austria, 3Canadian Centre for Electron Microscopy, McMaster University, Canada, 4Electron Microscopy Group of Materials Science, Ulm University, Germany, 5FEI Electron Optics, Eindhoven, Netherlands, 6Department of Physics, Humboldt University Berlin, Germany
stefan.loeffler@tuwien.ac.at

The energy, position, and momentum distributions of electrons inside a material are decisive for most of the material's properties, ranging from optical, electrical and magnetic properties to hardness, durability, or the melting point. Therefore, the electron distribution is a key quantity in many fields of research. Unfortunately, it is also elusive and directly imaging electronic orbitals and bonds in the bulk has not been possible so far.

In the last few years, several authors reported measurements of orbital properties using electron energy loss spectrometry (EELS) [1-3]. While these are great advances, it would be even better to actually "see" the orbitals in an image. Recently, the possibility to record maps of transition probabilities – from which orbitals can be deduced – was predicted theoretically [4]. In this work, we investigate the requirements and the feasibility to realize that prediction in an actual experiment.

On the one hand, the point group symmetry of the sample atoms plays a crucial role. Intuitively, this is readily understandable. Taking an isolated atom, for example, one is faced with a spherically symmetric problem. Clearly, its solutions must produce rotationally symmetric images. Hence, the symmetry of the system has an important influence on orbital maps. Here, we investigate the requirements on the crystal structure in order to be able to see a directional dependence of transition probabilities, orbitals, and bonds (see Fig. 1).

On the other hand, experimental parameters such as the signal to noise ratio (where the signal is the difference from the average; see Fig. 2), as well as the stability of the specimen and of the microscope are vital for successfully recording high-resolution maps. Based on experimental energy filtered images recorded with very high spatial resolution, we evaluate the requirements on both the sample and the microscope to obtain reproducible and directly interpretable maps. This nurtures the hope that orbital mapping will become a reality in the near future and will become an invaluable tool for many fields of research.

[1] Löffler et al., Ultramicroscopy 111 (2011) 1163
[2] Neish et al., PRB 88 (2013) 115120
[3] Hetaba et al., Micron, in print
[4] Löffler et al., Ultramicroscopy 131 (2013) 39


The authors acknowledge financial support by the FWF (I543-N20), the DPG, and the MWK Baden-Württemberg.

Fig. 1: Comparison of the K-edge maps for an Oxygen atom with full O(3) point-group symmetry (left) and with C2v symmetry (right). For the simulations, an acceleration voltage of 80 kV, a collection semi-angle of 24 mrad and ideal imaging conditions (Cs=0, Cc=0, df=0) were assumed.

Fig. 2: Predicted maps of the Ti L edge for a thin Rutile sample in [001] zone axis for different signal-to-noise ratios (SNR) as indicated (a-d). Preliminary experimental map as acquired (e). EELS signal of 1 px (dots) and averaged over 14000 px (line) (f). An acceleration voltage of 80 kV and an energy-window of 4 eV on the L2 edge were used.

Type of presentation: Invited

IT-5-IN-1656 Single atom imaging and spectroscopy with aberration-corrected STEM

Zhou W.1, Lupini A. R.1, Kapetanakis M. D.2,1, Lee J.3, Oxley M. P.2,1, Prange M. P.2,1, Pantelides S. T.2,1, Idrobo J. C.3, Pennycook S. J.4
1Materials Science & Technology Division, Oak Ridge National Laboratory, Oak Ridge, TN 37831, USA, 2Department of Physics and Astronomy, Vanderbilt University, Nashville, TN 37235, USA, 3Center for Nanophase Materials Sciences, Oak Ridge National Laboratory, Oak Ridge, TN 37831, USA, 4Department of Materials Science and Engineering, University of Tennessee, Knoxville, TN 37996, USA
wu.zhou.stem@gmail.com

Aberration-corrected scanning transmission electron microscopy (STEM) at low voltage can now provide real space imaging and spectroscopy measurements at the atomic scale with single atom sensitivity. This opens new opportunities for quantitative study of structural defects in 2D materials. Such studies, especially when combined with first-principles calculations, serve as an important step to correlate the defect structure with local properties, and help to create new functionalities in 2D materials via controlled defect engineering.

Figure 1 shows an experimental annular dark field (ADF) image of Se-doped MoS2 monolayer. Quantitative image analysis allows us to identify the chemical nature of the dopant and map out their distribution in MoS2, one atomic-layer at a time, providing a feasible way to quantify the local composition and measure the local band gap at the 10 nm scale [1].

Combining the imaging and spectroscopy power on a STEM, the changes in local optical response and electronic structure can be directly measured at defect sites. We show that the presence of a single Si atom in the graphene lattice can enhance the low-energy interband transitions with sub-nm spatial confinement [2]. Furthermore, the fine structure in electron energy loss spectra acquired under optimized dose levels provides the sensitivity to determine the nature of the chemical bonding of single atoms. We show that three-dimensional and planar bonding configurations for individual Si atoms in graphene can be directly discriminated (Figure 2) [3].

[1] Y. Gong et al., Nano Lett., 14, 442-449 (2014).
[2] W. Zhou et al., Nat. Nanotech., 7, 161-165 (2012).
[3] W. Zhou et al., Phys. Rev. Lett., 109, 206803 (2012).


This research was supported by a Wigner Fellowship of Oak Ridge National Laboratory (ORNL), the U.S. DOE Basic Energy Sciences, and ORNL’s Center for Nanophase Materials Sciences.

Fig. 1: Atom-by-atom Se dopant analysis in monolayer MoS2 adapted from Ref. [1]. (left) ADF image of Se-doped MoS2. (right) Structure model obtained from histogram analysis showing the distribution of single- and double-Se substituted S2 sites.

Fig. 2: Direct determination of the chemical bonding of single Si atoms via combination of ADF and spectrum imaging with first-principles calculations. (Left) ADF images of 3- and 4-fold coordinated Si atoms in graphene, and their respective Si L-edge fine structure extracted from spectrum images (Right). [3].

Type of presentation: Invited

IT-5-IN-5756 Advancing the spectroscopic frontiers of STEM

Stephan O.1, Arenal R.2, Bocher L.1, Bourrellier R.1, Colliex C.1, Gloter A.1, Kociak M.1, Losquin A.1, March K.1, Marinova M.1,3, Tararan A.1, Tencé M.1, Tizei L.1,4, Zobelli A.1
1LPS, UMR8502 CNRS, Université Paris-Sud, Orsay, France, 2LMA, INA, Universidad de Zaragoza, Zaragoza, Spain, 3UMET, Université Lille1, Villeneuve d'Ascq, France, 4AIST, Tsukuba, Japan
odile.stephan@u-psud.fr

The field of electron energy-loss spectroscopy (EELS) has recently achieved a succession of impressive successes linked with the development of aberration correctors, enabling atomically-resolved spectroscopy, which are now spreading worldwide. In addition, a new generation of monochromators is emerging, providing improvements in energy resolution of at least one order of magnitude and giving unprecedented access to low energy-loss ranges. The possibilities in elemental analysis have become exceptional, especially as the progress in EELS has been accompanied by advances in energy-dispersive X-ray (EDX) analysis thanks to improvements in signal detection efficiency. Similarly, recent progress in the collection of visible-range photons emitted by a sample illuminated by a focused beam, has enabled novel cathodo-luminescence (CL) experiments in the scanning transmission electron microscope (STEM). In addition, new ways of exploiting fast electron beams, including combining them with beams of photons, have opened up the field of nano-optics, providing a high-spatial resolution alternative to more conventional optical techniques. Thus, STEM instruments are now extremely versatile, allowing for the simultaneous detection of an increasing variety of signals. Some of these new possibilities will be illustrated. For example, going beyond elemental mapping to measure fine structure (ELNES) variations at the scale of individual atomic columns or atoms will be described. Spectroscopic data acquired with EM or low-noise CCD cameras will be discussed in connection with the quantitative measurement of electron densities and the identification of charge ordering in oxide materials [1] or probing the chemical bonds of heteroatoms hosted in a carbon lattice [2]. Novel experiments combining EELS and CL, using a dedicated, home-made, high-efficiency nano-CL system will be presented, demonstrating how the usual macroscopic concepts such as extinction, absorption, and scattering cross-sections are no longer sufficient to describe optical phenomena at the nanoscale [3]. When used in combination with TEM structural investigations, nano-CL experiments have proved to be a unique way to explore the intimate link between a crystal structure (h-BN), its defects and its optical properties [4].
These experiments open stimulating perspectives for the development of further new spectroscopic techniques, combining photons and electrons in time-resolved applications for example, or entering the field of quantum optics.

[1] L. Bocher et al Phys.Rev. Lett. 111 (2013) 167202

[2] R. Arenal de la Concha et al, arXiv:1401.5007

[3] A. Losquin et al, submitted

[4] R. Bourrelier et al, arXiv:1401.1948


The work has received funding from the french CNRS-CEA METSA network and the European Programme ESTEEM2

Fig. 1: Schematics of the Orsay STEM set-up dedicated to nano-optics experiments. The HAADF image of a gold nanoprism is shown in combination with the EELS and CL signal intensity maps (not acquired simultaneously in that specific case) revealing the spatial variation of the « tip » plasmon mode of the particle.

Type of presentation: Oral

IT-5-O-1510 Quantitative electron magnetic circular dichroic signals acquired by 1000 kV (S)TEM-EELS

Tatsumi K.1, Kudo T.2, Muto S.1, Rusz J.3
1EcoTopcia Science Institute, Nagoya University, Nagoya 464-8603, Japan, 2Graduate School of Engineering, Nagoya University, Nagoya 464-8603, Japan, 3Department of Physics and Astronomy, Uppsala University, Box 516, SE-75120 Uppsala, Sweden
k-tatsumi@nucl.nagoya-u.ac.jp

Electron magnetic circular dichroism (EMCD) in EELS is an attractive technique for studying spin-related properties in a higher spatial resolution than the well-utilized x-ray magnetic circular dichroism (XMCD). Even though several effective experimental geometries of the EMCD measurement have been established, these are still substantially in the testing stage due to the intrinsic difficulties associated with quantitative analysis, such as a low signal-to-noise ratio (SNR) and complicated dynamical and plural inelastic scattering effects depending on the sample thickness/orientation. The present study shows the EMCD measurements using a (S)TEM-EELS with an acceleration voltage V of 1000 kV, toward more quantitative local magnetic analysis.
Figure 1 shows the example results of Co L2,3 EELS near edge structures (ELNES) with 200 and 1000 kV acceleration voltages in the intrinsic EMCD method [K. Tatsumi et al., Microcopy, 2014; doi: 10.1093/jmicro/dfu002]. The dichroic signals as the difference spectra showed a better SNR in the 1000 kV case, because of the larger fraction of the dichroic signals.
Figure 2 shows theoretical fractions of the dichroic signals, here represented by a quantity, fd = (IB - IA) / (IA + IB) at the L3 peak energy, as a function of sample thickness t, with several different collection angles φcol represented as EELS aperture radii in units of g shown in Fig. 1. The calculations were performed based on the dynamical diffraction and single core loss scattering. The desirable collection angle is less than 0.2 g because the larger φcol significantly decreases fd. fd at 1000 kV are significantly larger than 200 kV for relatively large thicknesses (t = 30 to 40 nm), because of the larger extinction distance and inelastic mean free path. The simulated and experimental ratios of fd at 1000 kV to fd at 200 kV for t= 35 nm are 2.5 and 3.0, respectively, showing reasonable consistency.
Finally, this advantage is utilized to statistically acquire quantitative EMCD signals distributed over the diffraction plane, demonstrating that quantitative magnetic information can be routinely obtained using electron beams of only a few nanometers in diameter without any restriction regarding the crystalline order of the specimen [S. Muto et al., Nature Commun., 2014; doi: 10.1038/ncomms4138].


This work was partly supported in Grant-in-Aids for Scientific Research of JSPS (Wakate A: 24686070 and Innovative areas: 25106004) and Swedish Research Council.

Fig. 1: Experimental Co-L2,3 ELNES collected at two different EELS aperture positions A and B. Two sets of results obtained by using different TEM-EELS systems with different acceleration voltages, 200 kV (a) and 1000 kV (b), are shown.

Fig. 2: Theoretical fd with V = 200 and 1000 kV. Numbers inset are φcol, represented by EELS aperture radii in g. Filled circles are results with the experimental φcol.

Fig. 3:
Type of presentation: Oral

IT-5-O-1718   Towards Quantitative EDX Results in 3 Dimensions.

Goris B.1, Freitag B.2, Zanaga D.1, Bladt E.1, Altantzis T.1, Sudfeld D.2, Bals S.1
1EMAT, University of Antwerp, Antwerp, Belgium , 2FEI Company, P.O. Box 80066, KA 5600 Eindhoven, The Netherlands
sara.bals@ua.ac.be

Over the last 10 years, electron tomography has evolved into a versatile tool to investigate (hetero)nanostructures [1]. Nevertheless, resolving their chemical composition in 3D remains challenging. In principle, energy dispersive X-ray (EDX) mapping can be combined with electron tomography since the number of generated X-rays increases with sample thickness. However, early attempts to perform 3D EDX experiments were complicated by the specimen-detector geometry [2]. Recent efforts therefore led to a novel EDX detection system, enabling the extension of EDX mapping to 3D [3]. An example of a 3D EDX reconstruction is shown in Figure 1, showing a Au@Ag nanocube of which the Au core yields an octahedral shape. This example clearly illustrates the potential of 3D EDX mapping, but one needs to be careful when extracting quantitative information from such reconstructions. In order to obtain quantitative 3D reconstructions using EDX, different steps in the experiment need to be optimized.

The Super-X detection system consists of 4 EDX detectors that are symmetrically arranged around the sample. As a result, it is expected that shadowing effects are minimized and that the total number of detected characteristic X-rays for a spherical nanoparticle is independent of tilt angle. Figure 2 presents the EDX counts that were acquired from a Au particle using a Model 2030 Fischione tomography holder. Using this dedicated holder, shadowing is kept at a strict minimum, but even in this case, an asymmetric collection efficiency of the detector is still observed. This problem, caused by remaining shadowing of the sample grid, can be overcome by combining EDX signals, unaffected by shadowing, that are collected by different detectors during the tilt series.

Quantification of the EDX maps is typically performed using the “Cliff-Lorimer” method, originally developed for the investigation of thin films. Here, we evaluate the use of the “ζ (zeta)-factor” method to obtain quantitative 3D chemical data using the following equation [4]:

ρt=ζ I ⁄ (CDe)

In this formula, ρ is the density of the material and t equals sample thickness, which can be obtained from 3D high angle annular dark field STEM (HAADF-STEM) reconstructions. First, the ζ-factor can be determined by measuring the intensity I and the electron dose for monometallic nanostructures. After estimation of the ζ-factors for different elements, quantitative 3D elemental analysis becomes possible for heteronanomaterials having unknown composition.

[1] PA Midgley, RE Dunin-Borkowski, Nature Materials 8 (2009), p.271

[2] G Möbus, RC Doole and BJ Inkson, Ultramicroscopy 96 (2003), p.433

[3] P Schlossmacher et al, Microscopy Today 18 (2010), p.14

[4] M Watanabe and DB Williams, Journal of Microscopy 221 (2006) p.89


The authors acknowledge support from the European Research Council (ERC Starting Grant -COLOURATOMS) and the FWO.

Fig. 1:  (a) 2D EDX map of a Au@Ag nanocube. Based on a tilt series of such 2D EDX maps, 3D reconstructions (b) could be obtained showing the 3D distribution of the different chemical elements.

Fig. 2: (a) Detected X-ray counts as function of tilt angle for each individual detector of the super-X system. At certain tilt angles, shadowing effects may block the X-rays preventing them to reach the detectors. (b) Total X-ray count when adding the signal from different detectors.

Type of presentation: Oral

IT-5-O-1642 Surface-plasmon-polariton coupling between adjacent submicron slits in a Au-film investigated by STEM-EELS

Fritz S.1, Walther R.1, Schneider R.1, Gerthsen D.1, Matyssek C.2, Busch K.2,3, Maniv T.4, Cohen H.5
1Laboratorium für Elektronenmikroskopie, Karlsruher Institut für Technologie (KIT), Karlsruhe, Germany, 2Humboldt-Universität zu Berlin, Institut für Physik, AG Theoretische Optik & Photonik, Berlin, Germany, 3Max-Born-Institut, Berlin, Germany, 4Schulich Faculty of Chemistry, Technion – Israel Institute of Technology, Haifa, Israel, 5Department of Chemical Research Support, Weizmann Institute of Science, Rehovot, Israel
stefan.fritz@partner.kit.edu

Surface plasmon polaritons (SPP) and associated cavity modes (CM) were recently analyzed by STEM-EELS in submicron slits in thin metal films.1, 2 Moreover, a strong enhancement of the CM was observed upon introduction of neighboring slits.3 Such nanostructures exhibit extraordinary optical transmission, which is further enhanced due to SPP coupling between slits.4
In this work, EEL spectra were acquired in a monochromated FEI Titan3 80-300. Background subtraction was performed by a fit to the ZLP tail. Slits with a size of 180 x 900 nm2 were milled in a 200 nm Au-film by FIB milling. Numerical simulations were performed with the Discontinuous Galerkin Time Domain method adapted for EELS.5
Fig. 1a shows spectra acquired at 10 nm distance to the walls of a double slit (cf. dots in the HAADF STEM image). In addition to the Au surface plasmon (SP) at ~2.4 eV, signals at 0.5 and ~1.5 eV are resolved which correspond to the fundamental ω1 and 3rd harmonic ω3 of a CM hybridized with SPPs supported by the metal wall. A significant enhancement of ω1 and ω3 is found close to the inner wall. Also, ω3 is red-shifted and the SP is reduced in intensity. Fig. 1b shows corresponding simulated spectra which agree well with the experiments.
The coupling was studied in double slits with inter-slit distances (p) of 280, 450, 900, 1080, and 1980 nm. Fig. 2a shows spectra acquired at 10 nm distance from the inner walls. For increasing p values, ω1 is red-shifted from 0.5 to 0.4 eV and reduced in intensity, nearly vanishing at p=1080 nm. At an even larger p value, ω1 is observed again at ~0.5 eV. Fig. 2b shows corresponding simulations agreeing well with the experiments. For p > 880 nm, a second signal at higher energy is observed which red-shifts and increases in intensity up to p=1580 nm. These two signals correspond to symmetric and anti-symmetric coupling of the hybridized SPP CM between both slits. At p=1080 nm, two weak modes are observed in the simulation which corresponds to the almost vanishing loss intensity in the experimental spectrum. For further increasing p the symmetric mode increases in intensity which is also observed in the experimental spectra for p=1980 nm. The latter effect impressively demonstrates the effect of coherent interference between SPPs of adjacent slits across the metal bar.

1 I. Carmeli et al., Phys. Rev. B 85 (2012).
2 B. Ögüt et al., ACS Nano 5, 6701 (2011).
3 R. Walther et al., arXiv:1212.1987 (2013).
4 F. J. Garcia-Vidal et al., Rev. Mod. Phys. 82, 729 (2010).
5 C. Matyssek et al. Photonics Nanostruct. 9, 367 (2011).


Funding by the DFG Research Center for Functional Nanostructures, the Ministry of Science, Research and the Arts of BW, and DFG project Bu 1107/7-2 (KB and CM).

Fig. 1: a) Experimental and b) simulated EEL spectra at the outer (black line) and inner (red line) walls in a double slit system (cf. HAADF STEM image in inset). The scale bar corresponds to 500 nm.

Fig. 2: a) EEL spectra from double-slit systems with varying p detailing the evolution of ω1. The signal at 0.9 eV for p=1980 nm corresponds to the energy of ω2. This signal was excited due to SPP coupling between the two slits despite having a node at the measurement position (see red dot in Fig. 1a).b) Simulated EEL spectra as a function for increasing p

Type of presentation: Oral

IT-5-O-1653 New EM signals made accessible by sub-20 meV resolution EELS

Křivánek O. L.1, Lovejoy T. C.1, Aoki T.2, Crozier P. A.2, Rez P.3, Egerton R. F.4, Dellby N.1
1Nion Co., 1102 Eight St, Kirkland, WA 98033, USA, 2Center for Solid State Science, Arizona State University, Tempe, AZ 85287, USA, 3Department of Physics, Arizona State University, Tempe, AZ 85287, USA, 4Department of Physics, University of Alberta, Edmonton T6G 2E1, Canada
krivanek@nion.com

Nion’s High Energy Resolution Monochromated EELS-STEM (HERMES) instrument [1] is able to combine Scanning Transmission Electron Microscopy (STEM) spatial resolution of a 1-10 Å with 12-50 meV Electron Energy Loss Spectroscopy (EELS) energy resolution. These capabilities promise to make new signals available in analytical EM, and thereby revolutionize it even more than aberration correction has revolutionized EM imaging. Here we explore two new signals: spatially-resolved phonon spectroscopy, and the detection of very light elements by energy-filtered imaging of electrons scattered to high angles.

Fig. 1 illustrates how important signals have up to now been “hidden in plain sight” – obscured by a broad EELS zero loss peak (ZLP). The solid green spectrum was recorded with the beam passing through the monochromator but the energy-selecting slit retracted. The full-width at half-maximum (FWHM) of the ZLP is ~250 meV. The red (line) spectrum was recorded with the slit in, in 0.1 s, and shows FWHM of 14 meV.

The blue (x1000) spectrum in Fig. 1 was recorded from a ~2 nm Ø area of SiO2. The optical phonon peak visible at 140 meV energy loss is in good agreement with the energy of the strongest feature in infrared spectra of SiO2, at 1100 cm-1. (To convert cm-1 to meV, divide by 8.)

Fig. 2 demonstrates that some phonon signals can be spatially resolved with a resolution of a few nm, and hopefully better in the future. The phonon intensity decays close to zero within ~3 nm inside the Si and there is also an initial sharp intensity drop-off at the SiO2–vacuum interface. There is also a long tail stretching tens of nm outside the sample, which suggests that damage-free phonon spectroscopy may be possible with an aloof electron beam.

Imaging phonons in compounds containing light elements such as H should allow the spatial distribution of the compounds to be mapped. It may, however, also be possible to image the light elements in a more general way, by using the fact that electrons scattered incoherently by atomic nuclei to high angles (Rutherford scattering) transfer small amounts of energy to the recoiling nuclei, inversely proportional to their mass.

Fig. 3 shows proof-of-principle energy-filtered high-angle dark field (EFHADF) mapping of light vs. heavy atoms: 60 keV spectrum-image data from Au particles supported on an amorphous carbon foil ~20 nm thick, next to a hole in the foil. Energy window B is centred on the ZLP (±10 meV) and the corresponding image 3(b) shows mainly Au particles. Energy window C is placed over energy losses of 85±10 meV, and image 3(c) shows only carbon. The fact that we are able to image only the carbon shows that we now have sufficient energy discrimination to map very light elements such as H and Li [2].

[1] OL Krivanek et al., Microscopy 62 (2013) 3-21.
[2] TC Lovejoy et al., M&M meeting (2014, Hartford).
[3] We are grateful for the use of LeRoy Eyring Center facilities at ASU.


Fig. 1: EEL spectra recorded under various conditions by Nion HERMES at 60 keV. Solid (green) spectrum: monochromator slit out; red spectrum: slit in; blue (x1000) spectrum: slit in, electron probe on SiO2, acquisition time 10 s, beam current ~10 pA, probe convergence angle ±12 mrad, collection angle ±12 mrad

Fig. 2: a) HAADF image of a Si-SiO2 cross-section; b) profile of SiO2 phonon intensity and sample thickness along the red line in (a). The SiO2 phonon intensity was measured from a series of 100 spectra in 10 s each, normalized by the ZLP. The sample thickness was determined from spectra of all energy losses up to 180 eV, recorded separately.

Fig. 3: EFHADF recoil mapping of Au on am. carbon: a) EEL spectra from a Au particle and from carbon film; b) image formed with window B showing only Au; c) image formed with window C showing only carbon. The angles admitted into the spectrometer were 120±30 mrad, and scattering from carbon nuclei was expected to give a broad peak centered on 40 meV.

Type of presentation: Oral

IT-5-O-1665 Atom-by-atom chemical imaging of topological insulator nanostructures by ChemiSTEM

Jiang Y.1, Wang Y.1, Zhang Z.1
1Center of Electron Microscopy and State Key Laboratory of Silicon Materials, Department of Materials Science and Engineering, Zhejiang University, Hangzhou, China
jiang0209@zju.edu.cn

Topological insulators (TIs) have attracted ever-increasing attention due to their exotic physical phenomena, however, the overwhelming majority of reported work was focused on the physical properties [1,2]. In contrast, limited effort has been made to gain an accurate picture for their chemical compositions at atomic level, although such information is of critical importance to comprehend their demonstrated properties. Here by employing a state-of-the-art atomic-mapping technology (ChemiSTEM), we present a direct atom-by-atom chemical identification of nanostructures and defects in TIs [3]. We first identify and explain the layer-chemistry evolution of Bi2Te3-xSex TIs (Fig. 1). Significantly, we reveal a long neglected but crucially important defect which is universally present in Bi2Te3 films, the seven-layer Bi3Te4 nano-lamella (Fig. 2). This nano-lamella may explain inconsistencies in measured conduction type as well as open up a new route to manipulate the bulk carrier concentration. This work may pave the way to thoroughly understand and tailor the nature of the bulk, and to secure controllable bulk states for their future dissipationless devices.

References
1. Xiu, F.; He, L.; Wang, Y.; Cheng, L.; Chang, L.-T.; Lang, M.; Huang, G.; Kou, X.; Zhou, Y.; Jiang, X.; Chen, Z.; Zou, J.; Shailos, A.; Wang, K. L. Nature Nanotechnology 6, 216 (2011).
2. Wang, Y.; Xiu, F.; Cheng, L.; He, L.; Lang, M.; Tang, J.; Kou, X.; Yu, X.; Jiang, X.; Chen, Z.; Zou, J.; Wang, K. L. Nano Letters 12, 1170 (2012).
3. Jiang, Y.; Wang, Y.; Sagendorf, J.; West, D.; Kou, X.; Wei, X.; He, L.; Wang, K. L.; Zhang, S. B.; Zhang, Z. Nano Letters 13, 2851 (2013)


We acknowledge the support of NSFC (No. 11174244), the National 973 Program of China (2013CB934600), Zhejiang Provincial NSFC (LR12A04002).

Fig. 1: Layer-chemistry evolution of Bi2Te3-xSex (x=0, 1, 2, 3).

Fig. 2: Structural and chemical identifications of the 7-layer lamellae in Bi2Te3.

Type of presentation: Oral

IT-5-O-1860 Nanocharacterisation of vanadium and niobium carbide and carbonitrides precipitates in austenite high manganese steels with DualEELS and HAADF techniques

Bobynko J.1, MacLaren I.1, Craven A. J.1, McGrouther D.1, Paul G.2
1Kelvin Nanocharacterisation Centre, SUPA School of Physics and Astronomy, University of Glasgow, Glasgow G12 8QQ, UK, 2ThyssenKrupp Steel Europe AG, Research and Development, FuE-E Modelling and Simulation, Kaiser-Wilhelm-Straße 100, 47166 Duisburg, Germany
j.bobynko.1@research.gla.ac.uk

In order to fulfil the twin goals of safety enhancement and reduced CO2 emissions combined with fuel efficiency for automotive applications, steels combining high strength and high ductility are required for structural members. One alloy system of interest for this application is high-manganese steel microalloyed with either Nb or V to provide high densities of nanoscale carbide or carbonitride precipitates to provide additional strengthening by the dispersion hardening mechanism. We report significant progress in the understanding of the atomic structure of such precipitates, together with their nanoscale chemistry using aberration-corrected analytical STEM. Specifically, we report the details of the precipitate morphology, and absolute measurements of the chemical composition, and use this to make plausible reconstructions of the 3D morphology of these few nm precipitates. This is achieved using DualEELS spectrum imaging datasets recorded with a Gatan GIF Quantum on a JEOL ARM 200F cold FEG, probe-corrected STEM at pixel spacing of a few Ångströms.
We show that a processing approach consisting of subtracting all artefacts from the dataset, followed by the subtraction of the matrix components from the dataset allows the creation of a spectrum image just representing the precipitate. This can then be quantitatively analysed both to produce an accurate thickness map using the t/λ method, and also to produce chemical maps of both the metal cation and carbon contents. Excellent correlation is found between all the maps showing that all elements are uniformly distributed throughout the precipitate, and it is shown that absolute thicknesses can be calculated using crystallographic parameters together with standards for the cross sections of Nb, V, C and Ti (present as impurities). It is shown that there is little evidence for any significant content of N in the precipitates, as is expected for an Al containing steel, and consequently we see little evidence of TiN nanoclusters acting as nucleation points for the precipitates. Rather, Ti impurities have exactly the same spatial distribution in the precipitates as the V or Nb. Nevertheless, we also show that in at least one case, a tiny concentration of N could be detected and mapped within a precipitate, which seems in this case to have been distributed across the whole precipitate.


We are grateful to the European Commission for the funding of this work as part of a Research Fund for Coal and Steel Project (Precipitation in High Manganese Steels, RFSR-CT-2010-00018). This work was only possible because of the generous provision of the MagTEM facility by SUPA and the University of Glasgow.

Fig. 1: Qualitative maps processed from spectrum images of vanadium carbide and niobium carbide precipitates, both from steels isothermally treated at 900°C for a fixed time of 100s followed by quenching. The vanadium precipitate after the background subtraction reveals small traces of N.

Fig. 2: The background subtracted spectrum of the core loss EEL spectrum from the central region of the NbC precipitate, and profiles of the integrated counts in a horizontal line across the centre of the precipitate. Ti counts have been multiplied by a factor of 10 for the purpose of this graph.

Type of presentation: Oral

IT-5-O-1775 Study of layer by layer graphitization of 4H-SiC, through atomic-EELS at low energy

Nicotra G.1, Deretzis I.1, Ramasse Q.2, Longo P.3, Scuderi M.1, Twesten R. D.3, La Magna A.1, Giannazzo F.1, Spinella C.1
1CNR-IMM, Strada VIII, 5, 95121 Catania, Italy, 2SuperSTEM Laboratory, STFC Daresbury Campus, Daresbury WA4 4AD, United Kingdom, 3Gatan, Inc., 5794 W Las Positas Blvd, Pleasanton, CA, 94588, USA
giuseppe.nicotra@imm.cnr.it

Epitaxial graphene (EG) grown on Si-polarized SiC, play a crucial role by the presence of a so-called carbon “buffer layer”. Such layer has been shown to present a certain degree of sp3 hybridization since it is partially bound to the outmost Si atoms of the SiC (0001) surface [1].
Our results indicate a layer by layer graphitisation of the SiC as the Si evaporates. Atomic resolution EELS measurements show that the relative Si concentration across the buffer layer [2]. Moreover, the presence of oxygen has been revealed across the buffer layer as shown in Figure 1b. The presence of oxygen could be responsible of the slower decomposition of the SiC into graphitic layers. This and other aspects will be discussed.
All the STEM and atomic EELS measurements were performed at 60k. This consists of a probe corrected STEM microscope, capable to deliver a probe size of 1.1 Å, and equipped with a C-FEG and a fully loaded GIF Quantum ER as EELS spectrometer. Low- and core-loss spectra were nearly simultaneously acquired using the DualEELS capability. In this way an accurate measurement of the π*/σ* peaks ratio that is proportional to the sp2 contribution can be carried out. Low- and core-loss EELS spectra were taken across the green box in the ADF STEM image in Figure 1a using a pixel step size of 0.6Å and an exposure time of 20 ms for each pixel. The spectrometer was set to 0.25eV dispersion yielding 0.75eV energy resolution. Such energy resolution is sufficient to reveal different features in the fine structure of the C K-edge and Si L2,3-edges. The ADF STEM image in Figure 1a shows the presence of the buffer layer between the SiC substrate and the 3 graphitic layers. EELS spectra of the O K-edge, Si L2,3-edges and C K-edge are shown in Figures 1b,c,d respectively and were extracted from the selected positions in the sample as shown in Figure 1a. In Figure 1b, the O K-edge peak shows up only in regions 4 - 6 and is particularly strong in region 6. There seems to be in this region of the buffer layer an increase of the oxygen concentration. No oxygen is detected in either the SiC substrate or the graphitic layers. Particularly interesting are the C K-edge spectra in Figure 1d. The spectra in positions 1-3 in the SiC substrate region show different π* peak, indicating chemistry changes. The spectra extracted from the graphene layers in positions 7,8,9 show much higher contribution in the π* peak that leads to the fully sp2 hybridization indicating transition to graphitic structure. [1] G Nicotra et al, ACS Nano 7 (4), (2013) p. 3045.2 [2] G Nicotra et al, to be published


This work was performed at Beyondnano CNR-IMM, which is supported by the Italian Ministry of Education and Research (MIUR) under project Beyond-Nano (PON a3_00363); The SuperSTEM  Laboratory is supported by the U.K. Engineering and Physical Sciences Research Council (EPSRC)

Fig. 1: a) ADTE STEM survey image. b-d) EELS spectra extracted from the selected positions in Figure 1a each spectrum was corrected for the effects of energy drift and plural scattering; b) O K at 532 eV slightly enlarged for better visualization; c) Si L2,3-edges at 99 eV; d) C K-edge at 284eV. Spectra from positions 7 – 9 are from the graphitic layers.

Type of presentation: Oral

IT-5-O-1847 Atomic resolution mapping of localized excitations using STEM-EELS spectroscopy

Gauquelin N.1, Egoavil R.1, Martinez G. T.1, Van Aert S.1, Van Tendeloo G.1, Verbeeck J.1
1EMAT,University of Antwerp, B2020 Antwerp, Belgium
nicolas.gauquelin@uantwerpen.be

Atomically resolved EELS experiments are commonplace in modern aberration-corrected transmission electron microscopes. Energy resolution has also been increasing steadily with the continuous improvement of electron monochromators [1]. Improving the energy resolution further seems attractive in order to study phonon lattice vibrations which typically occur between a few meV and 1 eV, but with very large scattering angle (10-1000 mrad).
However these interactions are known to be strongly delocalized due to the long range interaction of the charged accelerated electrons with the electrons in a sample, and were found to scale approximately inversely proportional to the energy loss. This has made several scientists question the value of combined high spatial and energy resolution for mapping interband transitions and possibly phonon excitation in crystals.
For phonon excitations, the fast electron couples to a lattice vibration mode via Coulomb interaction, which might result in strong localization of the scattering. [2]
On the other hand, preservation of elastic contrast in low-loss EELS mapping has been reported by S. Lazar et al. [3], where the filtered image of the zero loss peak (ZLP) intensity shows the complementary nature of the high angle annular dark field (HAADF) intensity and the elastic contrast. Furthermore, atomically resolved signatures were observed at 3 eV using high collection angle (124 mrad) and addressed as being possibly related to phonon assisted losses.
In this work we demonstrate experimentally that atomic resolution information is indeed available at very low energy losses of a few hundreds meV expressed as a modulation of the broadening of the zero loss peak. [4] Careful data analysis allows us to get a glimpse of what are likely phonon excitations. On figure 1, one can note the strong presence of atomic resolution contrast hinting to localized inelastic phonon scattering. The contrast vanishes for higher energy losses where delocalized electronic excitations prevail. On Figure 2e and 2f important deviations from the average when changing probe positions within a unit cell can be observed. The spectra show both gain and loss contributions in region where multiple phonon losses are expected. These experiments confirm recent theoretical predictions on the strong localization of phonon excitations [2] as opposed to electronic excitations and show that a combination of atomic resolution and recent developments in increased energy resolution will offer great benefit for mapping phonon modes in real space.
[1] O. L. Krivanek et al., Philos. Trans. A367, 3683 (2009).
[2] C. Dwyer, http://arxiv.org/abs/1401.6305 (2014).
[3] S. Lazar et al., Microsc. Microanal. 16, 416 (2010).
[4] R. Egoavil et al., in preparation (2014).


This work was supported by funding from the ERC grant 246791 - COUNTATOMS and ERC Starting Grant 278510 VORTEX. R. E. acknowledges funding from the  FP7 program grant nr NMP3-LA-2010-246102 IFOX. All authors acknowledge support from the FP7 Program Reference No. 312483-ESTEEM2. The fund for scientific research Flanders is acknowledged for funding FWO project G.0044.13N and G.0064.10N.

Fig. 1: Fig.1: (a) ADF survey image of an EELS SI acquisition taken on [100] SrTiO3 at 120 kV and collection angle β = 129 mrad. (b,c) Corresponding integrated EELS signal in a small window above the ZLP obtained by dividing each spectrum by a scaled average spectrum (b) or by subtracting (c) a scaled averaged spectrum from each individual spectrum.

Fig. 2: Fig.2: Same data averaged over 16 unit cells (a) ADF image (b,c) subtracted and divided maps over the same energy range. Normalized EELS spectra can be extracted for the 3 different types of atomic columns Sr (green), TiO (red) and O (blue) (d) logarithmically scaled zero loss peaks and (e,f) spectra for both division and subtraction treatments.

Type of presentation: Oral

IT-5-O-1900 Calculation and Measurements of XEDS Collection Solid Angle in the AEM

Zaluzec N. J.1
1Electron Microscopy Center, Argonne National Laboratory, Argonne IL, USA
zaluzec@microscopy.com

      

        One of the most often misquoted parameters of a solid state x-ray detector interfaced to the AEM is its collection solid angle (Ω). Closed form analytical equations have been developed to calculate the solid angle of six of common geometries of solid-state x-ray detectors[1]. These include cylindrical, rectangular and annular configurations, with the detector in either a tilted or untilted configuration as illustrated in figure 1. These formulae have been integrated into an on-line calculator which is freely accessible and only requires a Javascript compatible browser. It can be accessed at: http://tpm.amc.anl.gov/NJZTools. The use of this tool, removes the ambiguity which besets the community in assessing the relative merits of different manufacturer’s claims, by providing an independent procedure to assess the characteristics of difference detector sizes and their geometries.

       

        Due to the advances in the design and construction of modern Silicon Drift Detectors (SDD) the collection solid angle which in the past has hovered about 0.1-0.15 sR, is now routinely quoted in the 0.2-0.8 sR range. There is is a very little real data published which can be used as a direct measurement of the absolute solid angle or real comparisons of individual system other than anecdotal observations. To illustrate this variation we show measurements of the Ni Kα shell x-ray intensity on 4 different instruments all normalized to the same operating conditions (Figure 2). The performance of a series of 8 analytical electron microscopes, operated at 200 kV were tested independently and their collection solid angle determined. This was accomplished through the measurement of the absolute intensity/nA/nm of the x-ray signal emitted from an amorphous 10 nm thick Germanium specimen[2]. The results of these measurements for the different microscopes, with 18 different detectors configurations are summarized in Figure 3.  The largest individual detector tested had a nominal solid angle of ~ 0.24 sR, which agrees well with the calculated value obtain for it using the on-line calculator. For systems with multiple detectors each individual detector was measured, the net solid angle in such situations is the simple sum of from each independent detector value (i.e. instruments 4, 6, 9). The relative variation in the experimental measurements was very large (0.02 sR – 0.24 sR) illustrating the need to quantitatively assess detectors in terms of their real solid angle, rather than the often misinterpreted “Detector Area”. 

     

 References

[1] N.J. Zaluzec, Microsc Microanal, 20, in press (2014)

[2] N.J. Zaluzec, Microsc Microanal, 19 (Suppl 2) , 1262-1263, (2013)

doi: 10.1017/S1431927613008301


This work was supported by the U.S. DoE, Office of Basic Energy Sciences, Contract No. DE-AC02-06CH11357 at the Electron Microscopy Center at Argonne National Laboratory.

Fig. 1: Figure 1.) User interface to the on-line XEDS Solid Angle Calculator (http://tpm.amc.anl.gov/NJZTools)

Fig. 2: Figure 2.) Experimental variation of the performance of 4 different 30 mm2 detector systems for the same NiO specimen.

Fig. 3: Figure 3.) Experimental variation of the solid angle as a function of instrument for a 10nm thick amorphous Ge specimen.

Type of presentation: Oral

IT-5-O-1937 In situ observation of symmetry and bond length induced ELNES changes of the CaCO3 to CaO phase transformation by in-column energy-filtered low-kV TEM

Golla-Schindler U.1, Klingl T.1, Kinyanju M. K.1, Benner G.2, Orchowski A.2, Kaiser U.1
1Group of Electron Microscopy of Material Science,Ulm , Germany, 2 Carl Zeiss Microscopy GmbH, Oberkochen, Germany
ute.golla-schindler@uni-ulm.de

Calcite (CaCO3) is an important system for biomineralization and of enormous importance in the construction industry. When used as material for the construction of buildings, calcite is transformed into CaO by the reaction: CaCO3→CaO+CO2↑ releasing CO2. The same phase transformation of calcite is initiated in the electron microscope. Low dose rates and the low accelerating voltages achievable with the SALVE prototype microscope, enable to monitor right from the start in situ the phase transformation and track the change of chemistry and electronic environment by EELS and ELNES. The coordination of the Ca atom in calcite is close to an octahedral coordination if the ligands were aligned along the x, y and z axis and the octahedron is stretched along the <111> direction. The angles between the Ca-O bonds are 87.25° and 92.75° and the bond length is 2.357 Ǻ in all directions. In cubic CaO the calcium atom is octahedral coordinated with a bond length between Ca and O atoms of 2.407 Ǻ. Induced by the octahedral coordination of calcium in calcite and CaO the degenerated energy levels of the 3d shell split up into t2g and eg energy levels with an energy difference of Δo. Figure 1 presents the ELNES of the Ca L2,3 edge for the 40 kV and 20 kV series. Four peaks can be separated, with two of them correlated to the Ca 2p3/2 →3d (t2g eg) and two of them to the Ca 2p1/2 →3d (t2g eg) transitions. For 40 and 20 kV, we obtained starting peak positions of the Ca-L2,3 edges-characteristic for calcite, indicated by blue dashed lines. The second end peak-position-characteristic for CaO are indicated by red dashed lines. In the mid of the phase transformation a superposition of the starting and end spectra with resolved peak positions is detectable. Solely these two peak positions exist in all spectra. This proved that the change in the peak position cannot be initiated by energy drift of the experimental setup but is caused by the change of the electronic environment of the Ca atom. All peaks shift to lower energies and additionally the energy level splitting Δo increases. During the phase transformation of calcite to CaO, the octahedral coordination of calcium with oxygen as binding partner is preserved, but the distortion is removed and the bond length changes by 5 pm. Calculation of theoretically fitted EELS spectra with the CTM4X4S and Wien code are started to separate both effects. Muller [1] showed that changes in bond lengths generate core level shifts and his findings show the same relationship as our results, but the bond length variation and following the energy level shift are reduced approximately by a factor of 10 [2].

[1] Muller D. (1999). Ultramicroscopy 78

[2] Golla-Schindler U., Benner G., Orchowski A., Kaiser U. Microsc. Microanal. accepted


This work was supported by the DFG (German Research Foundation) and the Ministry of Science, Research and the Arts (MWK) of Baden-Wuerttemberg in the frame of the SALVE (Sub Angstrom Low-Voltage Electron microscopy and spectroscopy project.

Fig. 1: Figure 1. 40 kV (a) and 20kV (b) EELS spectra (raw data) of the time dependent changes of the Ca-L2,3 edge . The blue dashed lines are aligned to the peak position of the Calcite ELNES and the red dashed lines to the peak positions of CaO ELNES.

Type of presentation: Oral

IT-5-O-2005 Electron energy losses and cathodoluminescence from complex plasmonic nanostructures : spectra, maps and CL radiation patterns from a generalized field propagator

Arbouet A.1, Mlayah A.1, Girard C.1, Colas des Francs G.2
1CEMES-CNRS, Université de Toulouse, 29 Rue Jeanne Marvig, 31055 TOULOUSE FRANCE EU, 2Laboratoire Interdisciplinaire Carnot de Bourgogne (ICB), UMR 6303 CNRS, Université de Bourgogne, 9 Avenue Savary, BP 47870, 21078 Dijon Cedex, France
arnaud.arbouet@cemes.fr

Stimulated by both instrumental (monochromators, detectors) and methodological (signal deconvolution and processing) advances, fast electron based spectroscopies have demonstrated their unique potential in probing surface plasmons (SP) of metallic nanostructures. Their nanometer spatial resolution and ability to probe so-called dark modes have given Electron Energy Loss Spectroscopy (EELS) and Cathodoluminescence spectroscopy (CL) a central role in experimental nano-optics. Today, these techniques are used to investigate nanostructures of increasing complexity in which the particle morphology, the substrate, or the interparticle interactions strongly influence their optical response[1]. Several examples of recent breakthroughs in combined electron/optical spectroscopy techniques such as Electron Energy Gain Spectroscopy demonstrated in ultrafast Transmission Electron Microscopes or surface plasmon three-dimensional imaging push forward the need and development of novel simulation techniques. In this context, we have developed a novel simulation technique allowing to describe thoroughly the interaction of fast electrons with metallic nanostructures. Building on the 3D Green Dyadic Method, our technique yields accurate predictions of the energy losses and CL photon emission consecutive to the interaction of a moving charge with a metallic nanostructure. It can be applied to nanostructures of arbitrary morphology, both penetrating and non-penetrating trajectories and rigorously takes into account the dielectric response of the substrate. Several examples will be presented which show an excellent agreement with recent experimental results. The influence of the substrate on the EELS spectra will be addressed. EELS spectra and maps (Fig. 1), CL spectra, maps and radiation patterns (Fig. 2) of several gold nanostructures from well-known textbook examples (nanoprisms, rods...) to more complex architectures (nanoporous films, particle aggregates, Fig. 3-4) will be presented. Finally, the potential of our technique will be illustrated on complex scenarii involving electron/photon interactions.


The authors acknowledge financial support from the European Union under the Seventh Framework Programme under a contract for an Integrated Infrastructure Initiative Reference 312483-ESTEEM2 and the National Research Agency under the program ANR HYNNA (ANR-10-BLAN-1016)

Fig. 1: EELS spectra computed at 9 nm from a gold nanoprism edge (edge length a = 950 nm, thickness t = 15 nm lying on a Si3N4 substrate (εsub = 3.9). The electron kinetic energy is 200 keV. Inset: EELS maps computed at 1 eV. The EELS probability is per electron and per unit energy.

Fig. 2: Above: CL intensity per electron per unit energy range at 1.42 eV induced by a 200 kV electron incident on a gold nanowire (length L=700 nm, diameter D = 50 nm) deposited on a substrate with εsub = 4. The collecting mirror has a numerical aperture NA = 0.8. Below: CL radiation pattern for an electron incident at the center of the nanowire.

Fig. 3: Randomly generated nanostructure composed of 29 gold spheres (diameter D = 9 nm). The particles are on a glass substrate (εsub = 2.25).

Fig. 4: Corresponding computed EELS probability maps for a 200 kV electron incident computed at 2 eV.

Type of presentation: Oral

IT-5-O-2070 Electron impact investigation of void plasmon ring resonators

Talebi N.1, Ögüt B.1, Sigle W.1, Vogelgesang R.2, van Aken P. A.1
1Max Planck Institute for Intelligent Systems, Heisenbergstr. 3, 70569 Stuttgart, Germany, 2University of Oldenburg, Oldenburg, Ammerländer Heerstr. 114-118, 26129, Germany
vanaken@is.mpg.de

Electron energy-loss spectroscopy (EELS) provides a fast and accurate way of investigating the optical density of local states (ODLS), especially plasmons. By detecting the amount of loss of the electron energy, one can study plasmonic modes. Furthermore, energy-filtered transmission electron microscopy (EFTEM) performed in the Zeiss SESAM microscope [1] is an efficient detection tool for mapping the optical modes in two spatial dimensions.

A specific domain of the possible ODLS excitable by the electrons is the localized plasmon resonance (LPR) excitation. LPR is attracting particular attention due to its possibility to introduce optical key elements such as waveguides [2] and resonators [3] below the diffraction limit of light. Due to the small spatial dimensions of these optical systems, investigating their behaviour in a wide energy band is challenging for optical spectroscopy techniques.

Here, using EFTEM and EELS, we investigate the possible modes of void hexamer and heptamer plasmon resonators, with respect to their symmetries and topologies. The hexamer nanocavity is composed of 6 holes with a diameter of 70 nm and rim-to-rim spacing of 30 nm, drilled into a 100 nm thick silver film. The heptamer resonator is composed of 7 holes with one hole located at the center, and the other 6 holes located along the circumference of a ring. Each hole has a diameter of 60 nm and a rim-to-rim spacing of 50 nm.

The proposed structures sustain similar symmetry; however they differ according to the number of holes and topology. In order to investigate the spatial distribution of LPR modes, a peak-finding algorithm [4] has been utilized. Four distinguished modes could experimentally be observed (Figure 1a,b). These LPR modes can be classified into those modes related to topology, such as toroidal plasmonic modes [5, 6], and those only related to symmetries, such as radially and azimuthally polarized modes [6]. In order to investigate the symmetries more systematically, FDTD calculations have been performed, which provides modal decomposition analysis (Figure 2) with selective excitation of different resonances, and hence yielding their eigenenergies and spatial distributions. In this presentation the concept of symmetry- and topology-related classification of LPR modes is thoroughly discussed.

References:

[1] C.T. Koch et al., Microscopy and Microanalysis 12 (2006) 506

[2] S. M. Raeis Zadeh Bajestani, M. Shahabadi, and N. Talebi, J. Opt. Soc. Am. B. 28 (2011) 937

[3] N Talebi, A Mahjoubfar, and M. Shahabadi, J. Opt. Soc. Am. B 25 (2008) 2116

[4] N Talebi et al. Langmuir 28 (2012), 8867

[5] B Ögüt et al. Nano Lett. 12 (2012) 5239

[6] N. Talebi et al. Appl. Phys. A, accepted for publication (2014)


The research leading to these results has received funding from the European Union Seventh Framework Program [FP/2007-2013] under grant agreement no 312483 (ESTEEM2).

Fig. 1: Peak maps obtained from the acquired EFTEM images at energy losses depicted in the figure for (a) a heptamer resonator and (c) a hexamer resonator. Acquired EELS spectra at the depicted impacts are shown in (b) for heptamer and (c) hexamer nanocavities.

Fig. 2: The eigenmodes simulated with 3D-FDTD which depict the spatial field distribution for the magnitude of the z-component of the electric field for (a) the heptamer nanocavity system at energy loss values of 2.1, 2.5, 3.0, 3.5, and 3.7 eV, and (b) the hexamer nanocavity structure at energy loss values of 1.8, 2.7, 2.9, 3.4, and 3.8 eV.

Type of presentation: Oral

IT-5-O-2097 Characterization of hybrid plasmonic waveguide by dispersion measurement

Saito H.1, Kurata H.1
1Institute for Chemical Research, Kyoto University
saito@eels.kuicr.kyoto-u.ac.jp

 Dielectric waveguides combined with plasmonic waveguides, so called hybrid waveguides [1], have a great potential for providing subwavelength confinement and long propagation length, leading to highly integrated photonic circuits. The simplest example of the hybrid waveguide structures is a three-layered film consists of a high-permittivity semiconductor layer separated from a metal substrate with a thin low-permittivity insulator gap. We performed dispersion measurements on Si/SiO2/Al films using angle-resolved EELS (AREELS) combined with a TEM [2].

 Figure 1(a) and (b) show AREELS patterns taken from the Si(157 nm)/SiO2(6 nm) film and the Si(157 nm)/SiO2(6 nm)/Al(35 nm) film, respectively. The dispersion curves of the first and second order Si waveguide modes are observed in Fig. 1(a), while the dispersion curves in Fig. 1(b) shift to the low energy side compared to those of the Si waveguide modes. These curves are the dispersion curves of the hybrid waveguide modes. According to the coupled-mode theory [3], to a first approximation, the hybrid waveguide mode can be described as a superposition of the Si waveguide mode and the surface plasmon-polariton (SPP) mode excited on the Al/SiO2 interface. The amplitude of the Si waveguide mode ASi is determined by the wave vector of the Si waveguide (kSi), the SPP (kSPP) and the hybrid waveguide (khyb) modes as follow [1],

|ASi|2 = (khyb-kSPP)/(2khyb-kSi-kSPP).    (1)

 Using the experimental dispersion plots (Fig. 2(a)) and the calculated dispersion relation of the SPP mode excited on the interface Al/SiO2, the square norm of ASi can be determined by equation (1). Figure 2(b) shows the resultant mode character depending on the energy of coupling. The hybrid waveguide mode with high energy has large component of the Si waveguide mode, so the electromagnetic energy of waveguide is mainly stored inside the Si layer, while it should be transferred from the Si layer to the Al/SiO2 interface with the decrease of energy because of the increase of the SPP component as shown in Fig. 2(b).

[1] R. F. Oulton et al. Nat. Photon. 2, 496 (2008).

[2] H. Saito et al. J. Appl. Phys. 113, 113509 (2013).

[3] A. W. Snyder and J. D. Love, Optical Waveguide Theory (Chapman and Hall, London, New York, 1983).


Fig. 1: Figure 1. AREELS patterns of (a) Si/SiO2 film and (b) Si/SiO2/Al film. The white curves and line are the calculated dispersion relations of light in Si bulk, SiO2 bulk and vacuum. The red curve is the dispersion relation of Čerenkov radiation in Si bulk.

Fig. 2: Figure 2. (a) The dispersion plots of the first order modes of the Si waveguide obtained from Si/SiO2 film (blue) and the hybrid waveguide obtained from Si/SiO2/Al film (red). (b) The energy dependence of the square norm of ASi calculated using the equation (1).

Type of presentation: Oral

IT-5-O-2167 Electro-optical characterization of single InGaN/GaN core-shell LEDs inside an SEM

Ledig J.1, Scholz G.1, Popp M.1, Steib F.1, Fahl A.1, Wang X.1, Hartmann J.1, Mandl M.1,2, Schimpke T.1,2, Strassburg M.2, Wehmann H. H.1, Waag A.1
1Institut für Halbleitertechnik, Technische Universität Braunschweig, Hans-Sommer-Str. 66, 38106 Braunschweig, Germany, 2OSRAM Opto Semiconductors GmbH, Leibnizstr. 4, 93055 Regensburg, Germany
j.ledig@tu-bs.de

Three dimensional light emitting diodes (LEDs) with a shell geometry around a columnar GaN core are supposed to have substantial advantages over conventional planar LEDs. The active area along the sidewalls of the GaN pillars can considerably be increased by high aspect ratios - leading to a lower current density inside the InGaN multi quantum well (MQW) at the same operation current per substrate area. Due to the 3-dimensional (3D) shape, the electrical and optical characterization of such device structures is a substantial problem because most of the conventional characterization techniques (e.g. Hall effect, capacitance/voltage) cannot be used with 3D geometries.
A nano-manipulator setup inside a scanning electron microscope (SEM) has been used in combination with a cathodoluminescence (CL) system to characterize the electro-optical properties by directly contacting single facets of the 3D structure. The investigated core-shell LEDs are grown by selective area metal organic vapor phase epitaxy on templates consisting of a patterned SiOx mask layer on an n-type GaN layer on 2” sapphire wafers. The light extraction of optical emission from a small region is related to the structure shape, this is estimated using ray tracing simulation and observed by spatially resolved and angle resolved microscope images inside the CL system.
Electron beam induced current (EBIC) images obtained on 3D structures contacted inside the SEM via the substrate and a probe tip clearly prove that a conjunct p-type shell is wrapped around the entire n-type column with an aspect ratio of about 5 forming a depletion region. By comparing spatially resolved CL and EBIC, the rate of charge carrier generation, trapping and separation in different regions of the structure are discussed.
We will present results of electroluminescence (EL) of MQW as well as defect related emission from single core-shell LED structures obtained at different injection currents. A wavelength shift of the MQW emission by 60 nm is observed along the structure height for both excitation methods (CL and EL), indicating a gradient of the indium incorporation caused by changing local growth conditions. The spectra are corrected with respect to the spectral sensitivity of the optical system - including the collection optics, monochromator and the CCD parallel detector.
In addition, metal contacts have been fabricated in order to get a defined contact area. By evaluating the contact area and the EL spectra we gain an insight to the internal efficiency of a single structure versus current density and the average spatially resolved extraction properties.


We thank Dr. Uwe Jahn for support regarding optical characterization. The financial support of the European commission (SMASH and GECCO) as well as the endorsement of the NTH and the JOMC are acknowledged.

Fig. 1: SE image of a core-shell LED structure on the cleaved growth template contacted by a probe tip on the sidewall at a height of 2.6 µm at an FOV =11.4 µm, EHT = 15 kV, tilt = 30°. EBIC image (right) at a reverse bias of VR = 7 V obtained by contacting a core-shell LED with a probe tip, the core is contacted via the n-type GaN buffer layer.

Fig. 2: Photograph of the CL-SEM chamber showing the arrangement of the parabolic mirror (1) for light collection, the micromanipulator (2) and sample (3) tilted by 30°. The sample is a contacted inside the focal point of the optical collection system by a tungsten probe tip attached to the micromanipulator.

Fig. 3: EL spectra of the core-shell LED shown above obtained by different injection currents through a tip contact placed on the sidewall at a height of 4.4 µm and CL spectrum (upper curve) obtained by electron probe excitation at the same height. The spectra are captured with a spectral FWHM of about 7.5 nm using a CCD parallel detector.

Type of presentation: Oral

IT-5-O-2181 Probing Near Fields of Nanoparticle Surface Plasmons with Electron Energy Loss Spectroscopy

Collins S. M.1, Nicoletti O.1, Ostasevicius T.1, Rossouw D.2, Duchamp M.3, Botton G. A.2, Midgley P. A.1
1Department of Materials Science and Metallurgy, University of Cambridge, Cambridge, UK, 2Department of Materials Science and Engineering, McMaster University, Hamilton, Canada, 3Ernst Ruska-Centre for Microscopy and Spectroscopy with Electrons, Forschungszentrum Jülich, Jülich, Germany
smc204@cam.ac.uk

Localized surface plasmon resonances (SPRs) of metal nanoparticles enable nanoscale manipulation of electromagnetic fields for a variety of sensing and light concentration applications. Electron energy loss spectroscopy in the scanning transmission electron microscope (STEM-EELS) is emerging as a key technique for near field characterization of SPRs. However, far field light and electron beam excitation produce distinct responses in plasmonic nanoparticles [1]. In recent work, three-dimensional imaging of nanocube resonances described modal responses consistent with light scattering near fields [2]. In this presentation, single nanoparticle resonances observed experimentally by STEM-EELS in lower symmetry systems (e.g., nanorods, right bipyramids) are compared using discrete dipole approximation (DDA) simulations for both light and electron excitation sources [3].

Silver nanorod monomers exhibit Fano-like resonances among longitudinal modes in far field light scattering studies [4]. Electron energy loss spectroscopy (EELS) experiments and EELS electrodynamics simulations, however, exhibit symmetric spectral line shapes. Experiments and simulations do display spatial amplitude modulation of the longitudinal nanorod mode, consistent with near field interference effects (Figure 1). In this presentation, key differences in the near field responses of silver nanorods to far field light and electron beam excitation will be examined (Figure 2). Interference effects among longitudinal nanorod modes will be discussed in terms of near field amplitude modulation as well as coupled oscillator modelling to explain experimental EELS mapping results as well as simulated EELS and cathodoluminescence signals.

[1] Collins, S. M.; Midgley, P. A. Phys. Rev. B 2013, 87, 235432.
[2] Nicoletti, O.; de la Peña, F.; Leary, R. K.; Holland, D. J.; Ducati, C.; Midgley, P. A. Nature 2013, 502, 80.
[3] Bigelow, N. W.; Vaschillo, A.; Iberi, V.; Camden, J. P.; Masiello, D. J. ACS Nano 2012, 6, 7497.
[4] López-Tejeira, F.; Paniagua-Domínguez, R.; Rodríguez-Oliveros, R.; Sánchez-Gil, J. A. New J. Phys. 2012, 14, 023035.


S.M.C. acknowledges the support of a Gates Cambridge Scholarship. This work has received funding from the European Research Council under the EU’s Seventh Framework Programme (FP7/2007-2013)/ERC grant agreement 291522-3DIMAGE and a contract for an Integrated Infrastructure Initiative (Reference 312483-ESTEEM2). G.A.B. is grateful to NSERC for a Discovery Grant supporting part of this work.

Fig. 1: STEM-EELS maps and line profiles (4 nm from rod side) of modal components m = 4 – 5 for a 540 nm long Ag nanorod on a 30 nm silicon nitride substrate (processed by non-negative matrix factorization). Extracted line profiles are compared with simulated line profiles. The electron trajectory is along the z-axis.

Fig. 2: (a)-(b) Phase analysis of light scattering and EELS responses of Ag nanorod simulated by DDA. The respective light absorption (Qabs) and light scattering (Qsca) efficiencies and EELS probability are plotted for comparison. (c)-(d) Net induced dipole moment along the axis of the rod (y¬-axis) calculated for light scattering and EELS.

Type of presentation: Oral

IT-5-O-2490 Single Eu atom M shell spectroscopy and X-ray fluorescence yield

G Tizei L. H.1, Nakanishi R.2, Kitaura R.2, Shinohara H.2, Suenaga K.1
1Nanotube Research Center, National Institute of Advanced Industrial Science and Technology (AIST), Tsukuba 305-8565, Japan, 2Department of Chemistry, Nagoya University, Nagoya 468-8602, Japan
Luiz.tizei@aist.go.jp

Europium (Eu) has a half filled 4f shell and two valences (2+, 3+). The transition between these is induced by pressure or oxidation. We report single atom EELS and EDX spectroscopy of Eu M shell. This shell is ideal for such experiments as its energy falls in a range where EELS and EDX are feasible. Atomically resolved EELS experiments were performed in the JEOL-CREST double corrected microscope operated at 60 kV. EDX-EELS experiments have been performed in an ARM-JEOL 200 operated at 60 kV, equipped with a Centurio-JEOL silicon drifted detector (SDD, 0.80 sr collection angle).
Eu double atomic chains confined inside carbon nanotubes (Figure 1a) have been studied. EELS analysis shows that all chains contain 2+ atoms. The crystal structure is related to that of bulk Eu hcp.
Firstly, Eu M edge EELS analysis (spectrum in Figure 1d) demonstrates the possibility of measuring the spectral signature of single Eu atoms and possibly their valence. HAADF (Figure 1b) and EELS (Figure 1c) intensity maps show the atomic positions and the measured signal.
To compare the absorption and emission of single atoms (Figure 2) EDX and EELS spectrum images have been acquired. Figure 2 shows the maps of the HAADF, EELS M edge and EDX M lines (Figure 2a-c). Atomic positions can only be distinguished in the first two due to the low signal level for the EDX. Profiles along the second atomic pair (arrow in Figure 2a) are shown in Figure 2d. The M and the L emission lines and the M edge absorption signal have been measured. The M EDX and M EELS spectra for one atom are shown in Figure 2e and 2f (equivalent to 12.5 s exposure), respectively. An average EDX spectrum (198 s exposure) is shown in Figure 2g.
We have estimated the total X-ray emission and absorption events (red square on Figure 2b). The absorption signal has been corrected for the CCD efficiency and the finite convergence and collection angles. Events not counted due to the finite energy integration window (100 eV) have been estimated using power laws (largest source of uncertainty). The X-ray signal was corrected by a geometric factor due to the detector’s solid angle. The creation of M-shell vacancies due to L shell transitions was estimated from the EDX signal. Coster-Kronig M sub-shell transitions were not considered. Estimated fluorescence yield lies between 0.02 and 0.03 (the theoretical value is 0.0136). The uncertainty stems from the necessity to extrapolate the tail of the M edge.


This work is partially supported by a JST Research Acceleration programme. The authors would like to thank Niimi Yoshiko for her assistance during initial experiments.

Fig. 1: a) HAADF image of a double Eu chain. b-c) HAADF and integrated M intensity (100 eV window) of a double Eu chain. d) Spectrum from one atom integrated on the red square in b (total exposure time 9x50 = 450 ms). The inset shows the signal after background subtraction.

Fig. 2: a-c) HAADF, M edge EELS and M line EDX for a double Eu chain, respectively. d) Profiles of the Eu M and L EDX, M edge EELS and HAADF signals along the arrow in a). e-g) Single atom EDX and EELS spectra, respectively. f) Average EDX signal in the spectrum image. The exposure time was 198s. Not all peaks are marked but all are identified.

Type of presentation: Oral

IT-5-O-2552 Imaging and X-Ray Microanalysis at the Nanoscale with a Cold-Field Emission Scanning Electron Microscope

Gauvin R.1, Brodusch N.1, Demers H.1, Woo P.2
1Department of Mining and Materials Engineering, McGill University, Montreal, Quebec, Canada, 2Hitachi High-Technologies Canada Inc., Toronto, Canada
raynald.gauvin@mcgill.ca

The scanning electron microscope (SEM) was primary developed for imaging applications. With the introduction of the Si(Li) energy dispersive spectrometer (EDS), simultaneous imaging and x-ray microanalysis became possible. However, long working distance and high current were needed because the position and small solid angle of the EDS detector. SEM was initially and is still optimized for imaging applications, where the high spatial resolution is generally obtained at short working distance. This problem is still relevant today and unfortunately x-ray microanalysis is never performed in the best imaging conditions, i.e., not with the smallest probe size. With the introduction of an annular silicon drift detector (SDD) system, scanning electron microscopy is facing a revolution. This detector is inserted below the objective lens which gives a higher solid angle (up to 1.2 sr). In consequence, a lower working distance and probe current can be used. An improved spatial resolution becomes possible during x-ray microanalysis. At this point, the time required for x-ray imaging will be of the same order as for the atomic number contrast images achieved through backscattered electrons (BSE) imaging.

Carbon nanotubes (CNTs) decorated with platinum (Pt) nanoparticles are often used to evaluate the spatial resolution of cold-field emission scanning electron microscope (CFE-SEM). Figure 1 shows an example of high spatial resolution imaging and x-ray microanalysis of CNTs at low accelerating voltage (2.5 kV). A resolution of 19 nm and 24 nm were measured with SMART-J on the SE micrograph and the Pt x-ray map, respectively. Figure 2 shows another example of high spatial resolution imaging x-ray map obtained with an annular SDD of CNTs with low voltage scanning transmitted electron microscope (LVSTEM) mode at 20 kV. The dark-field micrograph had a spatial resolution of 6.5 nm and the Pt x-ray map had a spatial resolution of 8.9 nm. Currently, this system is limited to accelerating voltage below 20 kV and the shortest working distance is around 10 mm, which is shorter than the one used with a conventional SDD (15 mm on our system).

With this x-ray detector installed on a HITACHI SU-8230 cold-field emission scanning electron microscope, quantitative x-ray microanalysis with high spatial resolution at low beam energy and low current becomes possible with the possibility of using the various different type of imaging at the same time. Also, since the count rate can be as high as 1,500 kcps with our system, which lowers significantly the detection limit of elements as well as the minimum feature sizes of different phases that can be distinguished.


Fig. 1: Secondary electron micrograph of CNTs decorated with Pt nanoparticles was acquired at an accelerating voltage of 2.5 kV and a working distance of 9.4 mm. The Pt X-ray map was acquired with an annular silicon drift detector. The map acquisition time was 1433 s with a count rate of 81 kcps.

Fig. 2: Dark field micrograph of CNTs decorated with Pt nanoparticles was acquired in LV-STEM mode. The Pt X-ray map was acquired with an annular silicon drift detector. An accelerating voltage of 20 kV and a working distance of 10.5 mm were used. The map acquisition time was 412 s with a count rate of 7 kcps.

Type of presentation: Oral

IT-5-O-2673 Probing the electronic structure of functional oxides with high energy resolution EELS

Bugnet M.1, Rossouw D.1,2, Liu H.1, Radtke G.3, Botton G. A.1
1Materials Science and Engineering, McMaster University, 1280 Main Street West, Hamilton, Ontario L8S 4L7, Canada, 2Department of Materials Science and Metallurgy, University of Cambridge, Cambridge, United-Kingdom., 3IMPMC-CNRS, UMR 7590, Université Pierre et Marie Curie-Paris 6, Campus Jussieu, 4 place Jussieu, F-75252 Paris Cedex 05, France
bugnetm@mcmaster.ca

The development of monochromators in the transmission electron microscope (TEM) has allowed improvements in energy resolution comparable to what is currently achieved in x-ray absorption spectroscopy. High-resolution electron energy loss spectroscopy (HREELS) in the TEM provides an invaluable tool to probe subtle changes in the electronic structure of materials at the nanoscale. Selected examples highlighting how HREELS has provided insight into the electronic properties of functional oxides will be persented. The perovskite BaTiO3 is of particular interest because of its intrinsic ferroelectricity at room temperature, and its applications as a piezoelectric material and in capacitors. The ferroelectricity of BaTiO3 at room temperature arises from the off-center position of the Ti atoms in the tetragonal lattice. From a fundamental point of view, the understanding of this structural characteristic, based on the investigation of BaTiO3 complex electronic structure, is essential [1]. In the present study [2], the O 1s excitation is probed by HREELS, and the interpretation of the spectral features is performed with ab initio calculations. The effect of the core-hole potential is investigated, and the correlation between its effect and the geometry of the excited atomic site, i.e. the relative position of the two nearest Ti atoms with respect to the excited O atom, is shown (Figure 1). The link between the core-hole effect and the off-center position of the Ti atom appears as a broadening in the near edge structures, which is resolved with monochromated EELS. This broadening effect is highlighted by probing the O 1s excitation during the phase transitions between the low and high temperature phases of BaTiO3. The effects of continuous light illumination on the structural and electronic modifications of TiO2, a prospective material used for photocatalysis and water splitting, will also be shown. By using a recently-built in situ laser-illumination setup in the TEM [3], we explore the exposure of titania to intense light irradiation. The electronic structure modifications are probed by HREELS and are interpreted in terms of local reversible changes in the material. Finally, HREELS was used to probe the valence changes upon cycling of Li(Mn,Co,Ni)O2 battery cathode materials. Using scanning transmission electron microscopy combined with HREELS, it is shown that valence maps provide exquisite spectroscopic information on local changes from the charge and discharge process in battery materials.

[1] B. Zalar et al. Physical Review B 71, 064107 (2005)

[2] M. Bugnet et al. Physical Review B 88, 201107(R) (2013)

[3] D. Rossouw et al. Physical Review B 87, 125403 (2013)


The Authors are grateful to NSERC for financial support. The experiments were carried out at the Canadian Centre for Electron Microscopy, a national facility supported by NSERC and McMaster University.

Fig. 1: (a) Evidence of the wider fine structure at low energy in the O-K edge for tetragonal BaTiO3, as compared to cubic SrTiO3. (b) Experimental O-K edge and calculated contributions of the two independent O positions in tetragonal BaTiO3 (O1 in red dashed line, O2 in black solid line). (c) Illustration of BaTiO3 off-centre position of Ti.

Fig. 2: The valence state of the transition metals in pristine Li(Mn,Co,Ni)O2 battery cathode material probed by HREELS: (a) Mn4+, Co3+, and (b) Ni2+. The evolution of the electronic structure of Ni upon charging is highlighted in (b), indicating a change in valence state.

Type of presentation: Oral

IT-5-O-2679 Challenges and Opportunities in Materials Science with Next Generation Monochromated EELS

Crozier P. A.1, Zhu J.1, Aoki T.2, Rez P.3, Bowman W. J.1, Carpenter R. W.2, Krivanek O. L.4, Dellby N.4, Lovejoy T. C.4, Egerton R. F.5
1School for Engineering of Matter, Transport and Energy, Arizona State University, Tempe, AZ 85287, USA , 2LeRoy Erying Center for Solid State Science, Arizona State University, Tempe, AZ 85287, USA , 3Department of Physics, ASU, Tempe, AZ 85287, USA, 4Nion Co., 1102 8th St, Kirkland, WA 98033, USA, 5Department of Physics, University of Alberta, Edmonton T6G 2E1, Canada
crozier@asu.edu

The development of monochromated scanning transmission electron microscopes (STEM) offering energy resolutions of better than 20 meV and electron probes of 0.1 nm in size provides a new tool for materials characterization. Unique opportunities opened by access to ultra-high energy resolution low loss EELS include determination of optical properties in the IR, bandgap mapping, detection of defect interband states and localized vibrational spectroscopy. At ASU we are currently applying ultra-high energy resolution low-loss EELS to a variety of materials that are important in fields such as energy, environmental science and information technology. Here we show representative initial results acquired on a newly installed Nion UltraSTEM equipped with a probe corrector and monochromator [1].

The optical properties of carbonaceous atmospheric aerosols are an important contributor to radiative forcing for climate change. By applying Kramers-Kronig techniques to energy-loss spectra acquired from the Nion, Figure 1 shows that the refractive index can be determined out to photon wavelengths of 2500 nm, thus covering most of the incoming solar spectrum [2].

Local measurement of bandgaps and states within the gap is of great importance for opto-electronic materials. Figure 2 shows the low-loss spectra from ceria (CeO2) and a ceria co-doped with Gd and Pr (Ce0.85Gd0.11Pr0.04O2-δ). From EELS, local bandgaps were about 3.5 eV and in some regions additional peaks were detected within the bandgap (Figure 2b). Interestingly, all the ceria based samples showed significant uniform intensity within the bandgap which will be discussed in terms of Cerenkov radiation, defects, and surface layers.

At lower energy transfers, localized phonon spectroscopy becomes possible. We have been able to identify vibrational peaks in a variety of compounds like SiO2 which match the Raman spectrum [1]. Figure 3 shows two regions of the low-loss spectrum from TiH2. The peak at 150 meV is prominent under aloof beam conditions.

Ultrahigh energy resolution EELS is an exciting new tool for characterization of materials. However, to realize its full potential, considerable experimental and theoretical work must be undertaken to develop a fundamental understanding of this form of EELS.

[1] O.L. Krivanek et al, these proceedings (2014)

[2] J. Zhu et al, these proceedings (2014)


The authors acknowledge support of the NIST 60NANB10D022, NSF Graduate Research Fellowship Program (DGE-1311230), NSF DMR 1308085,NSF MRI-R2 959905 and DOE DE-SC0004954. The authors acknowledge the use of facilities in the John M. Cowley Center for High Resolution Microscopy at Arizona State University.

Fig. 1: (a) EELS from two forms of carbonaceous particles. (b) Complex refractive index derived from EELS covering photon wavelength range 200-2500 nm.

Fig. 2: Low-loss spectra from a) CeO2, b) Ce0.85Gd0.11Pr0.04O2-δ and c) hexagonal BN.

Fig. 3: EELS from TiH2 showing a) wide energy range on sample and b) very low energy-loss region in aloof beam mode (~5 nm off sample).

Type of presentation: Oral

IT-5-O-2752 FIB-SEM Instrument with Integrated Raman Spectroscopy, Scanning Probe Microscopy and Secondary Ion Mass Spectroscopy

Jiruše J.1, Haničinec M.1, Havelka M.1, Hollricher O.2, Schaff O.3, Oestlund F.4, Whitby J.5, Michler J.5
1TESCAN Brno, s.r.o, Brno, Czech Republic , 2WITec GmbH, Ulm, Germany, 3SPECS Surface Nano Analysis GmbH, Berlin, Germany, 4TOFWERK AG, Thun, Switzerland, 5EMPA - Materials Science & Technology, Thun, Switzerland
jaroslav.jiruse@tescan.cz

Integration of a number of techniques in a single tool gives the possibility of correlating multiple measurements and analyses of the same sample area, all made in-situ. A multifunctional tool comprising an SEM-FIB, an SPM and a TOF-SIMS has been presented recently [1]. Newly, a Confocal Raman Microscope (CRM) has been added to yield information about molecular composition and chemical bonds. The CRM image complements the high resolution SEM image, topographic image from SPM, chemical map from TOF-SIMS and sample modification by FIB. Fig. 1 shows the arrangement of the presented apparatus.

State-of-the-art Raman analyzers in SEMs use a parabolic mirror for focusing and lateral resolution is usually no better than 2-5 µm. The presented system provides a resolution of 360 nm by integrating a full confocal light microscope. The important property is the capability of Raman imaging. When a spectrum from a single point is acquired, one can never be sure if the position calibration is correct. Fig. 2 shows overlaid Raman and SEM micrographs of diorite sample. Besides lateral scanning, vertical movement is supported, which allows non-destructive 3D tomography.

Integration is possible with two alternative electron optical columns, each with a Schottky field-emission gun: the LYRA with a conventional objective lens or the GAIA with an immersion lens. The immersion lens column [2] is recommended for non-conductive or fragile samples, because it offers better resolution at low energies (1 nm at 15 kV and 1.4 nm at 1 kV).

The FIB is used to modify the sample and it also enables 3D tomography techniques by sequential FIB slicing followed by imaging to create 3D datasets with analytical information such as elemental composition, crystallographic information, etc.

The FIB also acts as a primary ion beam for the TOF-SIMS analysis. It allows 2D as well as 3D spectral maps, carrying elemental, isotopic and chemical information about the investigated sample. Fig. 3 shows a TOF-SIMS 3D tomography of sodium contamination on solar cell sample. Lateral resolution of TOF-SIMS maps can be better than 50 nm.

The integrated Scanning Probe Microscope (SPM) supports work in STM and AFM modes. Its compact design allows it to sit on the SEM stage. Simultaneous use of SPM, SEM and FIB enables a true depth calibration of TOF-SIMS depth profile as well as the calibration of 3D tomography techniques. The SPM head is designed for a depth resolution of 0.1 nm and an imaging speed of up to 20 s per image. Fig. 4 shows the AFM topography map of gold particles on carbon and corresponding SEM micrograph.

References:

[1] J Jiruše et al, Microsc. Microanal. 18 (Suppl. 2) (2012) p. 638.

[2] J Jiruše et al, Microsc. Microanal. 19 (Suppl. 2) (2013) p. 1302.


The research leading to these results has received funding from the European Union Seventh Framework Program [FP7/2007-2013] under grant agreement No. 280566, project UnivSEM.

Fig. 1: Geometrical arrangement of the presented apparatus.

Fig. 2: Raman (in color) and SEM (in gray) overlaid micrographs of diorite sample. Different colors correspond to various phases in the sample.

Fig. 3: TOF-SIMS 3D tomography of sodium contamination on solar cell sample.

Fig. 4: AFM topography map of gold particles on carbon and corresponding SEM image. Field of view is 1 µm.

Type of presentation: Oral

IT-5-O-2917 Full Optical Properties of Carbonaceous Aerosols by Very High Energy Resolution Electron Energy-loss Spectroscopy

Zhu J.1, 2, Crozier P. A.1, Aoki T.2, Anderson J. R.1
1School for Engineering of Matter, Transport and Energy, Arizona State University, Tempe, AZ, USA, 2LeRoy Eyring Center for Solid State Science, Arizona State University, Tempe, AZ, USA
jiangtao.zhu@asu.edu

Carbonaceous aerosols have a strong impact on the global climate by direct radiative forcing via light absorption and scattering, and/or indirect radiative forcing via influencing cloud formation. It is critical to determine their optical properties to understand their contribution to direct radiative forcing. It is also important to understand their local chemical composition which would affect their role in cloud dynamics. By employing monochromated electron energy loss spectroscopy in a newly installed Nion UltraSTEM 100 at ASU with a sub 20 meV energy resolution, we can now determine the optical properties of carbonaceous aerosols over the full range of incoming solar radiation 200-2500 nm including the infrared. In addition, the compositional variation in aerosol spherules can also be studied.

Depending on the sources and the combustion conditions, different types of carbonaceous aerosols [1], can be present in the atmosphere. Here two types of aerosols, graphitic and amorphous carbon collected from East Asia, were investigated. Low loss spectra of these two types of aerosols were collected from Nion UltraSTEM 100 at 60 kV with an energy dispersion of 10 meV/channel giving a zero loss peak (ZLP) full width at half maximum (FWHM) of 40 meV. As shown in Fig. 1a, the EELS of graphitic and amorphous carbons are different, which is related to their different microstructure and chemical bonding. We found that the standard thin film formulation of Kramers-Kronig analysis can be employed to make accurate determination of the dielectric function for carbonaceous particles down to about 40 nm in size [2]. Figure 1b and c show the complex refractive indices (n-ik) of graphitic carbon spherules and amorphous carbon spheres over the photon wavelength range of 200-2500 nm. The absorption in the infrared range is obtained although it is smaller than that in the UV and visible range. The variances in the refractive indices of different particles were related to their variances in composition. We also found that on the surface of both graphitic spherule and amorphous carbon, there is a several nanometer thick layer rich in silicon and oxygen (Fig. 2). The nano-size coating could modify the chemical interactions of the carbonaceous aerosols, for example, by changing their ice nucleation properties.

References

[1] J. Zhu, P.A. Crozier, J.R. Anderson, Atmos Chem Phys. 13 (2013), p. 6359.

[2] J. Zhu, P.A. Crozier, P. Ercius, J.R. Anderson, accepted by Microsc Microanal.


The authors acknowledge support of the National Institute of Standards and Technology under Award 60NANB10D022 and the use of facilities in John M. Cowley Center for High Resolution Electron Microscopy at Arizona State University.

Fig. 1: (a) Low loss EELS of graphitic and amorphous carbon (a) and their refractive index (n-ik) in the range of 200-2500 nm in optical wavelength (b, c), respectively.

Fig. 2: (a) The angular dark field image of a graphitic carbonaceous aerosols.(b) Intensity distributions of C, O, S, K, Ca, Ga and Si along the line in the image (a).

Type of presentation: Oral

IT-5-O-2999 Direct observations of local electronic states in a quasicrystal by STEM-EELS

Seki T.1, Abe E.1
1University of Tokyo, Tokyo, Japan
seki@stem.t.u-tokyo.ac.jp

    Quasicrystals (QCs) have long-range ordered complex structure without periodicities. Stability of QCs has been discussed in terms of energetic gains in electron systems, because most QCs reveal pseudogaps in their density of states around Fermi level. In fact, many QCs have been discovered by tuning valence electron density based on Hume-Rothery rule. Therefore, understanding electronic structures in QCs may provide an important clue for their stabilization mechanism. Generally, it has been frequently discussed based on an interaction between Fermi surface and Brillouin zone boundary within the framework of nearly free electron model, providing an underlying physics of a Hume-Rothery’s empirical criteria. However, the electronic structures of QCs have not yet been fully understood, particularly being in microscopic-macroscopic relations. In the present work, we investigate local electronic states in Al-based QCs using electron energy loss spectroscopy (EELS) combined with scanning transmission electron microscopy (STEM).
    The AlCuIr decagonal phase was used for the present study [1]. A cluster with a diameter of ~2 nm emerges as a building unit for AlCuIr decagonal phase (Fig. 1, 2). Principal components analysis clearly showed up the atomic-site dependence of plasmon loss spectra in a two-dimensional map correlated with the cluster arrangement. Qualitatively, there seems to be certain correlations between the plasmon peaks and the core-loss edges, Al L1, Ir O23, Ir N67 and Cu L23, all of which reveal different behaviours at the cluster centers and the edges (Fig. 3). All results indicate the cluster centers have metallic states, while the cluster edges have covalent states. First-principles calculations confirm the unusual electronic state. We analyse a distribution of covalent electrons by Fourier transformation of electron localization function. The distribution seems like a 10-fold charge density wave with Fermi wave length. It suggests that the Hume-Rothery mechanism play a key role even when the hybridization effect mainly contributes to pseudogap formation. Along with a context of orbital hybridization, the covalent electrons might reduce their energy by mimicking 10-fold charge density wave. On the other hand, the metallic regions at the cluster centers may have no contributions to the Hume-Rothery mechanism, since there are no distinguished peaks appeared along the 10-fold charge density wave at the relevant regions.

[1] 1. P. Kuczera, J. Wolny and W. Steurer, Acta Crystallographica B68 (2012), 578.


This work was conducted in Research Hub for Advanced Nano Characterization, The University of Tokyo, supported by the Ministry of Education, Culture, Sports, Science and Technology (MEXT), Japan. It is acknowledged that T. Seki is a research fellow of Japan Society for the Promotion of Science.

Fig. 1: HAADF-STEM image of AlCuIr decagonal quasicrystal. Yellow circles indicate clusters with a diameter of ~2 nm. The image shows only Cu and Ir atomic columns.

Fig. 2: Structure model of the cluster. Blue, green and red circles correspond to Al, Cu and Ir atoms, respectively.

Fig. 3: EELS spectra consisting Ir-O23, Ir-N67 and Al L23 obtained from the cluster centers and edges. Intensity ratio of Ir-O2 and N67 to Ir-O3 depends on atomic sites.

Type of presentation: Oral

IT-5-O-3107 Luminescence of structural defects in opto-electronic materials studied by CL-STEM

Perillat-Merceroz G.1, Alexander D. T.2, Zamani R. R.3, Arbiol J.3, Stadelmann P.2, Grandjean N.1, Hébert C.2
1Laboratory of Advanced Semiconductors for Photonics and Electronics (LASPE), École Polytechnique Fédérale de Lausanne (EPFL), Lausanne, Switzerland, 2Interdisciplinary Centre for Electron Microscopy (CIME), École Polytechnique Fédérale de Lausanne (EPFL), Lausanne, Switzerland, 3Institut de Ciència de Materials de Barcelona (ICMAB-CSIC), Barcelona, CAT, Spain
duncan.alexander@epfl.ch

The cathodoluminescence (CL) signal from opto-electronic materials is a sensitive function of their structural properties. CL mapping in scanning TEM (STEM) mode offers a number of advantages for studying this relationship. The interaction volume of the electron beam has a diameter of a few nm within the TEM lamella or sample, giving the potential for high spatial resolution CL studies (e.g. [1]). Further, the CL signal can be acquired simultaneously with defect contrast images taken with bright-field (BF) or low angle dark-field (LDF) detectors, and atomic number contrast images from a high angle annular dark-field (HAADF) detector. This allows direct correlation of luminescence to structure and defects.
Here we present studies using CL-STEM to probe the effects of structural defects on optical properties. Data are taken using a JEOL 2200FS equipped with the Gatan XiClone CL spectrometer, at 80 keV beam energy to increase CL signal and reduce beam damage. The instrument has BF, LDF, and HAADF STEM detectors; sample holders allow specimen temperatures from 5 K to room temperature. Para-CL spectrum image data are acquired on a liquid nitrogen cooled CCD camera using Gatan Digiscan, typically using a 300 lines/mm grating and sub-pixel scanning for the simultaneously acquired STEM images.
To optimize experimental conditions and to interpret data it is important to characterize the sample-instrument system. Fig. 1 shows the effects of specimen thickness and temperature on the CL signal of a GaN epilayer prepared for TEM by mechanical wedge polishing. Specimen cooling increases CL signal intensity and also reduces non-radiative recombination at the polished surfaces, so giving luminescence even for 30 nm thick material. Fig. 2 then shows threading dislocations in an InGaN quantum well (QW) on a GaN substrate. Some dislocations demonstrate modulations in wavelength that could correlate to In enrichment and depletion in the strain fields around their core [2]. Fig. 3 instead shows data taken from a GaN nanowire that contains stacking faults (SFs) on the basal plane [3]. The spatial correlation of SFs to the near-bandgap (NBE) and sub-bandgap (SBE) emission is complex, demanding further investigation, while midgap states generate a diffuse background. The observation of Fabry-Pérot resonator effects for CL in the TEM lamella will also be discussed. Together these results illustrate the great potential of CL-STEM for investigating the structure-optical domain.
[1] Zagonel et al., Nano Lett. 11, 2011, 568
[2] Mouti et al., PRB 83, 2011, 195309
[3] Schuster et al., Nano Lett. 12, 2012, 2199


The Competence Centre for Materials Science and Technology (CCMX) is acknowledged for CL funding, and Anas Mouti of ORNL for earlier works with the CL-STEM system.

Fig. 1: CL intensity of the GaN peak as a function of thickness at different temperatures. Extracted from CL line scans taken on a GaN epilayer sample with a wedge shape. Thicknesses estimated by electron energy-loss spectroscopy.

Fig. 2: (a) BF STEM image of a single InGaN QW on GaN substrate in plan view with threading dislocations coming to the surface. (b) Map of the wavelength of the QW peak (colour scale shown below). A butterfly shape is visible around some of the dislocations with emission blue-shifted on one side and red-shifted on the other.

Fig. 3: (a) BF STEM image and (b) RGB CL map of different emission bands for a GaN nanowire (red: NBE; green: SBE; blue: midgap states). While the planar defects in the nanowire have an influence on CL emission, there is not a direct spatial correlation between the two.

Type of presentation: Oral

IT-5-O-3165 3D EDX in SEM and TEM: common problems and common solutions

Burdet P.1, Croxall S. A.1, de la Peña F.1, Rossouw D.1, Midgley P. A.1
1Department of Materials Science & Metallurgy, University of Cambridge, Cambridge, UK
pb565@cam.ac.uk

Energy dispersive X-ray spectrometry (EDS) has recently seen a step forward with the introduction of silicon drift (SDD) detectors for both SEM and TEM. The improvements in speed and detector efficiency have allowed EDS, a traditionally 2D technique, to be extended to 3D. For SEM, focused ion beam (FIB) ‘slice-and-view’ methods can be used. For TEM, tilt series of 2D EDS maps are recorded for reconstruction by back-projection. 3D EDS mapping in both the SEM and TEM faces acquisition and processing challenges, such as partial detector shadowing, detection of spurious X-ray and high level of noise. In this paper we explore which problems are common to both techniques as well as proposing common solutions. To compare the two techniques, samples of Ni-based superalloy are investigated with complex structures ranging from microns to nanometers (see figure 1) and containing more than 10 elements (see figure 2).
For 3D experiments, the acquisition time per spectrum is often reduced, to cover a large volume or to acquire a sufficient number of tilt images, in a reasonable time. The set of data contains millions of noisy spectra characterizing only a limited set of chemical phases. This is a favorable case for a multivariate statistical approach, such as principal component analysis (PCA), and, as shown in figure 3, an important reduction of noise can be obtained with this technique. However, such statistical approaches need to be used with care, especially with data containing few counts per channel. Alternative PCA algorithms and pre-processing methods will be explored. Before quantification, the X-ray line intensities are extracted from each EDS spectrum. The involved processing steps are similar for SEM and TEM. Due to the relatively low energy resolution of the EDS detector, X-ray lines often overlap, as observed in figure 2 for Hf Mα and Ta Mα. Moreover, the background needs to be subtracted. Given the high level of noise in SEM and TEM datasets, accurate intensity extraction is challenging. Different processing strategies based on curve fitting will be discussed.
For 3D EDS in SEM and TEM, other signals may be recorded simultaneously, such as secondary electron (SE) for SEM or energy loss spectra for TEM. This opens the way for new processing methods benefiting from complementary signals. For instance, SE images can be used to improve the spatial resolution of the segmentation obtained with the EDS map [1]. In order to facilitate the interactive data analysis of these complex multi-dimensional datasets, HyperSpy [2] a free, open-source and open-development software package, has been extended to EDS data for SEM and TEM.
1. P. Burdet, J. Vannod, A. Hessler-Wyser, M. Rappaz, M. Cantoni, Acta Materialia 61 (2013) 3090–3098
2. http://hyperspy.org


The research leading to these results has received funding from the European Research Council under the European Union's Seventh Framework Programme (FP7/2007-2013) /ERC grant agreement 291522-3DIMAGE. SAC and PAM acknowledge financial support from Rolls-Royce plc.

Fig. 1: 3D SEM-EDS reconstruction of the sample of Ni-based superalloy. The green volume shows the Ni rich γ’ phase, red show the Hf rich phase and blue the Ta rich phase. The bounding box measures 12.8 µm x 11.2 µm x 6.3 µm.

Fig. 2: Characteristic SEM-EDS spectrum acquired from a sample of Ni-based superalloy. The main lines excited at 15 kV are indicated. The inset shows a magnified picture of the low energy lines.

Fig. 3: Denoising a SEM-EDS spectrum with PCA. A running sum is used prior to the PCA.

Type of presentation: Oral

IT-5-O-3209 Oxygen Vacancies at Grain Boundaries in Doubly-Doped Ceria Determined using EELS

Bowman W. J.1, Zhu J.1, Hussaini Z.1, Crozier A. P.1
1Arizona State University
wjbowman@asu.edu

 In oxygen conducting ceramics like CeO2, O2- diffusion occurs via thermally-activated hopping through vacancies whose concentration can be modulated by doping with aliovalent cations such as Gd3+ or Pr3+. Sluggish ionic conductivity in these polycrystalline electrolytes has been attributed in part to highly resistive grain boundaries (GBs) which degrade total ionic conductivity. Here we use a combination of electrochemical impedance spectroscopy (EIS) and electron energy-loss spectroscopy (EELS) to characterize the electrical conductivity and vacancy concentration of GBs in Gd-doped CeO2 also containing Pr or Co.
Gd-doped (GDC), Gd/Pr doubly-doped (GPDC) and a series of Gd/Co doubly-doped CeO2 ceramics were prepared using a spray drying approach together with traditional ceramic processing techniques. EIS was performed using a Gamry Reference 3000 potentiostat, and EELS was performed using a Nion UltraSTEM100. To facilitate interpretation of the EELS data, FEFF codes [1] were employed to simulate the CeO2 O K-edge spectra as a function of oxygen vacancy concentration.
 Fig. 1a shows a simulation of the O K-edge in CeO2 as a function of O2- vacancy concentration. In this model, the vacancies were randomly distributed in the fluorite structure. These results indicate a decrease in the first peak in the O-K edge fine structure with increasing vacancy concentration. Figs. 1b and 1c show experimental Ce M45 and O-K near edge structure acquired at the edge and center of a CeO2 particle. The Ce white line intensity switch is characteristic of Ce4+ reduction to Ce3+, and in this case is accompanied by a drop in the first O K-edge peak similar to that in the calculated spectra (1a).
 Conductivity data (fig. 2a) shows that the GB conductivity is an order of magnitude higher in GPDC compared to GDC. Figs. 2b and 2c show Ce white lines and O K-edge fine structure acquired at a GB and grain interior in GPDC. In this case the drop in the first O K-edge peak is not accompanied by a Ce white line intensity switch.
 Here we probe and correlate the cation distribution, oxidation state and the O K-edge fine structure to elucidate the vacancy environment at GBs in various doped CeO2 electrolytes. These measurements coupled with macroscopic characterization of GB conductivity will be used to relate atomic-level GB structure and chemistry with bulk electrical properties. We also aim to refine our FEFF calculations to improve the quantitative robustness of our experimental approach to determining the distribution of O2- vacancies near GBs.

References
[1] Rehr, J.J, et al., Phy. Chem. Chem. Phy., 2010 12 5503-5513


We thank Kevin Jorrisen for FEFF help, NSF GRFP-1311230 & DMR-1308085, ASU NASA Space Grant & ASU Cowley HREM Center

Fig. 1: (a) Calculated EELS O K-edge near-edge structure as a function of O2- vacancy in CeO2. The intensity of the first peak drops with increasing vacancy concentration. (b & c) Experimental Ce M45 white lines and O K-edge near edge structure acquired at the edge and center of a CeO2 particle.

Fig. 2: (a) Conductivity data showing the much higher GPDC grain boundary conductivity, σgb (◊). (b & c) Experimental Ce M45 white lines and O K-edge near edge structure acquired at a grain boundary and grain interior in GPDC.

Type of presentation: Oral

IT-5-O-3427 Elastic scattering of atomic-size electron probes carrying orbital angular momentum in aberration-corrected scanning transmission electron microscopy

Idrobo J. C.1
1Oak Ridge National Laboratory
jidrobo@gmail.com

The pioneering work by Uchida and Tonomura in 2010 [1] showed that electron beams carrying orbital angular momentum (OAM) can be produced in a transmission electron microscope. Since then, there has been a large interest in the microscopy community to produce atomic-size electron probes carrying OAM [2,3]. The interest arises because using those probes, in principle, one could study magnetic dichroism at the atomic scale through electron energy-loss spectroscopy (EELS) in aberration-corrected scanning transmission electron microscopy (STEM) [2].

In this work, we will present calculations that show how an atomic-size electron probe carrying OAM (vortex probe) channels through the sample and how its OAM character is affected by channeling. We will discuss the reasons why STEM images using vortex probes seem to show lower intensity contrast than images obtained with conventional aberration-corrected probes (as the example illustrated in Figure1). The STEM images simulations were obtained with a multislice algorithm scheme, using a recently developed code in Python (pySTEM) at Oak Ridge National Laboratory. The code calculates electron probes (up to C7 aberrations) with OAM implemented following the electron optics setup outlined in Refs. 4 and 5.

References:

[1] M. Uchida and A. Tonomura, Nature 464, 737 (2010).
[2] J. Verbeeck et al., Nature 467, 301 (2010).
[3] B. McMorran et al., Science 331, 192 (2011).
[4] J.C. Idrobo & S.J. Pennycook, J. of Electron Micros. 60, 295 (2011).
[5] O.L. Krivanek, et al., Micros. Microanal. in press (2014).


This research was supported by ORNL’s Center for Nanophase Materials Sciences (CNMS), which is sponsored by the Scientific User Facilities Division, Office of Basic Energy Sciences, U.S. Department of Energy (JCI).

Fig. 1: (Left) Simulated ADF images of monolayer MoS2 with an electron probe with and without orbital angular momentum (OAM).  Simulations done at 100 kV, a converge semi-angle of 30 mrad, and ADF collection semi-angles of 81-200 mrad.  (Rigth) Intensity profile along the centers of a Mo atom and a S2 atomic column.  Intensity profile width of 0.16 nm.

Type of presentation: Poster

IT-5-P-1533 Characterization of EDS Systems with respect to the Geometrical Collection Efficiency

Terborg R.1, Hodoroaba V. D.2, Falke M.1, Käppel A.1
1Bruker Nano GmbH, Am Studio 2D, 12489 Berlin, Germany, 2BAM Federal Institute for Materials Research and Testing, 12200 Berlin, Germany
ralf.terborg@bruker-nano.de

Characteristic parameters are needed to compare the performance of different energy dispersive X-ray spectrometers (EDS). The ISO 15632 standard defines parameters such as energy resolution as the FWHM for the K lines of C, F and Mn. Another crucial feature is the solid angle Ω available for photon collection (Ω=A/r², A: active area of detector, r: distance between radiation origin and center of active detector surface). Ω is not an intrinsic spectrometer property. It can only be defined for a specific detector in combination with a specific system (e.g. SEM), since the minimal possible distance r is determined by the particular detector/microscope geometry. An approach to obtain Ω it is to simply determine A and r but, this is difficult if respective information is not provided by the manufacturer.

For TEM/EDS, the solid angle can be estimated from the ratio of the measured to the theoretical X-ray net count number in a specific element line using a sample of well-known thickness at well-defined acquisition conditions [1]. Parameters such as the detector quantum efficiency, detector and sample geometry as well as electron beam current and quality need to be known for this approach to deliver results close to the real geometric solid angle. Otherwise the strategy can be used to determine just a performance parameter to compare to other TEM-EDS systems.

A similar approach for SEM/EDS systems is to acquire an X-ray spectrum under defined conditions, e.g. a spectrum of a pure Cu bulk sample at 20 keV and known beam current and measure the number of counts in the Cu-K peaks [2]. Using high energy lines reduces the influence of absorption effects, sample surface morphology and contamination. However, some SEMs don't provide the possibility to measure the beam current or don’t have a well calibrated ampere meter. Again, the quantum efficiency must be known. With a significant dead time the input count rate must be used.

A practical approach we suggest for the determination of the real detector-sample distance without need of knowledge of the beam current and detector efficiency is to measure the count rate in a defined energy region for various detector positions retracting the detector in known steps without altering the take-off angle. The count rate I should be proportional to 1/r² and therefore 1/sqrt(I) vs. r should be a straight line through the ordinate origin. This can be used to determine the absolute distance, see Fig. 1, but also to find possible problems with, e.g. shadowing and alignment, which can cause lower count rates, Fig. 2. For the active area A the nominal value can be used. An alternative (if possible) are measurements with apertures of known area placed onto the front of the EDS in a fixed measurement position [3].


[1] R F Egerton, S C Cheng, Ultramicroscopy, 55 (1994), p. 43.

[2] F Schamber, ISO/TC202, Boulder, CO, USA, 2013.

[3] M Procop, Microsc. Microanal, 10 (2004), p. 481.

Fig. 1: Count rate parameter (expressed as cps-1/2) in dependence on the relative position of the EDS for the calculation of the absolute distance from the radiation origin to the detector chip.

Fig. 2: Count rate parameter in dependence on the relative position of the EDS showing shadowing or misalignment leading to a non-linear dependence for small distances.

Type of presentation: Poster

IT-5-P-1543 Optimizing spatial and energy resolution in the TEM at low dose - Lateral Resolved EELS of organic bulk heterojunctions

Kast A. K.1,2, Oster M.1, Pfannmöller M.1,3, Benner G.4, Kowalsky W.2,5, Schröder R. R.1,2,6
1CryoEM, CellNetworks, Universitätsklinikum Heidelberg, Germany, 2InnovationLab GmbH, Heidelberg, Germany, 3EMAT, Antwerp, Belgium, 4Carl Zeiss Microscopy GmbH, Oberkochen, Germany, 5Institut für Hochfrequenztechnik, TU Braunschweig, Germany, 6Center for Advanced Materials, Universität Heidelberg, Germany
anne.kast@bioquant.uni-heidelberg.de

The morphology of organic bulk heterojunction (BHJ) solar cells is strongly correlated to the efficiency of the device [1]. To improve device performance, knowledge of the nanoscale morphology is thus essential. We implement a novel analytical method using Energy Filtered Transmission Electron Microscopy (EFTEM), which allows visualization of donor and acceptor materials in these thin films by analyzing electronic excitation features in the optical and plasmonic electron energy loss region.
Segmentation by Electron Energy Loss Spectroscopy (EELS) reveals that the carbon based organic photovoltaic materials show characteristic optical excitations in energy-loss spectra. However, the blend materials are very sensitive to radiation damage, which impedes also spatial spectral mapping using EELS in conventional scanning beam mode [2].
We introduce an automated scheme which exploits the inherent spatial resolution in the EEL spectrum as it is obtainable from aberration corrected imaging energy filters. It involves automatic scanning of the image of a slit aperture in the illumination beam path. To eliminate residual image distortion in the EEL spectrum we apply additional correction algorithms to facilitate quantitative spectrum interpretation.
Fig. (1a) shows a bright-field image of a polymer:PCBM BHJ and (1b) the spectroscopic image. The polymer-rich areas are represented in green while PCBM-rich areas are red. The slit position during LREEL spectrum acquisition is marked in (Fig. 1a). It was divided into eleven segments from each of which an averaged spectrum was extracted. The acquired spectrum is shown for three segments, in Fig. (2) (electrons are spread horizontally according to energy loss, the zero-loss peak is shown in red) and the averaged spectra from these areas are seen in Fig (3a), (3b) and (3c), respectively. Distinct differences in the spectral information from the different areas are obvious.
The application of such automated laterally resolved EEL spectroscopy (LREELS) as described here allows spatial mapping of high-resolution spectra in two dimensions at low-dose conditions. This is crucial for our understanding of organic BHJ solar cells.

[1] Pfannmöller, M. et al. Nano Lett. 11 (2011) 3099–3107
[2] Egerton, R. F. et al. Micron 43 (2012) 2–7


Financial support by the BMBF (FKZ 03EK3505K, FKZ 13N10794) is gratefully acknowledged.

Fig. 1: a) Bright-field image of a Polymer:PCBM BHJ and b) spectroscopic image depicting the polymer-rich areas in green and the PCBM-rich areas in red. LREELS data was acquired using the slit position marked in a).

Fig. 2: Recorded spectrum after correlation with the slit position. Of the eleven marked blocks (Fig. 1), three are shown representatively. The zero-loss peak is depicted in red. The energy loss is on the horizontal axis. The vertical direction corresponds to the slit length.

Fig. 3: The averaged spectra from areas 1 (a), 2 (b) and 3 (c) as seen in Fig.2. The spectra from the different areas show distinct differences typical for the materials studied here.

Type of presentation: Poster

IT-5-P-1605 Comparisons on Energy and Wavelength Dispersive X-Ray Spectrometry Microanalytical Results: First Approach Based on Iron Ore Sinter Phases

Magalhaes M. S.1,2, Figueiredo e Silva R. C.1,3, Balzuweit K.1,4, Moreira B. B.1, Garcia L. R.5, Persiano A. C.4,5
1Center of Microscopy – Federal University of Minas Gerais (UFMG), 2Consulting & Research , 3Institute of Geosciences - UFMG, 4Department of Physics - UFMG, 5Microanalyses Laboratory of Department of Physics – UFMG
marilias@uol.com.br

Microanalyses of iron ore sinter constituents — hematite, magnetite (Mag), silicoferrite of calcium and aluminium (SFCA) and silicates — are particularly important to handle their impurities. For this reason, these phases need to be well-studied in characterization researches and this had been done for sinters produced with iron ores from Quadrilatero Ferrifero Mineral Province (MG – Brazil). Prior researches had included many qualitative/semiquantitative microanalyses processed by an energy dispersive system (EDS) on a thermo-ionic scanning electron microscopy (SEM). In order to compare different microanalytical approaches and spectrometric devices, which is the main concern of this work, two of these sinter phases have been chosen: SFCA, a typical phase of these sinters, and Mag. At this time, a field emission gun SEM has been employed to acquire qualitative/semiquantitative analyses as much as quantitative ones. In the latter, standards were used. In the same way, a microprobe with wavelength dispersive spectrometers (WDS) was also applied, using similar standards. In a first remark, the relative behavior among the major elements constituents of these phases was retained in spite of the applied method. However, some differences can be highlighted. For SFCA (fig. 1), a similar trend has been observed for a same element, but each one has its proper behavior. Considering all applied spectrometry devices and elaborating the assessment of microanalyses from one sample, some aspects have been observed: Fe stated as Fe2O3 occurs close to 60wt% and inferior to 90wt%, as a rule between 70-85wt%; Ca usually varies near to 10-15wt% of CaO; Si ranges from 3 to 6wt% of SiO2, reaching either higher contents like 8wt% or lower ones as 2wt%; Al commonly achieves 2-4wt% of Al2O3; Mg as major element ranges between 1-2wt% of MgO; in minor quantities, it is close to 1wt% of MgO. For magnetite, a similar trend was also observed for a same element: Fe quantified as Fe2O3 vary from 87 to 97wt% (all spectra), for all devices. For the other elements: SiO2 extends from almost 0.0 to 1.2wt%; Al2O3 varies from 0.5 to 2.0wt%; CaO reaches values from 0.5 to 1.7wt% and MgO achieves contents between 2.5- 9.5wt%. In Mag, Si, Al, Ca and Mg are impurities. Even though diverse behaviors have been observed, it is possible to reproduce the composition of SFCA and Mag with a relative similar evaluation regardless of the applied method, revealing that all types of microanalyses (EDS and WDS) represent the behavior of the discussed elements. In the following steps, the number of microanalyses will be enlarged aiming the validation of this first assessment and to ensure the most appropriate system to understand the phases and related individual requests.


We are grateful for the infrastructure of the Center of Microscopy and of the Microanalyses Laboratory of Department of Physics, both institutions from UFMG.

Fig. 1: The diagram shows Fe2O3 mean distribution in some portions of SFCA, a typical iron ore sinter phase, considering EDS and WDS devices. Stdless: Standardless; STD-A: General Standard; STD-B: Particular Standard for SFCA; STD-C: Particular Standard for Magnetite; No-N: no-normalized.

Fig. 2: Other diagram showing Fe2O3 mean distribution in some portions of SFCA, a typical iron ore sinter phase, considering EDS and WDS devices.Jeol: JSM-5410 thermo-ionic SEM with a NORAN TN-M3055 spectrometer; FEI Company: Quanta-200 (Q200) FEG SEM with an EDAX Sapphire Si(Li) spectrometer; Jeol: JXA 8900R microprobe with four WDS spectrometers.

Fig. 3: The diagram shows the elements mean distribution in some portions of SFCA, a typical iron ore sinter phase, using EDS and WDS devices. The elements are stated as oxides - SiO2; Al2O3; CaO; MgO. See figure captions on figures 1 and 2.

Fig. 4: Other diagram showing the elements mean distribution in some portions of SFCA, a typical iron ore sinter phase, using EDS and WDS devices. The elements are stated as oxides - SiO2; Al2O3; CaO; MgO. See figure captions on figures 1 and 2.

Type of presentation: Poster

IT-5-P-1618 Magnified pseudo atomic column elemental maps realized by STEM moiré method

Okunishi E.1, Kondo Y.1
1EM Business Unit, JEOL Ltd., 3-1-2 Musashino Akishima Tokyo 196-8558, Japan.
kondo@jeol.co.jp

In recent years, modern technologies such as aberration correction realized an atomic column elemental mapping with analytical tools such as electron energy loss spectroscopy (EELS) and/or energy dispersive X-ray spectroscopy (EDS) [1,2], which is useful, since atomic species and positions in a crystalline specimen can be determined directly. A crucial issue to perform the mapping is specimen damage due to high electron dose onto a specimen, since an excitation probability for core electrons is small.
STEM moiré fringes for a periodic lattice arise when a pixel interval is close to a lattice spacing, due to the under-sampling effect [3]. With the proper pixel intervals in x and y, the moiré fringe shows the pseudo 2D magnified moiré lattice, which is homothetic to the original lattice [4]. The magnification (M) of moiré lattice to the original lattice is determined as  M = |1 – r|-1, where r is the ratio of the lattice spacing to the pixel interval. A magnified moiré lattice is formed with under-sampled signals picked from original lattices, resulting in reduced electron dose by M-2 on the specimen to form an atomic column with an equal pixel resolution. This paper reports a method to observe the atomic column elemental map with less electron dose and higher pixel resolution, utilizing the STEM-moiré method.

The specimen for our experiment was SrTiO3 [001] that has a square lattice. The microscope used was a Cs corrected microscope (JEM-ARM200F), equipped with a SDD type EDS. Figs. 1(a-f) show the high angle annular dark field (HAADF) and annular bright field (ABF) [5] images of magnified moiré lattice at various r. The magnification (M) increases as r approaches one. Figs. 2(a-i) show simultaneously obtained elemental maps of Sr, Ti and O, detected with an EDS. Each element was clearly separated on each magnified moiré elemental map. No beam damage on the specimen was observed during the experiment. Fig. 3 shows three line profiles along a (Ti+O)-(O) row of O-Kα map, Ti-Kα map and ABF image shown in Figs. 2(g,d) and 1(d). The profiles clearly show the peaks at oxygen sites.

In conclusion, the STEM moiré method was successfully applied to atomic column elemental mapping. The method can be applicable to measure detailed physical properties such as delocalization or channeling in crystalline specimens with higher pixel resolution, better signal-to-noise ratio and less electron dose than the direct atomic column mapping.

References:
[1] K Kimoto et al., Nature 450 (2007), p. 702.
[2] E Okunishi et al., Microsc. Microanal. 12 (Suppl. 2) (2006), p. 1150.
[3] N Endo and Y Kondo, Microsc. Microanal. 19 (Suppl. 2) (2013), p. 346.
[4] D Su and Y Zhu, Ultramicroscopy 110 3 (2010), p. 229.
[5] E Okunishi et al., Microsc. Microanal. 15 (Suppl. 2) (2009), p. 164.


Authors thank to Mr. Hosokawa of JEOL Ltd. for valuable discussion on a theoretical consideration.

Fig. 1: High angle annular dark field (HAADF) and annular bright field (ABF) images of magnified moiré lattice at various r.

Fig. 2: Elemental maps of magnified moiré lattice at various r, detected with the EDS.

Fig. 3: Line profiles along a (Ti+O)-(O) row of the O-Kα map, the Ti-Kα map and the ABF image shown in Figs. 2(g), 2(d) and 1(d).

Type of presentation: Poster

IT-5-P-1630 How to Avoid Unexpected Artifacts from Multivariate Statistical Analysis on STEM Spectrum-Imaging Datasets

Ishizuka K.1, Watanabe M.2
1HREM Research Inc., Higashimatsuyama, Japan, 2Dept of Materials Science and Engineering, Lehigh University, Bethlehem, USA
ishizuka@hremresearch.com

The latest aberration-corrected scanning transmission electron microscope (STEM) makes possible to perform routinely not only atomic-scale imaging but also chemical analysis via electron energy-loss spectrometry (EELS) and X-ray energy dispersive spectrometry (XEDS) [e.g. 1]. In combination with the latest hardware, the advances in the recent software developments allow us to acquire large-scale datasets such as multidimensional image series and spectrum images (SIs). Therefore, it is challenging to deal with the large-scale datasets, e.g. extraction of unknown features and estimation of dominant trends. If the datasets were relatively noisy, which is very common for atomic-resolution EELS/XEDS SIs, data analysis would be much harder tasks. Multivariate statistical analysis (MSA) is one of efficient approaches to analyse the large-scale datasets in terms of feature identification and extraction.
Principal component analysis (PCA) is one of the MSA techniques [2]. Since a use of PCA is relatively straightforward, PCA has been applied to SIs as data-mining and noise-reduction tools [e.g. 3]. The PCA tries to explain the data variation (variance) as much as possible using a small number of the components. Here, the signal itself of course contributes the data variation. However, a small amount of signal will be buried with the whole random noise. Therefore, despite that the PCA approach is very efficient and useful, it may create unexpected artifacts especially in higher noise conditions [4] (Figure 1). Since these artifacts might mislead results, it is essential to avoid such artifacts. There may be two approaches to improve the PCA sensitivity: (1) reduction of random noise and (2) enhancement of true variations. The former requires modifications in experimental conditions (higher currents and longer acquisitions). Conversely, the latter can be achieved by PCA analysis to divided small segments within a SI, which is called the local PCA approach (Figure 2). The division can be made spatially and spectrally. The spatially local PCA will be especially useful to detect segregated element in the matrix. The spectrally local PCA is useful to detect a weak signal, if the weak signal is spectrally separated from the strong signal. Especially the spectrally local PCA is useful for EELS Sis, since the background intensity varies significantly. In this study, advantages of the local PCA approach will be addressed.
[1] S.J. Pennycook & P.D. Nellist ed. Scanning Transmission Electron Microscopy: Imaging and Analysis, Springer, NY, (2011).
[2] E.R. Malinowski, Factor Analysis in Chemistry, 3rd ed., Wiley, New York, (2002).
[3] M Watanabe et al., Microscopy and Analysis, 23, Issue 7, (2009), 5-7.
[4] S. Lichtert & J. Verbeeck, Ultramicrosc., 125 (2013), 35-42


The authors acknowledge J. Verbeeck for providing the simulated BN test data. M.W. wishes to acknowledge financial support from the NSF through grants DMR-0804528 and DMR-1040229.

Fig. 1: Failure of PCA [4]. (a) BN model, where there are one excess N or B atom at the positions 1 and 2, respectively. (b) Untreated and weight PCA N element maps for low and high noise. Note that the N map of high noise shows higher intensity at 2 than 1.

Fig. 2: (a) and (b): Spatially local PCA and spectral local PCA, respectively. (c): N maps reconstructed by the spectrally local PCA shows higher intensity at the position 1 even for the high noise case contrary to the normal PCA.

Type of presentation: Poster

IT-5-P-1706 Formation of oxidatively stable M@Fe3O4 and MPt@Fe3O4 (M = Fe, Co) core@shell nanoparticles using a simple and versatile synthetic procedure

Knappett B. R.1, Gontard L. C.2, Ringe E.1, Tait E. W.1, Fernandéz A.2, Wheatley A. E.1
1University of Cambridge, 2Instituto de Ciencia de Materiales de Sevilla
bk324@cam.ac.uk

Core@shell nanoparticle synthesis offers the ability to create materials with dual characteristics, such as a magnetic core and a functionalisable or catalytically active shell. They are currently the subject of extensive research due to the tuneability of their structure and therefore properties.1 Our research focusses on the coating of magnetically interesting materials to protect against oxidation. Iron and cobalt nanoparticles are both strongly magnetic, but are highly susceptible to oxidation. Much work has been carried out to coat these materials in organic surfactant and polymer layers in an attempt to protect against core oxidation;1 however the coating procedure used in the current work utilises a 2-3 nm inorganic layer of Fe3O4, formed by the decomposition of Fe(CO)5 in solution, to stabilize the particle cores. The recently published Co@Fe3O4 system has been shown to have detectable levels of carbon present after oxidative plasma cleaning. This carbon must therefore be contained within the particle structure, suggesting that surfactant molecules that capped the Co seeds became trapped during shell formation.2 This has been verified using scanning transmission electron microscopy (STEM) to record electron energy loss spectroscopy (EELS) line scans and point scans with very high spatial resolution. Figure 1 shows the use of EELS point scans on a Co@Fe3O4 sample to confirm that, after plasma cleaning, no carbon can be detected at the outer surface of the particle shell, yet carbon remains detectable at the core-shell interface. It is also observable that in spite of the highly oxidative plasma cleaning process, the Co core remains metallic in nature. Figure 2 shows a line scan through an Fe@Fe3O4 nanoparticle, confirming a structure similar to that of Co@Fe3O4.

Recently, research has focused on utilizing the coating procedure to encapsulate FePt and CoPt alloys. These particles have interesting magnetic properties for applications in magnetic arrays for data storage. However, once synthesised, these particles require annealing, which often causes sintering. It has previously been established for FePt particles that a coating of iron oxide will prevent the particles from sintering during annealing.3 However, the coating procedure used was different to that employed in our work. We are seeking to prove that the same stabilization is granted the particles by our coating procedure, and to further extend the annealing studies to the system of CoPt, the coating of which has not before been reported.

References
[1] Ghosh Chaudhuri, R.; Paria, S., Chem. Rev., 112 (2012), 2373.
[2] Knappett, B. R. et al., Nanoscale 5 (2013), 5765.
[3] Liu, C. et al., Chem. Mater. 17 (2005), 620.


The authors would like to acknowledge financial support from The Junta de Andalucia (FEDER PE2009-FQM-4554, TEP-217) and EU FP7 AL-NANOFUNC project (CT-REGPOT2011-1- 285895). B. R. K. thanks the UK EPSRC, The University of Cambridge and Downing College for grants. E. R. acknowledges support from the Royal Society in the form of a Newton International Fellowship.

Fig. 1: Electron energy loss spectroscopy (EELS) point scans of the support (Si3N4, shown in blue) and the boundary of the core and shell of a particle. The signal clearly shows the presence of carbon within the structure of the particle after plasma cleaning.

Fig. 2: EELS line scan of an Fe@Fe3O4 particle, evidencing the presence of carbon within the structure of the particle. Again, the sample had been plasma cleaned, thus removing all carbon-containing ligand molecules external to the particle structure.

Type of presentation: Poster

IT-5-P-1713 Detection and quantification of phosphorus dopants in germanium by energy dispersive spectrometry using an annular detector with large solid angle

Robin E.1, Mollard N.1, Guilloy K.1, Pauc N.1
1CEA-INAC, Grenoble, France
eric.robin@cea.fr

Energy dispersive spectrometry (EDS) is generally not considered as a sufficiently sensitive and accurate technique for dopant detection and quantification. Indeed, the concentrations of traditional dopants are typically below ≈ 1020 at cm-3, which is very close to the detection limit of conventional EDS.

We report here the detection and quantification by EDS of phosphorus (P) dopant concentrations in germanium as low as 5 1018 at cm-3 with a precision and detection limit around 1018 at cm-3. This is achieved by using the Flat Quad 5060F annular detector recently developed by Bruker-AXS. This new generation of silicon drift detectors is composed of four bean-shaped silicon diodes, each of 15 mm2, arranged in a ring around a 1.6 mm central hole for the electron beam passage (Fig.1a). It is positioned a few millimeters above the sample (Fig.2b), a geometry which results in a much wider solid angle (up to 1.1 sr) compared to traditional detectors (<0.1 sr), thus allowing a higher counting rate at any operating conditions (up to 1000 kcps). The passage of the electron beam through the detector precludes the use of a conventional electron trap, which role is to protect the diodes against the backscattered electrons. To prevent detector damage, three mylar windows are mounted on the detector, the first (1 µm thick) being permanent to operate in the range 0-6 kV, the two others (2 and 6 µm thick) being retractable to operate in the range 6-12 kV and 12-20 kV, respectively. Although it was not their primary function, the two retractable windows may act as a high-pass X-ray energy filter allowing enhancement of the detection sensitivity for high energy X-rays. For instance, the insertion of window 2 enhances by a factor of 2 the counting rate in the P region (Fig. 2a). The Ge pile-up is also reduced due to the absorption of the Ge L lines, which also improves the detection of low concentrations of P dopants.

We tested five Ge 2D layers previously analyzed by Tof-SIMS and containing 0.66, 0.71, 0.98, 2.5 and 36 1019 at cm-3 of P dopants. Samples were analyzed at 4 different voltages (3, 4, 6 and 8 kV) with the window 2 inserted. All spectra were acquired for 2 hours at ≈ 500 kcps and normalized to pure Ge spectra (Fig. 2b). The reproducibility was tested by repeating the analyses at least three times. Results show a relatively good consistency with Tof-SIMS results, even for the lowest concentrations of P dopants (Fig. 2c). The reproducibility is within the analytical uncertainty of the counting statistic. The precision and the detection limit depend on the voltage and the total acquisition time (Fig. 2d). Typically at 4 kV on P doped Ge nanowires, it is around 2 1018 at cm-3 and 1018 at cm-3 for 30 minutes and 2 hours of acquisition time, respectively.


Fig. 1: The Flat Quad 5060F annular detector from Bruker-AXS: a) bottom view showing the four bean-shaped Si diodes (d), the central hole (h) and the first retractable mylar window (w); b) top view with the sample in place.

Fig. 2: a) 6 kV EDS-FQ spectra acquired on pure Ge using increasing thickness windows; b) 4 and 8 kV Ge-normalized EDS-FQ spectra acquired with window 2 on P doped Ge 2D layers; c) Comparison of P dopant concentrations between EDS-FQ and ToF-SIMS; d) precision/detection limit for P dopants by EDS-FQ.

Type of presentation: Poster

IT-5-P-1781 Effect of the asymmetry of dynamical electron diffraction on intensity of acquired EMCD signals

Song D. S.1, Wang Z. Q.1, Zhu J.1
1National Center for Electron Microscopy in Beijing, School of Material Science & Engineering, Tsinghua University, Beijing 100084, China
todongsheng@126.com

One of the most challenging issues to characterize magnetic materials in the transmission electron microscopy is to obtain quantitative magnetic parameters on the nanometer scale. By the technique of electron energy-loss magnetic chiral dichroism (EMCD) is proposed and applying the sum rules, it is possible to quantitatively extract the orbital to spin magnetic moment ratio mL/mS with high spatial resolution. Compared with the technique of XMCD, the detection source of EMCD technology are the transmission electrons rather than the X-ray based on the precious synchrotron radiation. Therefore, the dynamical diffraction effects of electrons are quite remarkable in the periodic crystal structures, making the quantitative EMCD technique more complicated. By establishing the quantitative relation between EMCD and dynamical diffraction effects, spin and orbital moment of different elements and nonequivalent crystallographic sites are quantitatively determined in a spinel structure NiFe2O4 [1].
However, the diffraction geometry in EMCD experiment is strict and conditions of symmetric detector positions should be fulfilled. It has been reported that the inherent asymmetry of the two-beam geometry can lead to systematic errors in quantitative EMCD measurements [2]. Besides, the asymmetry of dynamical coefficients in the three-beam geometry also exists and is neglected in the previous study. Here, we point out that the asymmetry of dynamical electron diffraction should be accounted and its impact on the quantitative measurements of the EMCD signal needs to be evaluated.
Reference:
[1] Z.Q. Wang, X.Y. Zhong, R Yu, Z.Y. Cheng, J Zhu. Quantitative experimental determination of site-specific magnetic structures by transmitted electrons. Nature communications, 2013, 4: 1395.
[2] J. Rusz, P.M. Oppeneer, H. Lidbaum, S. Rubino, K. Leifer. Asymmetry of the two‐beam geometry in EMCD experiments . Journal of microscopy, 2010, 237(3): 465-468.


This work is financially supported by National 973 Project of China (2009CB623701) and Chinese National Nature Science Foundation (11374174,51390471 ). This work made use of the resources of the Beijing National Center for Electron Microscopy and Tsinghua National Laboratory for Information Science and Technology.

Type of presentation: Poster

IT-5-P-1792 Spatial and Temporal Coherences in Spin-Polarized Transmission Electron Microscopy

Kuwahara M.1, Kusunoki S.1, Nambo Y.1, Saitoh K.2, Ujihara T.1, Asano H.1, Takeda Y.3, Tanaka N.2
1Graduate school of Engineering, Nagoya University Nagoya, 464-8603, Japan, 2EcoTopia Science Institute, Nagoya University, Nagoya 464-8603, Japan, 3Nagoya Industrial Science Research Institute, Nagoya 464-0819, Japan
kuwahara@esi.nagoya-u.ac.jp

Great advances have recently been made in magnetic recording technology and spintronic devices, which are promising for high-density storage devices. Such devices are expected to lead to the development of systems that can analyze magnetic and spin states with a nanometer-order spatial resolution.

We have commenced a development of a spin-polarized transmission electron microscope (SPTEM), which consists of a polarized electron source (PES) and a conventional TEM [1-3]. Figure 1 shows a photograph of the SP-TEM. Spin-polarized electrons can be generated using an optical orientation of III–V semiconductors and vacuum extraction that uses a negative electron affinity (NEA) surface. Several beam parameters of the PES are vastly superior to those of conventional thermal electron beams. In addition, it has the ability to generate a sub-picosecond multibunch beam[4]. A high ESP of 92% and a high QE of 0.5% have been realized using a GaAs–GaAsP strained superlattice photocathode[5].

We have already demonstrated that the SPTEM can provide both TEM images and the diffraction patterns [1]. The TEM images can be obtained in a spatial resolution of 1 nm in a 30-kV acceleration voltage. The apparatus has a below 240-meV energy width of electron beam in the TEM without any monochrometors (Fig. 2). The energy width indicates the temporal coherence is about 2.7 fs (longitudinal coherence of 2.7×10-7 m) at 30-keV beam energy. A brightness is directly measured by taking a spot size and a convergent angle on an image plane. The measured brightness is about 4×107 A/cm2sr in a 30-keV beam energy with a polarization of 82 % and the drive-laser power of 800 kW/cm2 on the photocathode [6]. The brightness for a 200-kV beam energy is 3×108 A/cm2sr which is converted by using a Lorentz factor. The order of the brightness is enough to do an interference experiment. We also demonstrated interference fringes of spin-polarized electron beam by using a newly installed biprism as shown in figure 3. These results indicate the SP-TEM can provide enough coherence in both lateral direction and longitudinal direction even if the semiconductor photocathode is used for an electron emitter.

[1] M. Kuwahara et al., Appl. Phys. Lett. 101 (2012) 03310

[2] M. Kuwahara et al., AMTC Letters 3 (2012) 180.

[3] M. Kuwahara et al., J. Phys.:Conf. Ser. 298 (2011) 012016.

[4] Y. Honda, et al., Jpn. J. Appl. Phys. 52, 086401-086407(2013).

[5] X.G. Jin et al., Appl. Phys. Express 1 (2008) 045002.

[6] M. Kuwahara et al., to be submitted (2014).


The authors thank Drs. H. Shinada, M. Koguchi and M. Tomita of Hitachi Central Research Laboratory for fruitful discussions and encouragement. This research was supported by MEXT KAKENHI Grant Number 51996964, 24651123, 25706031 and Kurata Research Grants from the Kurata Foundation.

Fig. 1: Photograph of the spin-polarized TEM.

Fig. 2: Energy spread of spin-polarized electron beam as a function electron energy.

Fig. 3: Interference fringe of spin-polarized electron beam extracted from a GaAs-GaAsP strained superlattice photocathode using a biprism.

Type of presentation: Poster

IT-5-P-1793 Measurement of liquid vibrational spectra using monochromated STEM-EELS

Miyata T.1,2, Fukuyama M.1,3, Hibara A.3, Ikuhara Y.2,4, Okunishi E.5, Mukai M.5, Mizoguchi T.1
1Institute of Industrial Science, the University of Tokyo, Tokyo, Japan, 2School of Engineering, the University of Tokyo, Tokyo, Japan, 3School of Science, Tokyo Institute of Technology, Tokyo, Japan, 4Japan Fine Ceramics Center, Aichi, Japan, 5JEOL Ltd., Tokyo, Japan
tomo-m@iis.u-tokyo.ac.jp

  Liquid is widely used in daily life and industrial activities. The dynamic behavior of the molecules in liquid is an important factor to determine the various liquid properties. The dynamic behavior of liquid molecules has been extensively investigated using infrared (IR) spectroscopy and Raman spectroscopy. However, these spectroscopy techniques allow us to obtain only averaged information of the entire sample. On the other hand, a specific location, such as solid-liquid interface, plays important role for reactions in electrochemistry and organic chemistry, in which liquid is treated as reactants and reaction media. That is, the methods for analyzing the dynamic behavior of liquid molecules at high spatial resolution have been desired.
  In this presentation, thus, we will report the results of the measurements of liquid vibrational spectra by monochromated STEM-EELS. For the analyses, I used an aberration corrected STEM with a monochromator (JEM-2400FCS, JEOL Ltd., 120keV). The energy resolution reached 0.065eV using the monochromator. As a liquid sample I chose a popular ionic liquid, 1-ethyl-3-methylimidazolium bis (trifluoromethyl-sulfonyl) imide (C2mim-TFSI). In order to verify the vibrational spectra by STEM-EELS, an IR spectrum was measured from the same sample. In addition, first principles calculations were performed to interpret the peaks in the vibrational spectrum. The plane-wave pseudopotential method (CASTEP code) was used in the calculations.
  From the STEM-EELS measurement, the HOMO-LUMO gap of C2min-TFSI was estimated to be 5.3eV, which is consistent with the results of the first-principles calculations and a separately measured ultraviolet-visible (UV-vis) spectrum. The peaks ascribed to the molecular vibration were measured in the vicinity of 0.4eV. These peaks were also observed in the IR spectrum and the one from the first-principles calculations. From those analyses, it was confirmed that the peaks at the 0.4eV correspond to the CH bonds stretching peaks. Based on this study, we have demonstrated that the vibrational peaks of the nano area in liquid are available by the monochromated STEM-EELS.


This study was supported by of MEXT and JSPS (22686059, 25106003). Some calculations were performed using a supercomputer at Institute of Solid State Physics of the University of Tokyo.

Type of presentation: Poster

IT-5-P-1811 Features and Applications of an Analysis System using Double Silicon Drift-type Detectors for Transmission Electron Microscopy

Ohnishi I.1, Kawai S.1, Ishikawa T.1, Yagi K.1, Iwama T.1, Miyatake K.1, Iwasawa Y.1, Matsushita M.1, Kaneyama T.1, Kondo Y.1
1JEOL Ltd., Tokyo, Japan
ionishi@jeol.co.jp

  The elemental analysis using energy dispersive X-ray spectroscopy (EDS) requires considerably long time due to a small ionization cross section for core electrons of a specimen and a small solid angle for a detector. Therefore, the realization of an efficient, short time analysis not only serves the increasing needs of users but also reduces specimen damage due to electron irradiation. For this purpose, we have developed a new analysis system, which is extremely sensitive than the one at present. We report the features and applications of the newly developed analysis system, which is composed of two silicon drift-type detectors (SDDs).

  Our newly developed analysis system consists of two SDDs (double-SDD) with a large sensor area (100 mm2 in area). Figure 1 shows the schematic configuration of the system for a field emission electron microscope (JEM-2800). A new TEM column has two ports for the detectors. And a special analytical specimen holder, which is thinner than the present one, has been newly developed to reduce the distance between the detector and specimen.

  The total X-ray intensity of the new system has increased, because the X-ray signals collected from two detectors are integrated. The new specimen holder also promotes higher X-ray collection efficiency. As a result, total sensitivity has been significantly improved. For example, the peak intensity of Al K line obtained with a double-SDD has increased to be approximately 1.7 times higher than that obtained with a single SDD as shown in Fig. 2.

  Since the new analysis system can provide high analytical sensitivity, an atomic resolution elemental map with a high S/N ratio can be acquired in combination with Cs-corrected TEM/STEM machines such as JEM-ARM200F. In the upper parts of Fig. 3 are shown atomic resolution elemental maps sized 128 x 128 pixels for SrTiO3<100>. These maps clearly show atomic columns of Sr, Ti and O. The profiles of elemental columns, displayed in the lower parts of Fig. 3, show a significantly improved signal to noise ratio for the double-SDD compared with the one for the single SDD.

  Our new system has significantly high X-ray collection ability. Therefore, it provides a shorter acquisition time than a single SDD system. It helps us, beyond all doubt, to analyze a beam sensitive specimen and to detect trace elements in a specimen.


Fig. 1: Fig. 1: A schematic configuration of an EDS analysis system for a field emission electron microscope (JEM-2800).

Fig. 2: Fig. 2: EDS spectra of an Al foil specimen, obtained by using JEM-2800 (200 kV) with SDD1+SDD2 (red) and SDD1 (blue), respectively. The vertical axis has been normalized by the peak intensity of Al K line, obtained with SDD1.

Fig. 3: Fig. 3: STEM-ADF images and atomic resolution elemental maps (O, Ti, Sr) for SrTiO3<100>, obtained with JEM-ARM200F (200 kV) with SDD1 (left) and SDD1+SDD2 (right). The mapping sizes are 128 x 128 pixels. Line profiles of O Kα, Ti Kα, and Sr Lα gross intensity extracted from yellow lines in ADF images are also shown below the maps.

Type of presentation: Poster

IT-5-P-1826 Orientational dependence of EMCD signals of hcp Co with strong magnetocrsytalline anisotropy

Kudo T.1, Tatsumi K.2, Leifer K.3, Rusz J.4, Muto S.2
1Graduate School of engineering, Nagoya University, 2Eco Topia Science Institute, Nagoya University, 3Department of Engineering Science, Uppsala University, 4Department of Physics and Astronomy, Uppsala University
kudou.tomohiro@h.mbox.nagoya-u.ac.jp

Electron magnetic circular dichroism (EMCD) is an electron spin-related property of ferromagnetic samples revealed as the difference of the EELS inner shell spectra measured at two specific positions on the diffraction plane [1]. EMCD at transmission geometry can be advantageous to the X-ray counterpart, XMCD, in spatial resolution and probing bulk properties since the L2,3/M4,5 white-lines of transition metals/rare earth elements, to which the sum rule is applicable, are located in the soft X-ray energy region. For these advantages, a number of experimental schemes have been proposed for better quantitative measurement and higher spatial resolution [2].
In the intrinsic EMCD experimental scheme, with magnetization of a crystalline sample aligned along the strong magnetic field of the objective lens, the symmetric two- or three-beam condition is required [1]. The dichroic signal, Δσ, is acquired as the difference between the two ELNES spectra measured at the two positions, A and B (Fig.1-(a) and (b)). According to the inversion sum-rule [3], the dichroic signal intensity is approximately proportional to (q×q')∙M , where q and q' are the inelastic scattering vectors respectively pointing from O and G to the two detector positions lying on the Thale circle (cf., Fig. 1), and M is the magnetization of the sample. The measured signal intensity is proportional to M·H (H: external magnetic field, parallel to the optic axis) if the direction of M is not fully saturated in the direction of the external magnetic field, which has not yet been experimentally exploited.
We measured EMCD signals of L2,3 in hcp Co, a hard magnet, exhibiting relatively larger magnetocrystalline anisotropy. A thin sample was prepared by electrochemical polishing. In Fig.2-(a) and (b) are shown spectra collected at the two different geometries of Fig.1-(a) and (b), where the optical axis is nearly parallel and perpendicular to the [001] easy magnetization direction of hcp Co with the low-order systematic row excitation conditions. The spectral intensities are normalized by the L3 peak collected at the detector position A. The magnetic dichroism is clearly enhanced in the geometry (a), compared to (b), confirming the unsaturated magnetization along the external magnetic field.
Moreover, a theoretical simulation [3] predicts that with the specific EELS detector positions on the three-beam condition the spin moment in the plane normal to the optical axis can be probed, the trial of which is also presented.
References
[1] P. Schattschneider et al. Nature 441, 486 (2006)
[2] S. Muto et al., Nature Commun., 5, 3138 (2014): doi:10.1038/ncomms4138
[3] J. Rusz et al., Phys. Rev. B 84, 064444 (2011)


A part of this work was supported by a Grant-in-Aid on Innovative Areas "Nano Informatics" (grant number 25106004) and on Young scientist A (24686070) from the Japan Society of the Promotion of Science. J.R. acknowledge support from the Swedish Research Council and STINT.

Fig. 1: Schematics of two experimental geometries in the present study. Optical axis is nearly parallel (a) and perpendicular (b) to easy magnetization axis, respectively.

Fig. 2: Experimental Co-L2,3 ELNES and difference spectra respectively corresponding to geometries (a) and (b) in Fig. 1.

Type of presentation: Poster

IT-5-P-1842 Advanced SEM/EDS analysis using an Annular Silicon Drift Detector (SDD): Applications in Nano, Life, Earth and Planetary Sciences below Micrometer Scale

Salge T.1, Terborg R.1, Ball A. D.2, Broad G. R.2, Kearsley A. T.2, Jones C. G.2, Smith C.2, Rades S.3, Hodoroaba V. D.3
1Bruker Nano GmbH, Berlin, Germany, 2Natural History Museum, London, United Kingdom, 3BAM Federal Institute for Materials Research and Testing, Division 6.8 Surface Analysis and Interfacial Chemistry, Berlin, Germany
tobias.salge@bruker-nano.de

Analysis of fine-scale structures requires low accelerating voltages. Consequently, only low to intermediate energy X-ray lines with many peak overlaps can be evaluated which requires deconvolution. Examination of nano-scale structures also requires low probe currents which would give low X-ray count rates with traditional EDX detectors. The additional time required to acquire sufficient data for deconvolution risks altering the specimen as a result of beam-sample interaction or sample contamination. The BRUKER XFlash 5060F SDD has allowed us to overcome these limitations and offers additional benefits as we demonstrate here.

The annular SDD is inserted between the pole piece and sample and has a large solid angle (1.1sr). It is ideally suited for the analysis of topographically complex, three-dimensional samples. X-rays are collected from 4 separate detector segments and signals are processed in parallel by 4 detection units allowing high count rates at low dead-time. Even at lowest beam current (<10pA), samples can be investigated under high vacuum in natural state. In VP mode, sufficient statistics can be collected with reduced acquisition times, consequently reducing the likelihood of sample contamination.

4 studies are presented here. Experimental impact foil craters (Fig. 1) used to develop methodologies for the examination of samples from NASA’s STARDUST mission were examined (6kV, 1,100kcps, 45% dead time, 4096 x 3072 pixels, 62nm pixels, 7 min). Residues of glass projectiles can be clearly distinguished from the aluminium target with no detector shadowing across the field of view. Fig. 2 shows a portion of the Martian meteorite “Tissint”. This analysis was carried out at low beam current (4kV, <10pA, 733x853 pixels, 55nm pixels, 13 h). The mapped area revealed a thin coating and local enrichment of carbon (and nitrogen). The third study (Fig. 3) was carried out under low vacuum (20Pa, 5kV, 1.8nA, 20kcps, 320x240 pixels, 460nm pixels, 33 min) and shows mineralization in the ovipositor of a parasitoid wasp. Sufficient data quality allows deconvolution of overlapping element lines (Zn-L, Na-K). The final study (Fig. 4) demonstrates that the analysis of core-shell nano particles (5kV, 520pA, 14-72kcps, 250x250 pixel, 2nm pixels, 6 min) on thin film supports has become possible.

It can be concluded that improvements in SDD technology will stimulate new approaches for various fields. To minimize acquisition time and increase spectrum statistics, the total solid angle is relevant, not the active detector area. Element analysis at low kV and low beam current in combination with multi-segment SDDs provides high spatial resolution and high detection sensitivity without the necessity of applying a conductive coating or working in low vacuum.


Fig. 1: Stardust analogue crater experiment; Composite EDX map overlain with SE micrograph reveals residues of glass projectiles (red) on an aluminium target (green).

Fig. 2: Tissint Martian meteorite; Composite EDX map of carbon and oxygen overlain with SE micrograph shows a thin coating and local enrichment of carbon.

Fig. 3: Biomineralization in a parasitoid wasp Monolexis fuscicornis. The ovipositor (sting and egg-layer) reveals ZnO reinforcement and contamination with NaCl.

Fig. 4: Composite EDX map of fluorescent silica nano particles. At the line scan (net intensities, 229 points, 467nm length, 30kcps, and 6.9 s), five adjacent pixel/spectra were binned for each point in order to improve impulse statistics.

Type of presentation: Poster

IT-5-P-1944 Planer defect structure analysis based on electron channeling phenomena

Ichikawa T.1, Ohtsuka M.1, Muto S.2
1Nagoya University, Nagoya, Japan, 2EcoTopia Science Institute, Nagoya, Japan
ichikawa.takahiro@a.mbox.nagoya-u.ac.jp

 High-angular resolution electron channeled X-ray spectroscopy (HARECXS) is based on the principle that the electron wavefunctions (Bloch waves) in a crystalline sample change their symmetry with the incident beam direction. HARECXS has been applied to the analysis of point defects, partly because ICSC [1] has been only the theoretical simulation code available to date for predicting the inelastic scattering cross sections depending on the diffraction condition, which does not allow us to include 2D/3D defects. In this study an attempt is made to apply HARECXS to a planner defect lying on the {111} plane in heavily Si-doped GaAs [2] to extend its applicability.

 Thin films were prepared by dimpling the sliced single crystalline wafers followed by Ar ion milling. HARECXS was carried out using a JEM-2100 S/TEM equipped with an EDX spectrometer, operated at the beam-rocking mode at 200kV. The illuminated area was about 1μm in diameter and the incident beam angle was rocked by ±1.5 degrees with a step of 0.05 degrees under the systematic row excitation condition.

 A bright-field TEM image of the planer defect nearly viewed end-on is shown Fig. 1(a) and the corresponding HARECXS profiles in (b), as functions of the incident beam direction in units of g1-11, tilted in the direction perpendicular to the projected defect plane. The HARECXS profiles of the Ga- and As-K lines have a different symmetry due to the polarity of the GaAs. The slight asymmetric profile of Si-K similar to that of As-K suggests that Si mainly occupies the As sites. The theoretical simulations based on the model where a Si atom substitutes for the As and Ga sites are shown in Fig. 2(a), using the Bloch wave [1] method with the dynamical inelastic scattering process incorporated. For comparison, more realistic models where Si occupies the single (111) Ga or As layer are also simulated, as shown in Fig. 2(b), in which a multislice method [3] is developed for including a planar defect in the simulation. The simulation result with Si occupying the As single layer seems to be relatively more consistent with the experimental Si-K HARECXS profile.

References
[1] M. P. Oxley and L.J. Allen, J. Appl. Cryst. 36 (2003) 940-943.
[2] S. Muto, S. Takeda, M. Hirata, K. Fujii, and K. Ibe, Phil. Mag. A. 66 (1992) 257-268.
[3] M. Ohtsuka, et al., AMTC4 letters (2014) in publish


A part of this work was supported by a Grant-in-Aid on Innovative Areas "Nano Informatics" (Grant number 25106004) from the Japan Society of the Promotion of Science.

Fig. 1: (a)Bright-field TEM image of planer defect in Si-doped GaAs nearly viewed end-on. (b)HARECXS profiles of Ga-K, As-K, and Si-K characteristic X-ray peaks as functions of incident beam direction in units of g1-11.

Fig. 2: Theoretically simulated HARECXS profiles for models where Si atom substitute for Ga (red) and As (blue) sites (a) and Si occupies single (111) Ga (red) and As (blue) atom layers (b).

Type of presentation: Poster

IT-5-P-1976 EELS of Si-nanocrystals by hyperspectral segmentation and multivariate factorization.

Eljarrat A.1, López-Conesa L.1, López-Vidrier J.1, Hernández S.1, Estradé S.1, 2, Magén C.3,4, Garrido B.1, Peiró F.1
1MIND-IN2UB, Departament d'Electrònica, Universitat de Barcelona, c/ Martí i Franqués 1, 08028 Barcelona, Spain., 2TEM-MAT, Centres Científics i Tecnològics (CCiT), Universitat de Barcelona, Solís Sabarís 1, Barcelona, Spain., 3LMA-INA, Departamento de Física de la Materia Condensada, Universidad de Zaragoza, 50018 Zaragoza, Spain., 4Fundación ARAID, 50018 Zaragoza, Spain.
aeljarrat@el.ub.edu

This work is focused on advanced data analysis methods for the characterization of Si NCs by high angle annular dark field (HAADF) and electron energy loss spectroscopy (EELS) in the aberration corrected and monochromated scanning transmission electron microscope (STEM). These Si-NCs, of high interest for photovoltaic applications, are embedded in multilayer stacks where SiO2, SiC and Si3N4 are used as dielectric barriers.

A comparison will be made between different techniques that exploit the information within low-loss EELS-spectrum images (SI). In this sense, the generation of maps from measured properties on the spectrum, such as characterization of the plasmon peak and relative thickness from the measured spectra was complemented with segmentation of the EELS-SI using mathematical morphology (MM) and a detailed exploration of spectral factorization using multivariate analysis (MVA).

Plasmon energies determined at the EELS-SI reveal the approximate spatial distribution of the Si-NCs and barrier dielectric material (SiO2, SiC and Si3N4, depending on the case). This method is better suited than the examination of the HAADF images, because of the appearance of spurious features from the inhomogeneity of the sample, masking the Si-NC positions (see Fig. 1 and 2). Nevertheless, it was not possible to get a direct measurement of the pure contribution of the Si-NC to the spectra, as all measured data present at least a mixture of nanoparticle and substrate plasmon. Fitting these two peaks using a double plasmon model (DPM) is reliable only when they are well separated in energy and exhibit significant differences in FWHM, i.e. low energy narrow peak vs high energy wide peak (as in Si-NCs in a SiO2 substrate) [1]. However, for other non-favorable situations, segmentation of the EELS-SI by MM can be of help. Following this scheme, averages of the spectra in the particle and dielectric areas can be generated, along with slices of the EELS-SI. These slices are then analyzed using MVA algorithms (NMF and BLU) for a factorization of the EELS data (see Fig. 3).

The collection of computational tools enabling nanometric spatial resolution imaging of the Si-NCs using sub-eV energy resolution EELS will be presented. Maps of measured properties, such as mean free path to sample thickness ratio, will be plotted for the three studied systems with different dielectric barriers. Moreover, the extraction of particular features by segmentation and factorization of the EELS data will allow recovering the pure Si-NC plasmon in each sample. Finally, the possibility of extracting electro-optical properties by thickness-normalized Kramers-Kronig analysis of the spectra will be explored.

[1] A. Eljarrat et al. (2013) Nanoscale 5, 9963-9970


Fig. 1: HAADF (upper left panel) and EELS (blue dashed lines, lower left panel) simultaneously acquired of a SiC sample. The EELS is analyzed to form the plasmon energy map (central panel, with thresholded histogram at left). Si-NC and SiC regions are marked off in this map and the average EELS are overlayed to the raw EELS (black=Si-NC, red=SiC).

Fig. 2: Results from the SiO2 sample, showing the superior sensitivity of the plasmon energy map above the HAADF and relative thickness map. Si-NC and SiO2 positions are marked off in the map and in the histogram (lower panel) as thresholds, using the same color code as Fig. 1.

Fig. 3: MVA factorization results vs. average EELS from the same EELS-SI shown in Fig. 2. After segmentation of the upper Si-NCs region, factorization reveals two different nanoparticles, and their contribution to EELS (comp. 2) is separated from the background SiO2 spectra (comp. 1).

Type of presentation: Poster

IT-5-P-1978 Plasmon and dielectric function mapping of multiple InGaN QW sample by HAADF-EELS.

Eljarrat A.1, López-Conesa L.1, Magén C.2,3, García-Lepetit N.4, Gačević Ž.4, Calleja E.4, Estradé S.1,2, Peiró F.1
1LENS-MIND-IN2UB, Departament d'Electrònica, Universitat de Barcelona, c/ Martí i Franqués 1, 08028 Barcelona, Spain., 2LMA-INA, Departamento de Física de la Materia Condensada, Universidad de Zaragoza, 50018 Zaragoza, Spain., 3Fundación ARAID, 50018 Zaragoza, Spain., 4TEM-MAT, Centres Científics i Tecnològics (CCiT), Universitat de Barcelona, Solís Sabarís 1, Barcelona, Spain., 5ISOM, Universidad Politécnica de Madrid, Ciudad Universitaria s/n, 28040 Madrid, Spain
aeljarrat@el.ub.edu

A thorough examination of InGaN quantum wells (QWs) was carried out through high angle annular dark field (HAADF) and electron energy loss spectroscopy (EELS) by aberration corrected scanning electron microscopy (STEM). The considered nanostructure consists of ~1.25 nm QWs layers with 20% indium content, periodically distributed along ~6 nm InGaN barriers, with 5 % indium content.

High resolution HAADF images give structural information from the examined crystal lattice below the nanometer range. Z-contrast in these images reveals the position of the QWs, the occurrence of In diffusion or even the formation of In-rich islands. Moreover, this information can be exploited using geometric phase analysis (GPA) algorithms in order to obtain maps of the deformation along lattice directions. The resulting deformation maps reveal that the structure suffers a localized distortion along the growth axis related to the In-rich QW positions (see Fig. 1).

STEM-EELS spectrum images (SI) are used to gain insight into the material properties of the sample. For this purpose, maps of plasmon peak energy and width are obtained and compared with the HAADF images. Generally, for III-V materials, compositional information can be recovered from the analysis of the plasmon peak energy through Vegard law [1]. In the present case, the small size of the examined QW and plasmon delocalization make this approach difficult to apply. However, the analysis of the plasmon witdth reveals a consistent swelling of the peak related with the position of In-rich regions, along with some expected shift to higher energy (see Fig. 2).

Furthermore, Kramers-Kronig analysis (KKA) of the EELS allows recovering the complex dielectric function (CDF) which contains electro-optical information from the material. For instance, it is possible to calculate the electron effective mass (m*) from the recovered CDFs at each pixel of the EELS-SI. The obtained values of m*, ranging from 0.14·m0 to 0.18·m0 are among the expected for InGaN (m*GaN = 0.2·m0 m*InGaN=0.11·m0). Moreover, depression regions in which the m* values are consistenty lowered are found in regions related with the ones having wider plasmon peaks (see Fig. 3).

All the computational work has been performed using the Hyperspy toolbox. The collection of techniques that have been developed in order to perform these analyses, will be presented along with the obtained results.

[1] A. Eljarrat et al. (2012) Microsc. Microanal. 18, 1143-1154.


Fig. 1: Left panel, HAADF-STEM image of the structure, revealing the position and width of an In-rich QW. Right panel, deformation in the growth direction calculated from the previous image by GPA.

Fig. 2: HAADF image and results from the analysis of the simultaneously acquired EELS-SI on an In-rich QW region. The plasmon analysis reveals a shift towards higher energies (Ep) and a swelling (Γ) of the plasmon peak around the QW. The effective mass (m*) shows a characteristic depression around the same region as the swelling in Γ.

Fig. 3: The upper panel shows two average EELS in the narrow Γ (dashed line) and wide Γ (solid line) regimes. Also following this line code, the lower panel shows the real (red) and imaginary (black) parts of the average CDF from these same regions.

Type of presentation: Poster

IT-5-P-1986 Advantages of FEG-EPMA for Microphase Analysis in Nuclear Materials

BRACKX E.1, NONNET H.2, HOMBOURGER C.3, ALLEGRI P.1, DUGNE O.1
11 CEA, DEN, DTEC, SGCS, LMAC, Marcoule, 30207 Bagnols sur Cèze, France, 22 CEA, MAR, DTCD, SECM, LDMC, Marcoule, 30207 Bagnols sur Cèze, France, 33 CAMECA, 29 Quai des Grésillons, 92622 Gennevilliers, France
emmanuelle.brackx@cea.fr

Many nuclear materials include micrometer-scale particles or phases that require characterization to understand and improve material fabrication techniques and processes. The microphase inclusions can be characterized with better precision thanks to the new generations of Field Emission Gun Electron Probe MicroAnalyzers (FEG-EPMA) whose high-resolution electron beam operating at low voltage optimizes and reduce electron interaction volume [1] [2].
An example of characterization is presented in relation with a study of the crystallization of certain phases in a glass matrix used for nuclear waste applications. Depending on the glass composition and the melting and cooling conditions, crystals can form in the matrix. The characterization of these phases, often of micrometer size, and of the glass including them are primordial in basic studies to elucidate the mechanisms involved. The composition of the including glass can be characterized by means of a CAMECA SX 100 electron microprobe with a low-resolution electron beam. On the other hand, the microcrystals studied here, apatite containing rare earth elements and microparticles of platinum-group metals, require the use of FEG-EPMA to determine their chemical composition because of their small dimensions (less than 10 µm: Figure 1). The analyses were carried out with the CAMECA SX 100 and SXFiveFE at 12 keV and 10 nA. The improved analytical resolution obtained with the CAMECA SXFiveFE made it possible to optimize the analysis of the micro particles, and to determine their chemical formulas.
A second example of characterization is presented in the context of coating materials used in nuclear processes. The coatings must have satisfactory homogeneity to ensure material adhesion and sealing. Any chemical homogeneity defects must then be characterized in order to optimize the manufacturing process. In this context a sealing material containing impurities in the form of micrometer-scale layered inclusions was characterized by EPMA (Figure 2). The analyses were carried out with a CAMECA SX 100 and CAMECA SXFiveFE at 15 keV and 10 nA. A comparison of the results shows the optimization obtained with the CAMECA SXFiveFE due to the high resolution of the electron beam.
REFERENCES
[1] X. Llovet, E. Heikinheimo, A. Núñez Galindo, C. Merlet, J. F. Almagro Bello, S. Richter, J. Fournelle, J. G. van Hoek. Materials Science and Engineering, 32 (2012).
[2] D. E. Newbury. Journal of Research of the National Institute of Standards and Technology, 107, 605-619 (2002).


Fig. 1: Microparticles in a nuclear glass matrix

Fig. 2: Stratified material with a thin vein

Type of presentation: Poster

IT-5-P-1995 EDS Element Analysis with High Count Rates

Eggert F.1
1EDAX Inc., AMETEK Materials Analysis Division, Mahwah NJ, USA
frank.eggert@ametek.de

One major benefit of Silicon Drift Detectors (SDD) with count rates of X-ray acquisition was already highlighted with a high speed 250 kcps map in first publication about application for Scanning Electron Microscopes [1] . The high count rates are usual with X-ray spectral imaging in each SEM lab more than 15 years later. The modern SDD spectrometers are able to process high count rates without significant deterioration of their basic spectrometric properties. It will be given a short overview about state of the art.

In the past, analysts have acquired single EDS spectra after selecting objects. This is now performed already with scanning the electron beam across entire specimen surface, which is usually very heterogeneous. A complete spectrum is acquired at each point. The count rates vary about a high range, depending from different specimen composition and topography. Reaching very high count rates of about 200 kcps and more are usual in daily praxis. Despite all SDD electronics are equipped with X-ray coincidence rejection logics, so called pile-ups will pass and are then in spectrum as artefacts. This is very fundamental, with all systems and does not depend from vendors, never possible to neglect. The artificial counts will produce mistakes in qualitative and quantitative analytical results if not considered. An outline will be given about the effects and how to process (Fig 1) [2]. It will be demonstrated the quantitative results are stable up to 200…300 kcps, if a pile-up consideration is applied (Fig 2) [3].

But correction comes with higher result uncertainties and detection limits. Also, the pile-up consideration is with limits, e.g. the entire region must be homogeneous, were all X-rays in a spectrum came from [5]. Otherwise fundamental assumption about pile-up consideration with random emission was violated. It is not satisfied if the electron beam excitation involves areas of different specimen constituent. Phase determination by independent Principle Component Analysis (PCA) is useful to identify homogeneous specimen regions to avoid qualitative and quantitative analytical errors. This would be if total spectrum was taken from inhomogeneous area (Fig 3) [5]. Specialized single pixel spectra evaluation strategies are required for full quantitative maps.

References

[1] Strüder L, Meidinger N, Stötter D, Kemmer J, Lechner P, Leutenegger P, Soltau H, Eggert F, Rohde M and Schülein T (1998) Microsc. Microanal. 4 622
[2] Eggert F, Elam T, Anderhalt R, Nicolosi J (2012) IOP Conf. Ser.: Mater. Sci. Eng. 32, 012008
[3] Eggert F, Anderhalt R and Nicolosi J (2012) Microsc Microanal 18 (Suppl.2)
[4] Eggert F (2010) IOP Conf. Ser.: Mater. Sci. Eng. 7 012007
[5] Eggert F, Schleifer M, Reinauer F (2014) IOP Conf. Ser.: Mater. Sci. (in publication)


Fig. 1: The results of automated qualitative analysis [4] with two spectra of same specimen (a low; b very high count rate) are similar due to internal pile-up consideration (pile-up distribution is not included in reconstruction, blue line). This is despite big differences are visible in both spectra caused by pile-up artefacts (example from [2]).

Fig. 2: The quantitative results vary with count rate, if pile-up was not considered (a). They are much more stable with using pile-up consideration method (b) (example from [3]).

Fig. 3: Phase map of Kiruna mineral with very high count rates. Different phase areas indicate from which pixel regions sum spectra are possible to gather without analytical evaluation issues. The spectrum is from an inhomogeneous area to demonstrate the qualitative analysis challenge, even if the pile-up consideration was applied (example from [5]).

Type of presentation: Poster

IT-5-P-2431 Experimental detection of the Čerenkov limit in Si, GaAs and GaP

Horák M.1, Stöger-Pollach M.2
1Institute of Physical Engineering, Brno University of Technology, Brno, Czech Republic, 2USTEM, TU Vienna, Vienna, Austria
horak.michal@seznam.cz

Since the advent of monochromated electron energy loss spectrometry (EELS) the experimental detection of band gaps in semiconducting materials is of great importance. But due to the fact that the swift electron probe excites relativistic energy losses, like Čerenkov losses [1] and the corresponding light guiding modes, the band gap is hidden below them. Therefore a technique was developed to excavate them mathematically [2]. Another possibility is to reduce the beam energy [3] such that the speed of the swift probe electron v does not exceed the speed of light inside the sample c0/n (with n as the refractive index).

The investigated specimens are Si, GaAs and GaP single crystals. Sample preparation was performed by grinding and ion milling using a low voltage ion mill for the final preparation step in order to remove surface damage from prior milling.

With the TECNAI G20 at TU Vienna we are able to record EELS spectra in the energy range of 6–100 keV. Consequently we can experimentally verify the Čerenkov limit of Si, GaAs and GaP, which is 13.3, 20.6 and 30.1 keV, respectively (Fig. 1). The probability of Čerenkov photon excitation per unit path length of the electron inside the specimen PCP is proportional to

PCP = 1/c0 – 1/(v2ε1),

with ε1 as the real part of the samples dielectric function [3]. The theoretical values are computed for PCP equals to zero. The shaded area in Fig. 1 represents the beam energy needed in order to excite 0.2 to 0.4 Čerenkov photons.

Above the limit Čerenkov losses and light guiding modes cause a red shift of the signal onset in low loss spectra. The signal onset is shifted to higher energies when reducing the electron beam energy and there is no shift below the limit (Fig. 2). It must be noticed, that only the direct gap at 3.6 eV of Silicon is measured.

 

[1] E. Kröger, Z. Physik 216 (1968) 115-135.
[2] M. Stöger-Pollach, A. Laister, P. Schattschneider, Ultramicroscopy 108 (2008) 439-444.
[3] M. Stöger-Pollach, Micron 39 (2008) 1092-1110.


M. H. acknowledges FEI Company for financial support.

Fig. 1: Theoretical Čerenkov-limit of beam energy for refractive index between 1.5 and 6.5. The shaded area represents the beam energy needed in order to excite 0.2 to 0.4 Čerenkov photons.

Fig. 2: Zero loss deconvolved low loss spectra of Silicon at various beam energies. Below the Čerenkov limit at 13 keV no shift of the signal onset can be observed. Above the Čerenkov losses and light guiding modes cause a red shift of the onset.

Type of presentation: Poster

IT-5-P-2062 Nanoscale Luminescence Mapping of InGaN/GaN Multiple Quantum Dot Doped Nanowire LEDs with Scanning Transmission Electron Microscopy

Woo S. Y.1, Kociak M.2, Nguyen H. P.3, Mi Z.3, Botton G. A.1
1Department of Materials Science & Engineering, Brockhouse Institute for Materials Research, and Canadian Centre for Electron Microscopy, McMaster University, Hamilton, ON, Canada, 2Laboratoire de Physique des Solides, Université Paris-Sud XI, Orsay, France, 3Department of Electrical & Computer Engineering, McGill University, Montreal, QC, Canada
woosy@mcmaster.ca

Ternary InGaN compounds show great promise for light-emitting diode (LED) applications because of bandgap energies (0.7–3.4 eV) that can be tailored to have emission wavelengths spanning the entire visible spectral range. Complex III-N device heterostructures have been incorporated into GaN nanowires (NWs) recently, but exhibit emission linewidths that are broader than expected for their corresponding planar counterparts, as measured with photoluminescence (PL) spectroscopy. It is thus critical to understand how the structural and optical properties interplay, using spectroscopic methods that can resolve localized signals at the nanoscale.
Multiple InGaN/GaN quantum dot (QD) embedded nanowire (NW) LED structures, grown on Si(111) substrates by molecular beam epitaxy, were characterized by STEM. Elemental mapping using EELS has shown a systematic non-uniformity of the In-content between the InGaN QDs that are centrally confined within the active region, embedded between n- and p-doped GaN in the NW LED structure (Fig. 2(c,d)). To correlate these observations to the inhomogeneous broadening observed in PL from an ensemble of NWs, nm-resolution STEM-cathodoluminescence (CL) spectral imaging on single NWs was performed using a custom-made system on a VG HB-501 STEM as described in [2]. Individual NWs examined showed diverse optical responses, but most NWs exhibit one main emission peak centered at 500–550 nm in the yellow-green. Spectral features consisting of multiple sharp peaks (25–50 nm at FWHM) spanning a wavelength range of ~100 nm arise from the active region (Fig. 1(b)), showing an apparent spatial dependence of the spectral shifts (Fig. 1(a)). This is consistent with the PL, indicating that the broad emission originates from within single NWs and is not an inhomogeneous broadening. However, typical wavelength-integrated CL mapping was too ambiguous in the spatial assignment of some peaks that have overlapping intensities. Improved spatial-spectral correlation was achieved by inspecting orthogonal spatial slices from the spectrum image singly along x and y (Fig. 2(a,b)) to define various combined position and wavelength maxima. Multiple optical signals of varying emission wavelengths arising from well-defined locations within the QD active region were identified, and can be attributed to the observed In-content variation between successive QDs. Lastly, the evidence of localized emission intensity in the QDs towards the p-GaN, likely due to the diffusion of charge carriers generated by the electron beam, could suggest the accumulation of carriers within the active region towards the p-GaN.
[1] H.P.T. Nguyen et al., Nano Lett., 12(3), 1317-1323 (2012)
[2] Zagonel et al., Nano Lett., 11(2), 568-573 (2011)


This work was supported by the Natural Sciences and Engineering Research Council of Canada (NSERC).

Fig. 1: STEM-CL spectrum image (SI) of the NW structures. (a) HAADF and BF image acquired simultaneously with the CL, and spatial maps of spectral features centered about the wavelengths labeled. The three marked regions of interest (ROI) that exhibit unique emission spectra are shown in (b). (c) HAADF image to better resolve the same NW studied using CL.

Fig. 2: Spatial-spectral plots of the SI from Fig. 1, (a) across the SI in the y-axis, (b) along the SI in the x-axis with concurrent HAADF signal overlaid to show the structure; CL intensity is color-coded. (c, d) HAADF image and corresponding STEM-EELS In-map of the boxed area in another NW, showing the varying In-content in the 10 InGaN QDs.

Type of presentation: Poster

IT-5-P-2077 Application of Energy Filtering in STEM (EFSTEM) mode for Mapping of Elements and Chemical Bindings

Muehle U.1, Gluch J.1,2, Zschech E.1,3
1Fraunhofer Institute for Ceramic Technologies and Systems - Materials Diagnostics (IKTS-MD), Maria-Reiche-Str. 2, 01109 Dresden, Germany, 2Institute for Materials Science, TU Dresden, 01062 Dresden, Germany, 3Dresden Center for Nanoanalysis (DCN), TU Dresden, 01062 Dresden, Germany
uwe.muehle@ikts-md.fraunhofer.de

During the previous decade, Scanning Transmission Electron Microscopy (STEM) has gained a growing importance, enforced by the availability of TEMs with high-performance Cs correctors for the condenser system [1, 2]. One major benefit of STEM is that additional signals like X-ray emission (EDX) or energy loss of transmitted electrons (EELS) can be acquired with the same local resolution as the image of a sample.

The energy filtering in the STEM mode provides an improved time-to-data (or time-to-result) at a spatial resolution which is sufficient for a lot of application cases. Compared to the acquisition of complete EEL spectra at every pixel which leads to a long measuring time and a large data volume, , the proposed technique allows to obtain a highly resolved elemental distribution without leaving the STEM mode of the instrument [4].

The energy-filtered STEM data can be acquired using an in-column filter in combination with a BF/DF detector or a HAADF detector, positioned in the electron-optical path behind the filter (Fig. 1). After passing the energy selecting slit, the beam contains only electrons of the chosen energy range, and the acquired signal is comparable with that of the well-known EFTEM method [4].

One application is the improvement of the image quality by removing the inelastic scattered electrons. Shifting the energy of the primary beam allows to acquire STEM images using an energy window with a defined energy loss in the low loss or the core loss region. A combination of several images allows the application of the 3-windows method similarly to EFTEM or of the jump-ratio-method [2] (Fig. 2). The acquisition of a series of images with an energy window of about 2 eV and with stepwise increasing energy loss enables a detailed characterization of chemical bindings – either in the plasmon range of the EEL spectrum or above the ionization edge of an element.

Advantages of this technique are a better utilization of the available beam intensity, which is often weak for large energy losses. This approach enables to improve the focus of the image for large energy losses. In addition, some materials show less electron-beam damage in case of STEM imaging [5]. For these materials, the described technique is the technique of choice to avoid long illumination times as they are necessary for EFTEM. Finally, a sample drift does not influence the results as much as in the TEM mode.

Pennycook, S.J. Scanning Transmission Electron Microscopy Springer 2011

Brydson, R. Aberration-Corrected Analytical Transmission Electron Microscopy RMS 2011

Egerton, R.F. Electron energy-loss spectroscopy in the electron microscope Springer  2011

Muehle, U. et.al. Patent.; 10 2013 011 674.0 2013

Yeap, K.B. et al. IEEE IRPS 2013


Fig. 1: Schematic of an In-column STEM with HAADF-Detector behind the filter

Fig. 2: Elemental mapping of a semiconductor structure, acquired by energy filtered STEM and evaluated using the jump-ratio method for the elements nitrogen (a), oxygen (b) and titanium (c)

Type of presentation: Poster

IT-5-P-2130 Plasmon Mapping in Au@Ag Nanocube and their assemblies

Goris B.1, Guzzinati G.1, Fernández-López C.2, Pérez-Juste J.2, Liz-Marzán L. M.2,3,4, Trügler A.5, Hohenester U.5, Verbeeck J.1, Bals S.1, Van Tendeloo G.1
1EMAT, University of Antwerp, Groenenborgerlaan 171, 2020 Antwerp, Belgium, 2Departamento de Química Física, Universidade de Vigo, 36310 Vigo, Spain, 3BioNanoPlasmonics Laboratory, CIC biomaGUNE, Paseo de Miramón 182, 20009 Donostia , San Sebastián, Spain, 4Ikerbasque, Basque Foundation for Science, 48011 Bilbao, Spain, 5Institut für Physik, Karl-Franzens-Universität Graz, Universitätsplatz 5, 8010 Graz, Austria
bart.goris@uantwerpen.be

Plasmons are collective excitations of conduction electrons in metallic particles. For nanostructures, the resonant surface plasmon modes are highly sensitive to the geometry of the structure and can therefore be tuned by controlling their morphology and/or size. Here, we applied monochromated STEM-EELS to map the surface plasmon resonances in Au@Ag nanocubes and their assemblies. [1]

These assemblies reveal interesting plasmonic properties with an increased flexibility as compared to their single particle counterparts.
For the isolated nanocubes, EEL spectra were recorded at different locations, revealing the presence of three distinct plasmon resonances at energy values of 2.2 eV, 3.2 eV and 3.5 eV. As presented in Figure 1, the extracted plasmon maps indicate that the two modes with the lowest energy have the highest probability to be excited at the corners of the particles, whereas the third mode is best excited at the side faces, in agreement with previous reports. [2]

Interestingly, when the nanocubes are dispersed on a C support, they tend to orient themselves side by side yielding regular assemblies. As a first example, 3 nanocubes may form an approximately triangular array as shown in Figure 2. It can be observed that the main plasmon modes are obtained at energy losses of 1.2 and 1.6 eV and are in qualitative agreement to the plasmon modes of a perfect nanotriangle with the same dimensions. [3] Plasmon modes that occur at higher energy losses originate from the deviation of the overall shape from a perfect triangle, resulting in multiple regions of high intensity that are mainly located at the corners of the individual cubes. A nanotriangle constructed by the random ordering of multiple nanocubes could possibly act as the first half of a bow-tie antenna. This antenna has great potential due to the field enhancement in the central region between the two triangles caused by plasmon coupling. [4] As illustrated in figure 3, even with the simplified geometry for a bow-tie antenna, the field enhancement due to the coupling can be clearly visualized in the centre of both the experimental and the simulated plasmon maps. The enhanced field occurs at an energy loss of 1.3 eV which can also be observed as an increased probability for the energy loss in the acquired EEL spectra.


[1] S. Gómez-Graña, J. Phys. Chem. Lett. 4 (2013) p.2209
[2] O. Nicoletti et al., Nature 502 (2013) p.80
[3] J. Nelayah et al., Nature Physics 3 (2007) p.348
[4] A. Koh et al., Nano Letter 11 (2011) p.1323


The authors acknowledge support from the European Research Council and the FWO.

Fig. 1: Low loss EEL spectra of a Au@Ag nanocube showing three distinct major plasmon modes (a-c). The first two modes have the highest possibility to be excited at the corners of the cubes whereas the third one is best excited at the edges.

Fig. 2: Surface plasmon modes of a triangle comprising three Au@Ag core-shell nanocubes. (A and B) The EELS spectra were acquired at the positions indicated by the dots of the corresponding colors. When inspecting these spectra, several modes are observed, which are in agreement with BEM simulations of a perfect triangular shape.

Fig. 3: Surface plasmon maps of a bow-tie antenna created by a specific order of self-assembled nanocubes. Both the acquired EELS spectra (A) and the experimental and simulated near field maps (B-G) show a large field enhancement in the region between both nanocube structures.

Type of presentation: Poster

IT-5-P-2151 Non-stoichiometry and order-disorder in the SbxV1-xO2 (0<x<0.5) solid solution

Landa-Cánovas A. R.1, Vilanova-Martínez P.1, Agulló-Rueda F.1, Hernández-Velasco J.1
1Instituto de Ciencia de Materiales de Madrid, ICMM-CSIC. Sor Juana Inés de la Cruz, 3; 28049 Madrid, Spain
landa@icmm.csic.es

~SbVO4 plays a key role in the catalyst for the ammoxidation of propane to acrylonitrile. Besides, it exhibits an amazing structural flexibility involving cation vacancies, changes in oxidation states and different degrees of order-disorder, ranging from Short Range Order (SRO) to periodic superstructures and structural modulations [1,2]. In this work we have studied the solid solution that ranges from VO2 to SbVO4 according to the following reaction stoichiometry:

[Sb2O3 + V2O5] + VO2 ---> SbxV1-xO2

 During the reaction all Sb3+ cations are oxidized to Sb5+ while all V5+ cations are reduced to V3+. This implies the following substitution in the basic rutile-type VO2 matrix: Sb5+ + V3+ <---> 2V4+. In this way we have been able to synthesize a whole solid solution SbxV1-xO2 ranging from SbVO4 (x=0.5) to Sb0.1V0.9O2 (x=0.1) by heating the stoichiometric amounts of Sb2O3, V2O5 and VO2 at 800ºC under argon atmosphere. In the Sb-richest phase, SbVO4, most of the vanadium is V3+, as confirmed by EELS spectroscopy, magnetic susceptibility and neutron diffraction, showing the latter magnetic ordering at TN < 50K. Electron diffraction shows the presence of intense SRO in the form of wavy two-dimensional sheets of diffuse intensity in the reciprocal space, see Fig. 1. HRTEM demonstrates that the SRO is due to low correlation between ...Sb-V-Sb-V... chains running along c. This SRO disappears at Sb0.33V0.67O2 and magnetic ordering happens at lower temperatures (TN~6K). For Sb0.25V0.75O2 compositions, a different type of SRO appears and its intensity increases as the amount of Sb decreases, Fig. 2. At Sb0.1V0.9O2 this SRO can also be observed by electron diffraction as very intense two-dimensional sheets of diffuse intensity forming a three-dimensional net of edge-sharing octahedra in reciprocal space. This phase presents a structural transition similar to that of VO2 but at lower temperature (51ºC). A phase transition can be observed by electron diffraction (Fig. 3) when heating with the electron beam from Short Range Order (SRO, diffuse lines) (a) to a two-fold superlattice (b) and back to SRO (c) as the temperature is lowered. During the whole series we range from SbVO4 with V3+ to VO2 with V<4+ and that big change in stoichiometry has been accommodated in a "soft way" through SRO mechanisms.

[1] A.R. Landa-Cánovas, J. Nilsson, S. Hansen, K. Staahl and A. Andersson. J. Solid State Chem.116, 369-377 (1995); [2] A. R. Landa-Cánovas, F. J. García-García, S. Hansen. Catalysis Today 158, (2010) 156.


Authors thank Spanish Government (project FLEXOCAT-MAT2011-27192) for financial support.

Fig. 1: SAED patterns of a crystal of Sb1.0V1.0O4 showing sharp rutile maxima and wavy diffuse planes indicating Short Range Order between the cations.

Fig. 2: SAED patterns of Sb0.25V0.75O2 showing diffuse scattering produced by Short Range Order of the cations. Note that the shape of the diffuse scattering is very different to the one observed in Sb1.0V1.0O4, see Fig. 1.

Fig. 3: SAED patterns of the phase transition observed at the Sb0.1V0.9O2 sample, changing as the crystal temperature is raised with the electron beam over 51ºC from SRO (diffuse lines) (a) to a two-fold superlattice (b) and back to SRO as the temperature is lowered. SAED patterns are misaligned to increase the intensity of the SRO diffuse layers.

Type of presentation: Poster

IT-5-P-2159 Developments in FESTEM & EDS for the Characterisation of Dopant Distributions within Advanced Semiconductors

Dijkstra H.1, Thompson K.1, Stephens C. J.1
1ThermoFisher Scientific
chris.stephens1@thermofisher.com

Recent years have seen an acceleration in the pace of semiconductor development, driven by an ever increasing demand for high-performance, low-cost electronic devices. The ability to mass produce complex structures on a nanometre scale is critical to this process, requiring reliable characterisation techniques to measure and control key parameters of interest. The concentration and distribution of dopant atoms within semiconductors directly affects device performance, requiring analytical electron microscopes capable of accurate elemental quantification on a nanometre scale. This work focuses on the challenges in analysing advanced semiconductor structures, and the developments in microscope technology, EDS detectors and post-acquisition analysis routines which make this possible.

Figure 1 left shows a HAADF STEM image of an As/ P dopant distribution within a NMOS transistor, characterised using a JEM-2800 transmission electron microscope (TEM), a JEOL 100 mm2 (solid angle = 0.95 Sr) silicon drift detector (SDD) in conjunction with the NORAN System 7 microanalysis platform. The concentration of the dopant regions is low (<0.1%), within small regions (<5% area), as shown in the cumulative spectra in Figure 1 right. Whilst quantitative elemental mapping (peak deconvolution and background subtraction) eliminates many of the problems associated with traditional elemental analysis (Figure 2), such as overlapping peaks, many hours can be required to acquire statistically significant data. Furthermore, determining phases from such ‘Quant’ maps often results in end-user bias and the misidentification of chemically unique phases.

COMPASS is an ideal tool for EDS analysis under such extreme conditions, utilising multivariate statistical analysis (MSA) in order to extract the principle components of the spectra at each pixel and group statistically similar phases. Figure 3 left shows the composite phase map using COMPASS, Figure 3 centre shows the principle component map of As Doped Si and figure 3 right overlays spectra of each component. COMPASS extracts spectra relevant only to the specific phase of interest, enhancing the signal-to-noise ratio in comparison to the quant maps and assists in the detection of trace elements within a phase of interest. The end result is the significant reduction in acquisition time and the detection of physically significant phases Figure 4, otherwise missed by conventional spectrum-based phase approaches.

This work is set in the wider context of developments in analytical electron microscopy and the role this plays in improving advanced manufacturing processes.


Fig. 1: Figure 1 Left HAADF image of analysed area of a NMOS transistor. Right Cumulative spectrum for all pixels in the data set.

Fig. 2: Figure 2 Left Conventional peak count. Right quantitative elemental map of selected elements within the NMOS device. Ta is incorrectly displayed on a conventional elemental map due to overlaps with Ha

Fig. 3: Figure 3 Left Composite COMPASS element map of Si regions. Centre Principle component map of As doped Si Right selected area specta of Si phases.

Fig. 4: Figure 4 Left Composite map of non-Si phases. Right Spectra of principal components 10 and 12. Component 12 was not uniquely identified during spectrum based phase analysis

Type of presentation: Poster

IT-5-P-2324 Band gap measurements of ultra thin buried films using conical darkfield EFTEM and low voltage EFTEM

Stöger-Pollach M.1, Biedermann K.2, Beyer V.2
1USTEM, TU Vienna, Vienna Austria, 2Fraunhofer IPMS-CNT, Dresden, Germany
stoeger@ustem.tuwien.ac.at

New electronic devices require new techniques for characterization. We investigate a c-Si/a-SiOx/a-SiON/a-SiOx/pc-Si (SONOS) stack as used in flash memory devices by using valence electron energy loss spectrometry (VEELS) and energy filtered transmission electron microscopy (EFTEM).
In the present work we discuss the given physical limitations, which include relativistic energy losses – like Čerenkov losses – and the wide range Coulomb interaction. Whereas the first effect can alter the VEELS spectrum and can be easily avoided by reducing the beam energy, the latter affects the spatial resolution of the inelastically scattered electrons.
Although the range of the Coulomb interaction is smaller for slower electrons, the spatial resolution of the low voltage (LV-) EFTEM method will still be limited by this effect. In the case of conical dark field (conDF-) EFTEM we are going to locate the collection aperture in the reciprocal plane such, that we do not collect the small angle dispersed Čerenkov losses and such that we probe the indirect gap of Silicon. Still this method is critical, because its results do not give the optical properties. This is because the measurement is performed at q ≠ 0.
The EFTEM data cubes are recorded with a TECNAI G20. For the respective experiments we chose 40 keV for the LV-EFTEM experiment and 200 keV for the conDF-EFTEM experiment. In conDF the incoming electron beam is deflected by the Bragg angle of Si(111) and conically rotated during the EFTEM acquisitions. Therefore only dark field signal is used for the data cube.
Although the 40 keV experiment does not show Čerenkov losses inside the oxide-oxynitride-oxide (ONO) stack, it still shows some intensity in the direct gap of Silicon. Anyhow, the inelastic delocalization hinders an extraction of an SiO2 EELS signal. The major components of the spectrum extracted from the SiO2 layer positions are due to Si and SiON. The measured band gaps are 3.8 eV in SiON, 3.8 eV in SiO2, and 1.8 eV in Si (although the direct gap at 3.4 eV should be probed). They are all wrong due to delocalization and Čerenkov loss excitation.
In the case of the 200 keV conDF-EFTEM experiment, the spectra can be extracted quite well, although the SiO2 spectrum still suffers slightly from inelastic delocalization. The measured band gaps are 4.9 eV in SiON, 6.3 eV in SiO2, and 1.3 eV in Si (which is the indirect gap being probed under conDF conditions). The value for SiO2 still suffers from delocalization.
We demonstrate that the determination of optical properties of low is a problem with EELS as soon as the layer thickness is smaller than the inelastic delocalization. Probing the band gap can be done under the restrictions of a dark field experiment measuring q ≠ 0.


The authors acknowledge the USTEM facilities for providing the low-kV TEM.

Fig. 1: Figure 1: (a) Defocused bright field image of the area of interest. The ONO stack of the SONOS-transistor is located between the substrate and the pc-Si gate electrode. (b) elemental map of Silicon (green), Oxygen (blue) and Nitrogen (red). (c) HRTEM image of the ONO stack.

Fig. 2: Figure 2: (a) 40 keV spectrum image across the ONO stack extracted from the LV-EFTEM data cube. (b) 200 keV spectrum image across the ONO stack extracted from the conDF-EFTEM data cube. (c) Spectra of Si, SiO2 and SiON extracted from the corresponding data sets.

Type of presentation: Poster

IT-5-P-2163 An ELNES study of anisotropic materials using variable beam energies

Stöger-Pollach M.1, Hetaba W.1, Rodemeier R.2
1USTEM, TU Vienna, Vienna, Austria, 2GATAN GmbH, München, Germany
stoeger@ustem.tuwien.ac.at

The development of (S)TEMs in recent years is pointing towards a higher variability of beam energies. The reasons are manifold, as there are amongst others less beam damage [1], larger elastic and inelastic scattering cross sections, and less or even no excitation of Cerenkov losses for the analysis of optical properties [2,3]. Simultaneously the development of electron detectors being able to handle slower electrons efficiently allows detecting low signals with little noise. This is important, because low voltage electron beams have usually little beam current.
In the present work we demonstrate the effect of the varying beam energy on the ELNES of the B-K edge caused by the change in momentum transfer with respect to the forward (qz) and perpendicular (qperp) directions. With decreasing beam energy qperp is stronger decreasing than qz. For the experiment we orient the c-axes of a hexagonal BN crystal parallel to the electron beam. For low beam energies the van der Waals bonds contribute to the EELS spectrum stronger as compared to high beam energies showing a small qz. This can be seen in Figure 1, because the π* peak decreases with increasing beam energy. When orienting the beam axes perpendicular to the c-axes of h-BN the effect is inverted and the sp2 orbitals are contributing stronger with decreasing beam energy (see Figure 2).
We compare the experiments with ab initio calculations using the Wien2k code. Using the TELNES.3 routine, the effects of orientation dependence [4] as well as the influence of the beam energies on the fine structure of the Boron K-edge can be simulated. One can investigate orbital dependent properties by comparing these simulations to the experimentally acquired spectra. As changing the beam energy is in some cases a much easier task to perform than conducting ELCE experiments [5], this technique can be an alternative or a complementary procedure.

[1] U. Kaiser et al., Ultramicroscopy 111 (2011), 1239 - 1246
[2] M. Stöger-Pollach, Micron 39 (2008), 1092 - 1110
[3] M. Stöger-Pollach, Micron 41 (2010), 577 – 584
[4] C. Hebert-Souche et al., Ultramicroscopy 83 (2000), 9 – 16
[5] W. Hetaba et al., Micron, in press.


The authors aknowledge the USTEM facility for providing the low-KV TEM.

Fig. 1: Figure 1: a) Boron-K edge of BN recorded with various incident beam energies (normalized to the 195.8 eV peak). b) Bright field image of the BN specimen. c) Diffraction pattern of h-BN showing the [0001]-orientation. Consequently the c-axes of the h-BN is parallel to the beam axes.

Fig. 2: Figure 2: a) Boron-K edge of BN recorded with various incident beam energies (normalized to the 195.8 eV peak). b) Bright field image of the BN specimen. The circle indicates the position of the EELS experiments. c) Diffraction pattern of h-BN showing the (0002) spots only. Consequently the c-axes of the h-BN is perpendicular to the beam axes.

Type of presentation: Poster

IT-5-P-2169 QW emission shift along single InGaN/GaN core-shell LEDs evaluated by monochromatic cathodoluminescence image series

Ledig J.1, Fahl A.1, Popp M.1, Scholz G.1, Steib F.1, Wang X.1, Hartmann J.1, Mandl M.1,2, Schimpke T.1,2, Strassburg M.2, Wehmann H. H.1, Waag A.1
1Institut für Halbleitertechnik, Technische Universität Braunschweig, Hans-Sommer-Str. 66, 38106 Braunschweig, Germany, 2OSRAM Opto Semiconductors GmbH, Leibnizstr. 4, 93055 Regensburg, Germany
j.ledig@tu-bs.de

Three dimensional light emitting diodes (LEDs) with a shell geometry of p-GaN and InGaN multi quantum well (MQW) around a columnar n-GaN core are supposed to have distinct advantages over conventional planar LEDs. The active area along the sidewalls of the GaN pillars can substantially be increased by high aspect ratios - leading to a lower current density inside the MQW at the same operating current per substrate area.
The investigated core-shell LED structures are grown by selective area metal organic vapor phase epitaxy on templates consisting of a patterned SiOx mask layer on an n-type GaN layer on 2” sapphire wafers. Due to the 3-dimensional shape, the optical properties of the QWs in each structure show significant gradients along the height and variations between different facets on a micrometer scale.
The automatized capturing of monochromatic CL image series is realized by combining a scan generator (controlling the electron probe position on the sample) and monochromator control (grating rotation angle). Hyperspectral imaging as well as spectra from selected areas are generated by post processing of such image series with respect to the spectral sensitivity of the optical system - including the collection optics, monochromator and detector. The spatial 2-dimensional (2D)-resolution of the presented method is higher than that of using a parallel detector for capturing CL spectra for distinct points of excitation. In parallel, the flexibility of the wavelength range gives a benefit for hyperspectral investigation in different regions of emission from the 3D InGaN/GaN LEDs.
The CL shows near band edge emission (NBE), signals from the QWs and defect related yellow luminescence (YL) which proves that the InGaN is present on all sidewall facets. The properties of neighbor structures are similar – although their diameters might be different. By contacting a single facet with a tungsten probe tip electroluminescence (EL) spectra are obtained at different injection currents.
Investigation of those structures by electron beam induced current (EBIC) proves the geometry of a p-shell around an n-core. Similar to the CL intensity the spatially resolved EBIC analysis indicates how the generation rate is affected by topography (edges) and that the top facets show different properties. A wavelength shift of the MQW emission of 60 nm is observed along the structure height for both excitations by the electron probe (CL) and by a local current injection (EL). This shift is assigned to a gradient of the indium incorporation caused by diffusion mechanisms during growth.


We thank Dr. Uwe Jahn for support regarding optical characterization. The financial support of the European commission (SMASH and GECCO) and the endorsement of the NTH and the JOMC are acknowledged.

Fig. 1: SE image and monochromatic CL images of InGaN/GaN core-shell LED structures on the cleaved growth template using a spectral FWHM of about 7.5 nm at an FOV = 11.4 µm, EHT = 15 kV, tilt = 30°.

Fig. 2: Logarithmic contour plot of CL spectra obtained by exciting small areas at different structure height (17 positions on the sidewall and 6 positions on the top facets), captured with a spectral FWHM of about 7.5 nm using a CCD parallel detector.

Fig. 3: CL and EL spectra captured with a spectral FWHM of about 7.5 nm using a CCD parallel detector. The single core-shell LED was excited at different heights on the sidewall by an electron probe of EHT = 15 keV and a current injection of I = 1.5 µA via a probe tip contact, respectively.

Type of presentation: Poster

IT-5-P-2232 Distinguishing overlapping EMCD signals on oxidized metals in the TEM

Thersleff T.1, Rusz J.2, Rubino S.5, Hjörvarsson B.2, Ito Y.3, Zaluzec N. J.4, Leifer K.1
1Department of Engineering Sciences, Uppsala University, Uppsala, Sweden, 2Department of Physics and Astronomy, Uppsala University, Uppsala, Sweden, 3Department of Physics, Northern Illinois University, DeKalb, IL, USA, 4Electron Microscopy Center, Argonne National Laboratory, Argonne, IL, USA, 5Department of Physics, University of Oslo, Oslo, Norway
thth@angstrom.uu.se

Energy-loss Magnetic Circular Dichroism (EMCD) is a powerful electron microscopy technique capable of extracting quantitative magnetic information from nano-sized features [1–3]. This technique has potential for high impact in research and industry where understanding interfacial magnetism is crucial to the design of nanostructured magnetic materials yet hampered by a lack of reliable small-volume characterization techniques. A key challenge facing further development of the EMCD technique is to extract reliable data from very small volumes of materials with sufficient quality for quantitative analysis. This is especially difficult for metals, as a thin oxide typically forms on the surfaces of the as-prepared lamella during transfer into the transmission electron microscope. In the case of iron, this surface oxide layer may be itself magnetic, potentially complicating the quantification of an EMCD signal from thin regions.

In this study, we investigate variations in the EMCD signal on an iron thin film with a surface oxide layer. The experimental design is depicted in figure 1 and yields a fully convergent beam with a diameter that can be reduced to approximately 1 nm. Under these conditions, we are able to provide a detailed structural and chemical assessment of the surface oxide. Subsequently we explore the consequences of its magnetization as well as how this modifies the detected EMCD signal (see figure 2). We conclude by proposing a method to quantify this effect and distinguish between the EMCD signals from either the underlying metallic film or its surface layer.
References
[1] P. Schattschneider, S. Rubino, C. Hébert, J. Rusz, J. Kuneš, P. Novák, et al., Detection of magnetic circular dichroism using a transmission electron microscope, Nature. 441 (2006) 486–488.
[2] P. Schattschneider, M. Stöger-Pollach, S. Rubino, M. Sperl, C. Hurm, J. Zweck, et al., Detection of magnetic circular dichroism on the two-nanometer scale, Phys. Rev. B. 78 (2008).
[3] H. Lidbaum, J. Rusz, A. Liebig, B. Hjörvarsson, P. Oppeneer, E. Coronel, et al., Quantitative Magnetic Information from Reciprocal Space Maps in Transmission Electron Microscopy, Phys. Rev. Lett. 102 (2009).


The authors acknowledge STINT research grant (1G2009-2017) and the Electron Microscopy Center at Argonne National Laboratory, a U.S. Department of Energy Office of Science Laboratory operated under Contract No. DE-AC02-06CH11357 by UChicago Argonne, LLC. J. R. acknowledges the Swedish Research Council, Göran Gustafsson's Foundation and Swedish National Infrastructure for Computing (NSC center).

Fig. 1: Figure 1 - Illumination conditions for the EMCD experiment. The probe size is limited by the C2 aperture and can be reduced to approximately 1 nm in diameter.

Fig. 2: Figure 2 – ELNES spectra acquired at two detector positions using a probe with a diameter of approximately 1.5 nm.  Their difference is an EMCD spectrum.

Type of presentation: Poster

IT-5-P-2235 The Use ofVery Large Area Detectors for fast Light Element Mapping and Data Acquisition in STEM

Rowlands N.1, Phillips P.2, Bhadare S.3, Klie R.2, Nicholls A.2
1Oxford Instruments NanoAnalysis, Concord, USA, 2University of Illinois, Chicago, USA, 3Oxford Instruments NanoAnalysis, High Wycombe, UK
neil.rowlands@oxinst.com

In recent years silicon drift detectors have become the logical choice for characteristic EDS X-ray analysis in both SEM and TEM. Large sensor sizes offer increased solid angle, allow data to be acquired faster in low signal regimes and with only Peltier cooling required, the need for liquid nitrogen is removed. High count rate situations are also handled more easily due to lower noise and excellent energy resolution.
Now ultra-large non circular SDD detectors have been designed for (S)TEMs.  With their non-traditional geometries, they minimize the detector to sample distance and give very high solid angles. These large EDS detectors are able to detect and map elements down to the atomic level whilst minimizing analysis times and thus limiting the effects of beam drift, beam damage and sample contamination.
By taking further advantage of the clean, high vacuum regimes present in modern field emission (S)TEMs and the fact that these new detectors may be operated at relatively high temperatures, windowless SDD detectors can also now be easily and safely incorporated as part of the analytical system. 
By dispensing with the window and support grid, not only is true solid angle increased, but the increased collection efficiency dramatically improves signal to noise ratios for lower energy X-rays such as NKα and OKα. Figure1 shows this huge improvement in light element performance for the new Oxford Instruments X-MaxN 100TLE detector compared to an 80mm2 SDD detector.
This improved performance is especially useful for TEM instruments without EELS capabilities, where light element analysis can be problematic. Sensitivity can be further increased by mounting more than one detector on suitable instruments – this can maximize solid angle up to 2.0sr.  Not only is collection efficiency for un-tilted samples increased, but it also gives the ability to tilt the sample in a negative direction which can be useful when collecting simultaneous diffraction data.
Figure 2 shows an example of elemental distribution in an LED nanowire taken with a single X-Max 100 TLE in 10 minutes and structures as small as 2nm can be identified. The Al and Ga distribution is well defined and N concentrations are evenly distributed throughout the wire. In addition a thin out coating of oxide only a few nanometers thick is clearly seen in the mapping data. High resolution maps such as these can be collected in just 5 -10 minutes using a single detector.

Conclusions:
Large solid angle windowless detectors extend the capabilities of EDS analysis to application areas which were formerly regarded as being limited to EELS only. Higher collection efficiencies, especially at low energy now makes EDS a truly viable option for light element analysis on the nanoscale.


Fig. 1: Comparison of X-Max 80T (yellow - 80mm2 window) and Maxᶰ 100TLE (red - 100mm2 very large solid angle windowless) SDDs. Spectrum has been normalized on the Ni Kα peak to show the sensitivity enhancement of the windowless operation for low energy X-rays.

Fig. 2: Mapping of LED Al-Ga-N nanowires showing the elemental distribution throughout the wire. Each map has a resolution of 266x113 pixels.

Type of presentation: Poster

IT-5-P-2281 Spatially resolved EELS with an in-column Omega filter - characterization of energy filter aberrations and their correction by image processing

Entrup M.1, Kohl H.1
1Physikalisches Institut und Interdisziplinäres Centrum für Elektronenmikroskopie und Mikroanalyse (ICEM), WWU Münster, Wilhelm-Klemm-Str. 10, 48149 Münster, Germany
michael.entrup@wwu.de

Spatially resolved EELS (SR-EELS) [1] is a technique to preserve spatial information when recording EEL spectra. Essentially, many EEL spectra are recorded in parallel as a function of one spatial coordinate, perpendicular to the energy dispersive direction. This method is useful for investigating specimens like interfaces and layer systems. We apply SR-EELS in a TEM with an in-column Omega filter [2]. Remaining aberrations can be corrected by processing the recorded SR-EELS dataset, using the results of a previous characterisation measurement.
The characterization measurement is performed using the small filter entrance aperture - 100µm instead of 500µm used for the final SR-EELS measurement. The aperture is shifted along the lateral axis. At several positions a SR-EELS dataset is recorded. For each energy channel we can extract the position of the aperture borders (yb). From this information we can calculate the width (w) and the position (y) of the aperture. To increase the signal to noise ratio, up to 64 energy channels are binned.
One aberration is directly visible when inspecting a SR-EELS datasets. The width of the aperture decreases with increasing energy loss. Figure 1a) shows a superposition of 3 datasets recorded using the described method. The borders of the apertures are plotted in figure 2. In addition to the change of the aperture width, the borders are curved. A two dimensional polynomial of 2nd order (Σij Aij ΔEi yb(ΔE=0)j) is used to describe this aberration, where yb(ΔE=0) is the position of the border at E=200keV, the energy of electrons that are not deflected by the energy filter. The correction of the aberration is done by image processing. Figure 1b) shows the correction of figure 1a) using the polynomial plotted in figure 2.
The change of the aperture width is best visible when plotting the width of the aperture as a function of the position, for only one energy channel (see figure 3). With increasing distance to the image centre, the width of the aperture decreases. A polynomial of 2nd order describes this well. This aberration is depended on the excitation ΔQSinK7 of the 7th corrector of the energy filter. For ΔQSinK7=-32% there is nearly no change in width of the aperture. Figure 4 shows the variation of the aperture width for all recorded energy channels. A two dimensional polynomial of 2nd order (Σij Aij ΔEi yj) is used for fitting. The dependency w(ΔE) is clearly visible in both graphs, while only ΔQSinK7=0% shows the dependency w(y).

[1] L. Reimer et al., Ultramicroscopy 24 (1988) 339-354.
[2] S. Lanio, PhD thesis (1986), TH Darmstadt.
[3] The code that has been used to perform the characterisation is available on GitHub: https://github.com/EFTEMj/EFTEMj/Scripts+Macros


Fig. 1: a) A superposition of 3 SR-EELS datasets recorded while the excitation of the 7th corrector was changed by ΔQSinK7=-32%. A amorphous carbon film has been used to guaranty a uniform signal which simplifies the processing. b) A corrected version of a).

Fig. 2: The aperture borders extracted from the datasets shown in figure 1a). Only every 2nd data point is displayed. A polynomial of 2nd order can be used to fit each border separately. Introducing the position of the polynomial at ΔE=0eV, a single two dimensional polynomial of 2nd order can be used to fit all borders simultaneously.

Fig. 3: The width of the filter entrance aperture is plotted as a function of the position on the lateral axis (only the energy channel ΔE=0eV is considered). The graphs differ by the excitation ΔQSinK7 of the 7th corrector. A second order polynomial has been used for fitting.

Fig. 4: The width of the filter entrance aperture is plotted as a function of the position on the lateral axis. In contrast to figure 3 all recorded energy channels are considered. The optimal excitation (ΔQSinK7=-32%) of the 7th corrector is compared to the default excitation (ΔQSinK7=0%). Only every 4th data point is displayed.

Type of presentation: Poster

IT-5-P-2311 Characterising severe plastic deformation: a quantitative assessment of TEM imaging using transmission Kikuchi diffraction in the SEM

Trimby P. W.1, Xia J. H.2, Sha G.2, Schmidt N. H.3, Sitzman S.3, Tort M.4, Xia K.4, Ringer S. P.1,2
1The Australian Centre for Microscopy & Microanalysis, The University of Sydney, NSW 2006, Australia , 2School of Aerospace, Mechanical & Mechatronic Engineering, ARC Centre of Excellence for Design in Light Metals, The University of Sydney, NSW 2006, Australia, 3Oxford Instruments Nanoanalysis, Halifax Road, High Wycombe, Bucks, HP12 3SE, United Kingdom , 4Department of Mechanical Engineering, The University of Melbourne, VIC 3010, Australia
patrick.trimby@sydney.edu.au

The enhanced physical properties attributed to nanocrystalline materials have resulted in significant recent research focus on grain size refinement using severe plastic deformation (SPD). SPD-processed materials typically have ultrafine or nanocrystalline grain sizes, often with very high dislocation densities. These attributes make them difficult to analyse using conventional scanning electron microscope (SEM) based techniques such as electron backscatter diffraction (EBSD), resulting in most researchers relying on the higher resolution capabilities of the transmission electron microscope (TEM).
The recent emergence of transmission Kikuchi diffraction (TKD) in the SEM enables routine characterisation of materials with mean grain sizes below 50 nm, and with high intragranular dislocation densities [1, 2]. However, TKD analyses of samples previously characterised using TEM have sometimes produced discrepancies in the final grain size estimates.
This is the first in-depth direct correlation between TEM imaging and TKD in the SEM: we used a variety of Al-alloys that have been deformed using SPD, firstly imaging using 200 kV TEM (JEOL 2100) and then analysing the same areas using TKD in the SEM (Carl Zeiss Ultra Plus with Oxford Instruments AZtec EBSD).
The results reveal some startling differences. In an Al-6060 alloy deformed by high pressure torsion (5 revolutions at 180 °C, under 6 GPa), the recovery of dislocations enables relatively clear TEM imaging of the grain structure, as shown in fig. 1. TKD mapping (fig. 2) confirms the location of the high angle boundaries and shows that the intragranular dislocations and precipitates visible in the brightfield image are associated with very low misorientations (<1°). However, a misorientation profile across one grain shows that the cumulative lattice distortion can be significantly higher, in this case nearly 4° (fig. 3).
In Al-Cu-Mg alloys that have been deformed by equal channel angular processing (ECAP) at room temperature, the correlation between TEM and TKD is more challenging. The increased dislocation density makes clear TEM imaging difficult, and TKD results indicate that many grain-like features imaged in the TEM are in fact part of larger grains with significant intragranular subgrain structure. We will present numerous correlative analyses between the two techniques; the results have important implications for the characterisation of SPD materials, and show the benefit of the TKD technique for the rigorous measurement of microstructural properties.

References:
[1] R.R. Keller and R.H. Geiss, J. Microscopy, 245 (2012), p. 245-251.
[2] P. W. Trimby et. al., Acta Materialia, 62 (2014), p. 69-80.


Fig. 1: Bright field TEM image of an HPT deformed Al-6060 alloy.

Fig. 2: TKD orientation map (IPF colouring) of the same area (measurement step size 8 nm), with high angle boundaries (>10°) in black, low angle boundaries (1-10°) in red.

Fig. 3: Detailed quantitative analysis of a single grain in Figs 1 & 2. Top left: bright field TEM image. Bottom left: TKD map showing the change in lattice orientation relative to the central spot. Right: graph showing the change in orientation across the 700 nm transect A-B, relative to point A.

Type of presentation: Poster

IT-5-P-2335 WDX-measurement of Ta-, W- and Re-concentration profiles in a Nickel/Superalloy diffusion couple using Lβ-X-ray-lines

Nissen J.1, Berger D.1, Epishin A.2, Link T.2
1Technical University Berlin, Center for Electron Microscopy (ZELMI), Straße des 17. Juni 135, 10623 Berlin, Germany, 2Technical University Berlin, Institute of Material Science, Ernst-Reuter- Platz 1, 10587 Berlin, Germany
joerg.nissen@tu-berlin.de

Ni-base superalloys are multicomponent alloys used at temperatures up to about 1100°C. At such high temperatures, diffusion plays the principal role for structural stability and mechanical behaviour. The material under investigation is CMSX-10, which consists of 11 elements (Al, Ti, Co, Cr, Ni-base, Nb, Mo, Hf, Ta, W, Re). CMSX-10 is diffusion welded with pure Ni under vacuum at 1050°C, 10 MPa, 1 h, then annealed at 1050°C for 128 days. In order to quantify the diffusion kinetics in such a multicomponent system, the diffusion profiles in Ni/CMSX-10 diffusion couples have to be measured. However, for the key strengthening elements Ta, W and Re this task is not trivial because they are neighbours in the periodic system (atomic numbers 73, 74, 75) and their concentrations are quite small, 1-3 at%. Therefore, X-ray peaks of these elements are small and they overlap. For these reasons, an optimised method is presented.
Measurement of the diffusion profile by EDX-microanalysis in a SEM is not quite reliable because the energy resolution of the detector is too large (127 eV @ 5.9 keV), as can be seen from the overlapping of the M-lines of Ta, W and Re in Figure 1 and the L-Lines in Figure 2, respectively. Thus, WDX-analysis with high energy resolution becomes essential, in our case with the Field Emission Gun Electron Probe Microanalyser (FEG-EPMA) JEOL JXA-8530F, having a resolution of about 15 eV @ 5.9 keV (LIF). However, figure 1 shows, that even now the Ta-Mβ- and W-Mα-lines cannot be separated (ΔE=9 eV) as well as the W-Mβ- and Re-Mα-lines (ΔE= 8 eV).
Energies of the L-lines are in general about 5 times higher than those of the M-lines, thus also the separation of the lines. In Figure 2 it can be seen, that the Lα-peaks of W and Re are isolated, however, the very close and strong Ni-Kβ-line falsifies their background. Therefore, the Lβ-lines are used. The energy differences between TaLβ2 and WLβ1 (ΔE=20 eV) as well as WLβ2 and ReLβ1 (ΔE=50 eV) are large enough to allow a reliable peak deconvolution. To excite the L-lines of Ta, W and Re (E≈10 keV), a 20 kV accelerating voltage is applied. Anyhow, the small Ta-, W- and Re-concentrations give only small peak/background ratios, making necessary a careful background subtraction. The method was checked by measuring the element concentrations in CMSX-10, which gives results very close to the nominal composition.
The diffusion profiles were measured 1.8 mm across the interface with a step size of 5 µm. Figure 3 shows the profile scan for Ta-, W-, Re-, Ni and Al. Comparison of the experimental concentration profiles with such modelled by the software DICTRA shows a very close match. Therefore, it is proved that Lβ-lines might be used for the quantitative element analysis in WDX.


Fig. 1: M-lines of Ta, W and Re in an EDX/WDX-spectrum of CMSX-10

Fig. 2: L-lines of Ta, W and Re in an EDX/WDX-spectrum of CMSX-10

Fig. 3: WDX-Profile scan of Ta, W, Re, Ni and Al in CMSX-10

Type of presentation: Poster

IT-5-P-2350 Momentum-resolved electron energy-loss spectroscopy of MoS2 and graphene heterostructures

Mohn M.1, Hambach R.1, Wachsmuth P.1, Benner G.2, Kaiser U.1
1Electron Microscopy Group of Materials Science, Ulm University, Ulm, Germany, 2Carl Zeiss Microscopy GmbH, Oberkochen, Germany
michael.mohn@uni-ulm.de

By comparison of momentum-resolved electron energy-loss spectra and ab-initio calculations we analyze high-energy plasmons in 2D heterostructures made of graphene and few- or monolayer MoS2.
We are particularly interested in MoS2 monolayers covered by graphene (G/MoS2/G sandwiches) as it has been shown that such a configuration protects the MoS2 from beam damage [1].

Our experiments have been performed using a low-voltage (20–80 kV) transmission electron microscope with simultaneous acquisition of spectra for different momentum transfers. We have recorded energy-loss spectra in the range of 0–50 eV for momentum transfers along certain crystallographic axes within the Brillouin zone. For very high momentum and energy resolution, we have used a monochromated Zeiss Libra 200 based TEM in diffraction mode with an in-column Ω energy filter (SALVE I [2,3]).

The corresponding ab-initio calculations have been performed as follows: For the ground-state simulations, we have used the density-functional theory (DFT) software ABINIT [4] with pseudopotentials and local-density approximation (LDA). Energy-loss spectra have been calculated with the dp-code [5] within the random-phase approximation (RPA).

Eventually, deficiencies in both the experimental data and the simulations can be spotted by comparing the ab-initio calculations to the corresponding electron energy-loss spectra. These deficiencies include consequences of the approximations we made in the ab-initio calculations. Besides, our measurements are subject to the following experimental difficulties: First, despite the use of low acceleration voltages, the beam-sensitivity of the MoS2 monolayers limits the acquisition times of the spectra. During an exposure of only a few minutes, beam damage and contamination may lead to significant changes in the spectra. Second, the signal decays drastically with increasing momentum transfer, so that the background noise of the CCD plays a crucial role.

[1] G. Algara-Siller et al., Appl. Phys. Lett. 103, 203107 (2013)
[2] U. Kaiser et al., Ultramicroscopy 111, 1239-1246 (2011)
[3] P. Wachsmuth et al., Phys. Rev. B 88, 075433 (2013)
[4] X. Gonze et al., Comp. Mat. Sci. 25, 478 (2002)
[5] V. Olevano, L. Reining, F. Sottile, http://www.dp-code.org (1998)


Type of presentation: Poster

IT-5-P-2395 Toward X-ray Quantitative Microanalysis Maps with an Annular Silicon Drift Detector

Demers H.1, Brodusch N.1, Woo P.2, Gauvin R.1
1Department of Mining and Materials Engineering, McGill University, Montreal, Quebec, Canada., 2Hitachi High-Technologies Canada Inc., Toronto, Canada.
hendrix.demers@mail.mcgill.ca

The scanning electron microscope (SEM) was primary developed for imaging applications. With the introduction of the Si(Li) energy dispersive spectrometer (EDS), simultaneous imaging and x-ray microanalysis became possible. However, long working distance and high current were needed because the position and small solid angle of the EDS detector. SEM was initially and is still optimized for imaging applications, where the high spatial resolution is generally obtained at short working distance. This problem is still relevant today and unfortunately x-ray microanalysis is never performed in the best imaging conditions, i.e., not with the smallest probe size. The annular silicon drift detector (SDD) system is inserted below the objective lens and has four segments which give a higher solid angle (up to 1.2 sr). Also, a lower working distance and probe current can be used. An improved spatial resolution becomes possible during x-ray microanalysis. However, the effect of the detector geometry and position on the quantification microanalysis is unknown.

Because of the position of the detector, Mylar windows are used to prevent the backscattered electrons (BSEs) to damage the SDD segments. Three window thicknesses are available for this detector and their effect on the x-ray spectra is shown in Figure 1. The shape of the background was strongly affected by the window absorption at low x-ray energy. For accurate quantitative analysis, the calculation of peak net intensity depends on the background subtraction method used. Different approaches are currently studied with this annular SDD. Another artefact created by the window is the generation of C and O peaks and bremsstrahlung x-rays in the window by the BSEs. Figure 2 shows the variation of the output count rate for the spectrum and the Cu Lα peak with the working distance for the three window thicknesses. An optimum working distance was observed for the Cu Lα peak as predicted by the calculation of the solid angle of this detector. However, no decrease of the output count rate was observed. The x-ray emission in the window negates the effect of the solid angle. This effect is more pronounce at high accelerating voltage. An example of x-ray elemental maps of a mineral ore sample acquired with annular silicon drift detector (SDD) at low accelerating voltage is shown in Figure 3.

The effect of this detector geometry and position on the correction model is currently studied to obtain quantitative maps from the elemental maps. With adapted correction model, the annular SDD with its larger solid angle will clearly revolution the quantification microanalysis by moving from point analysis to quantitative micrograph with simultaneous electron imaging.


Fig. 1: Copper sample spectra with three different window thicknesses for annular silicon drift detector (SDD). The gray line shows the expected Duane-Hunt limit at 5 keV. The C Lα peak intensity decrease by 3.2 times with the 3 µm-thick window and 30 times with the 7 µm-thick window.

Fig. 2: Variation of the experimental output count rate with working distance for three different window thicknesses for annular silicon drift detector (SDD). The gray line shows the detector bottom position at 7.5 mm.

Fig. 3: X-ray elemental maps of a mineral sample acquired with annular silicon drift detector (SDD) at low accelerating voltage of 4 kV and working distance of 11 mm. The image size was 11 x 9 µm2 (1280 x 960 pixels) with an acquisition time of 1162 s (input count rate of 117.6 kcps).

Type of presentation: Poster

IT-5-P-2428 Obtaining an accurate quantification of light elements by EDX: K-factors vs. Zeta-factors

Lopez-Haro M.1,3, Bayle-Guillemaud P.1, Mollard N.1, Saint-Antonin F.2, van Vilsteren C.3, Freitag B.3, Robin E.1
1CEA, INAC/UJF-Grenoble 1, UMR-E, SP2M, LEMMA, Minatec, 38054 Grenoble Cedex 9, France, 2CEA, LITE-DTNM, L2N, 38054 Grenoble Cedex 9, France, 3FEI Company, P.O. Box 80066, KA 5600 Eindhoven, The Netherlands.
miguel.lopez-haro@cea.fr

The new energy dispersive X-ray (EDX) technology based on four silicon drift detectors (SDD) with a windowless design provides new possibilities in the field of analytical characterization at nanometer scale. The four detectors are symmetrically arranged with respect to the sample and this unique configuration provides very high collection efficiency, allowing high counting statistics and rapid acquisition of X-ray spectra, line scans and maps. However, new methodologies for the precise quantitative assessment of the elemental composition at nanometer scale are still needed.
Classically, EDX quantification has been carried out using “Cliff-Lorimer” ratio method. This method requires the knowledge of the k-factors and their precise determination is a key point to obtain an accurate quantification. They can be determined theoretically or experimentally, nevertheless, several limitations are found: i) the theoretical k-factors present large uncertainties, ii) the experimental determination of k-factors required multi-element samples with known compositions and iii) X-ray absorption correction may be important for low energy X-ray emissions, especially for light elements, which require the prior knowledge of the specimen mass thickness. To overcome such limitations, a new procedure named “zeta (ζ)-factor” method has been proposed [1]. In this method, the composition and mass thickness are computed simultaneously for each analysis point enabling X-ray absorption correction.
In this work, we present an accurate EDX quantification of various samples containing light elements or elements with low energy X-ray lines using the ζ-factor method. In this regard, a Super-X Tecnai-OSIRIS installed at PFNC-CEA-Grenoble and operating at 200kV has been used. Fig. 1 shows a representative HAADF image of a thin foil of wollastonite (CaSiO3) prepared by FIB, together with the EDX maps of Ca (3.69 keV), Si (1.74 keV) and O (0.53 keV). Individual profiles of the net X-ray counts are extracted from a line scan (see arrow in Fig. 1a) and quantified using k- and ζ- factors (Fig. 2). Quantification using the k-factors gives wrong results as a consequence of the strong absorption of oxygen (Fig. 2b). Conversely, an excellent agreement between the computed and expected results is obtained using the ζ-factor method (Fig. 2d), due to the specimen thickness is determined for each analyzed point (Fig. 2c) and allowing therefore the X-ray absorption correction.
This example clearly illustrates the potential of the ζ-factor method using the new Super-X detector. By measuring composition and mass thickness simultaneously, the ζ -factor method is a very promising tool for quantitative 3D reconstructions.

[1] M Watanabe and DB Williams, Journal of Microscopy 221 (2006) p. 89


Fig. 1: HAADF-STEM image of wollastinite recorded on a Tecnai-OSIRIS (a) and EDX elemental maps of Ca (b), Si (c) and O (d) using the Super-X detector.

Fig. 2: EDX Line scan extracted from the site marked with an arrow (a). EDX quantification using the k-factor method (b). Thickness measurement obtained by EDX analysis (c) and EDX quantification in atomic percentage using the ζ (zeta)-factor (d).

Type of presentation: Poster

IT-5-P-2448 A simple EDXperformance test for transmission electron microscopy

Van Cappellen E.1, Porcu M.2, Delille D.2, Sudfeld D.2
1FEI Company, Hillsboro, USA, 2FEI Company, Acht, The Netherlands
eric.van.cappellen@fei.com

In the last 5 years solid angles increased dramatically (a factor 10 to about 1srad) and some systems are windowless further improving collection efficiency. In practice this means that for most elements a few percent (up to 10% for heavy elements) of the ionization events are now detected. Besides speeding-up conventional X-ray analysis (point analysis and 2D EDX elemental maps) large solid angle detectors have also enabled new EDX applications such as atomic resolution elemental mapping and 3D EDX tomography but this last application only on condition that X-ray collection is possible over a large sample tilt range (like for FEI’s Super-X™ detector).
Although the geometrical definition of a solid-angle is straightforward it is tedious to experimentally verify specified numbers and not all solid-angles yield the same detection efficiency. Some areas of the available real estate around the sample are better than others for X-ray detection and this leads to the concept of “quality of solid-angle” (see fig. 1). Here we propose to determine the quality of solid angle by measuring the output X-ray count-rate per nA of primary electron beam current on a very well-defined sample. This calibration sample needs to have an undisputable and stable composition over time, as well as a fixed and known thickness and must be easily and reliably produced in large quantities to allow for comparisons between systems.
In this study we propose to use 200nm thick Si3N4 windows made in 200μm thick silicon wafers and cut into 3mm discs to fit in regular low-background TEM holders (see fig. 2). Wafer processing technology ensures very good thickness uniformity and thickness reproducibility. Furthermore Si3N4 is stoichiometric and stable certainly when the membrane is 200nm thick and the electron beam is defocussed. Last but not least each wafer yields over 300 TEM samples which keeps the price down and guarantees easy access and supply. Actual measurements will be discussed.


Fig. 1: A single large solid angle EDX: The average take-off angle is NOT θ as the lower part of the detector (dark area) doesn’t contribute at all to the signal. Below the red line no X-rays are counted and above the signal will gradually increase as the take-off angle increases (blue curve).

Fig. 2: The proposed sample: a Si3N4 window in a 3mm disc of 200μm thick silicon wafer, completely flat on the top to avoid shadowing.

Type of presentation: Poster

IT-5-P-2506 Dynamic, Analytical EDX studies of B/Ni composite nanowires with MEMS heating holder

Sudfeld D.1, Lourie O.1, Mele L.1, Dona P.1, Konings S.1, Delille D.1, Van Cappellen E.2, Barton B.1, Jinschek J. R.1, Freitag B.1
1FEI Electron Optics B. V., 5651 GG Eindhoven, The Netherlands, 2FEI Company, 5350 NE Dawson Creek Drive, Hillsboro, OR 97124, USA
Daniela.Sudfeld@fei.com

EDX measurements at high temperatures were a challenge for many years due to the technology constrains of the past related to the limitations of commercial Si(Li) EDX (energy-dispersive X-ray spectroscopy) detectors. With a newly designed MEMS holder and state-of-the-art SDD technology nowadays analytical EDX studies at elevated temperatures can be performed on a regular basis in a Scanning Transmission Electron Microscope (S/TEM), FEI’s TalosTM with ChemiSTEMTM Technology [1].
Here we present fast chemical maps of B/Ni composite nanowires on nanometer scale by EDX which were done with the new VeloxTM software, see Figure 1. Successful synthesis of crystalline nanowires composed of the refractory light materials such as Boron can enable novel applications for nanoelectronics [2-5]. Boron/Nickel composite nanostructures were prepared by a CVD-based synthetic procedure with a Ni-based compound catalyst; naturally blended with high conductivity and refraction index. The properties of this binary nanomaterial at room temperature, Figure 2, are compared to those achieved from heating experiments with temperatures up to 1000 °C. 2D-3D EDX chemical mappings show clearly the core-shell structure of the wires: B in the shell and Ni in the core. This is amplified at elevated temperatures of ca. 500 °C, see Figure 3. At ca. 1,000 °C EDX maps reveal also that Ni vanishes from the core, leaving behind hollow B nanowire (nanotube) structures.
For the given in situ experiments FEI’s NanoExTM heating holder was used with a small, consumable semiconductor (MEMS) device as the heater and providing a direct read-out of the temperature value at all times during the dynamic experiment with a known and reproducible temperature distribution over the heated area. The NanoEx solution is optimized for ChemiSTEM EDX experiments to trace compositional changes correlated with temperature and electrical stimuli. The holder geometry is suitable for high tilt angles, also allowing its use for 3D experiments.

[1] P. Schlossmacher et al., Microscopy Today 18(4) (2010) 14.
[2] CJ Otten, et al., J Am Chem Soc. 2002 May 1;124(17):4564.
[3] D. Wang et al., APL 2003, 183(25):5280.
[4] W. Ding et al., Mech, Comp. Sci. and Techn. 2006, 66:1109.
[5] J. Tian, et al., “Boron nanowires for flexible electronics”,APL 2008 93:122105-7-5.


Fig. 1: EDX compositional maps of Boron nanowires at room temperature. The total acquisition time is ca. 10 minutes for the map sized 256 x 256 Px with a speed of 9.1 ms/Px and ~50cs/Px signal for the B-K edge.

Fig. 2: Dynamic EDX compositional analysis at the begin of the experiment at room temperature.

Fig. 3: Dynamic EDX compositional analysis comparing the same area at the heating temperature of 500 °C showing the Ni particles before they disappear when the heating temperature gets further raised and the specimen finally got cooled down.

Type of presentation: Poster

IT-5-P-2551 Quantitative X-Ray Microanalysis in the Scanning and Transmission Electron Microscopes with the Generalized f-Ratio Method

Demers H.1, Brodusch N.1, Trudeau M.2, Gauvin R.1
1Department of Mining and Materials Engineering, McGill University, Montreal, Quebec, Canada, 2Materials Science, Hydro-Québec Research Institute, Varennes, Québec, Canada
hendrix.demers@mail.mcgill.ca

Quantitative x-ray microanalysis of bulk samples is usually obtained by measuring the characteristic x-ray intensities of each element in a sample and in a corresponding standard. The k-ratio of the measured intensities from the unknown material over the standard is related to the concentration using the ZAF or φ(ρz) correction methods. Under optimal conditions, accuracies approaching 1% are possible. However, all the experimental conditions must remain the identical during the sample and standard measurements. This is not possible with a cold-field emission scanning electron microscope (CFE-SEM) where beam current can fluctuate by 5% in its stable regime. To address this issue, a new method was developed using a single spectrum measurement (Horny et al., 2010; Gauvin, 2012). It is similar in approach to the Cliff and Lorimer (1975) ratio method developed for the analytical transmission electron microscope. However, corrections are made for x-rays generated from thick specimens using the ratio of the characteristic x-ray intensities of two elements in the same material. The proposed method utilizes the ratio of the intensity of a characteristic x-ray normalized by the sum of x-ray intensities of all the elements measured for the sample. Uncertainties in the physical parameters of x-ray generation are corrected using a calibration factor that must be previously acquired or calculated. With this method, relative accuracies better than 5% were obtained in a CFE- SEM.

The f-ratio method was generalized to more than two elements. The correction factors are still acquired experimentally relatively to two elements and they do not change with composition. They are obtained from measurement of one known phase. The concentration curves versus f-ratio are obtained by Monte Carlo simulations and the unknown concentrations are calculated from these curves and the measured f-ratios by combination of multi-dimensional interpolation and iterative procedure. An example is shown in Figure 1, where quantitative x-ray maps of ternary Al-Mg-Zn diffusion couple sample were obtained with a CFE- SEM at 5 kV. The generalized f-ratio method was also applied to a thin specimen in the transmission electron microscope. An example is presented in Figure 2, trace element concentration of Fe in a Zr-Nb alloy was determined. The generalized f-ratio method allows the quantification of multi-elements sample in both SEM and TEM. Furthermore, in which condition this method can be applied to heterogeneous sample is currently explored.

References:

P. Horny, E. Lifshin, H. Campbell and R. Gauvin, Microscopy and Microanalysis, 16, 821-830 (2010).

R. Gauvin, Microscopy and Microanalysis, 18, 915-940 (2012).

G. Cliff and G. W. Lorimer, Journal of Micrsocopy, 103, 203-207 (1975).


Fig. 1: Quantitative x-ray maps of ternary Al-Mg-Zn sample obtained with the f-ratio method at 5 kV. The weight fractions of each element (A: aluminum, B: magnesium, C: Zinc) are represented by a gray scale: black 0% and white 100%. D A ternary phase diagram was obtained from the x-ray maps.

Fig. 2: X-ray quantification of Zr-Nb-Fe alloy obtained with the f-ratio method in a TEM at 200 kV. A Correction factors measured with 76 wt% Zr, 19 wt% Nb and 5 wt% Fe alloy. B Iron concentration in a Zr-Nb phase was calculated with the f-ratio method and Monte Carlo simulations.

Type of presentation: Poster

IT-5-P-2570 Orbital ordering of A-site ordered SmBaMn2O6 studied by inelastic scattering accampanied by Mn-L shell excitation

Saitoh K.1, Toake Y.2, Tanaka N.1, Takenaka K.3
1EcoTopia Science Institute, Nagoya University, 2Department of Crystalline Materials Science, Nagoya University, 3Department of Applied Physics, Nagoya University
saitoh@esi.nagoya-u.ac.jp

Manganite perovskites have been drawing a grate attention from the unique properties such as metal-insulator transition, colossal magneto-resistance, etc. Such unique properties are attributed to a charge and orbital ordering (COO) of the 3d electrons in the eg orbitals of Manganese by resonant X-ray scattering experiments. SmBaMn2O6 shows the A-site ordering of Sm and Ba at room temperature. The crystal structure has a 2√2ap × 2√2ap ×4ap supercell with ap the fundamental cubic perovskite unit cell reported. In the present study, we determine the orbital ordering of SmBaMn2O6 by the convergent-beam electron diffraction (CBED) and inelastic scattering [3] accompanied by Mn-L shell excitation.

Samples of SmBaMn2O6 were synthesized by a solid-state reaction using Sm2O3, BaCO3, and MnO2. CBED patterns were taken from an area of about 10 nm in diameter. Inelastic scattering patterns accompanied by the Mn-L shell excitation were taken using an energy-filtering system fitted to the bottom of the electron microscope. A series of inelastic scattering patterns at successive energy losses from 620 eV to 670 eV with an energy step of 1-2 eV were taken with an energy window of 1-2 eV.

Figures 1(a), 1(b) and 1(c) show CBED patterns of SmBaMn2O6 taken at incidences in the [0 0 1], [0 1 0] and [0 4 1] orientations, respectively. The [0 0 1] and [0 1 0] patterns show two types of mirror symmetries and twofold rotation symmetry. Thus, the point group is determined to be mmm. The patterns does not show any systematic extinction rules of reflections, indicating that the lattice type is primitive P. From the dynamical extinction lines in the h00 (h = odd) reflections in the [0 0 1] pattern and 0 -4 17 reflection in the [0 4 1] pattern, the space group of SmBaMn2O6 was determined to be Pnam.

Figure 2(a) shows an inelastic scattering pattern of SmBaMn2O6 accompanied by the Mn-L shell excitation taken at an incidence in the [0 1 0] orientation. The pattern clearly shows an elongation along the a* axis. Figures 2(b) and 2(c) show inelastic scattering patterns simulated from two kinds of the orbital ordering composed of the 3z2-r2 type orbitals [1] and the x2-y2 type orbitals [2], respectively. The anisotropic feature of the experimental inelastic scattering patterns agrees well with that of the x2-y2 type model. An orbital-ordering model was constructed from the CBED symmetry and a qualitative comparison between the experimental and simulated inelastic scattering anisotropy.

References

[1] M. Uchida et al., J. Phys. Soc. Jpn. 71, (2002) 2605.

[2] M. García-Fernández, et al, Phys. Rev. B 77, 060402(R) (2008).


[3] 
K. Saitoh et al., J. Electron Microsc. 55 (2006) 281.; K.Saitoh et al., J. Appl. Phys. 112 (2012) 113920.


The present work was partly supported by the Grant-in-Aid for Challenging Exploratory Research (No. 23654117), the Ministry of Education, Culture, Sports, Science and Technology, Japan.

Fig. 1: Convergent-beam electron diffraction patterns of SmBaMn2O6 taken at incidences in the [0 0 1], [0 1 0] and orientations. The and patterns show two types of mirror symmetries and twofold rotation symmetry. Dynamical extinction lines in the [0 0 1] and [0 -4 17] patterns indicate that there exist a and n glide planes.

Fig. 2: Experimental inelastic scattering pattern of SmBaMn2O6 accompanied by Mn-L shell excitation taken at an incidence in the [010] orientation (a) and simulated patterns from the 3z2-r2 orbital model (b) and the x2-y2 orbital model (c). The elongation feature in the experimental pattern agrees well with the pattern simulated from the x2-y2 model.

Type of presentation: Poster

IT-5-P-2665 Electron energy-loss mapping of a perovskite-based solar cell

Divitini G.1, Peng X.1, de la Peña F.1, Saliba M.2, Snaith H. J.2, Ducati C.1
1University of Cambridge, 2University of Oxford
gd322@cam.ac.uk

The investigation of nanomaterials for solar cells, aimed at determining chemical structure and morphology, is a vital step in the development of novel materials to face the current energy crisis. Thin film solar cells have been evolving in the last decades, attaining interesting performance levels. In particular, highly efficient perovskite-based solar cells were demonstrated last year [Liu2013], and there is booming interest in the design of such systems. The addition of metal-oxide core-shell nanoparticles is also being investigated in several thin film devices for increasing light harvesting.
Morphology, both at the nano- and the micro-scale, plays a major role in several processes which affect the behaviour of thin film solar cells. Focused Ion Beam milling (FIB) processing is a powerful tool to extract samples from a full solar cell device for analysis with a transmission electron microscope (TEM). Characterisation in a TEM can shed light on the local elemental composition of the device, as well as spatially-resolved information on the electronic energy levels.
Recent developments in TEM analytical capabilities have enhanced the possibilities for investigation further: brighter electron guns and new detectors for energy-dispersed x-ray spectroscopy (EDX) and electron energy-loss spectroscopy (EELS) now allow large spectrum images to be acquired quickly. As a consequence, devices with organic components can be examined with a reduced electron irradiation, limiting beam damage artifacts.
In this work we apply FIB preparation to a perovskite-based solar cell and investigate the morphology of the photoanode using TEM. Elemental and EELS maps are acquired, and the effect of plasmonic nanoparticles is investigated. Information on the spatial distribution of metal nanoparticles is extracted to provide feedback on the device fabrication process.
Spectrum images are analysed using principal component analysis and blind source separation to optimise signal-to-noise ratio, thus obtaining high quality maps while limiting the electron dose on the specimen. This procedure also naturally separates different compounds without introducing operator bias, and is particularly effective in the presence of complex compounds, such as the perovskite active layer.

[Liu2013] M. Liu, M. B. Johnston and H. J. Snaith, Nature 501, 395–398, 2013


GD, XP, FDLP and CD thank ERC for funding.

Fig. 1: Cross-sectional view of a perovskite-based solar cell.

Type of presentation: Poster

IT-5-P-2705 Simultaneous panchromatic and color live cathodoluminescence imaging

Kološová J.1, Jiruše J.1
1TESCAN Brno, s.r.o., Brno, Czech Republic
jolana.kolosova@tescan.cz

Cathodoluminescence (CL) imaging is a standard non-destructive analytical technique. It provides information about composition and crystal structure of the studied material. Scanning electron microscope (SEM) equipped with a CL detection system allows panchromatic, monochromatic or color CL imaging, often in combination with other techniques (EDX, WDX, EBIC…). We can see growing needs for a seamless integration of multiple detection devices into one multi-analytical SEM system. In line with this trend, we have developed a new versatile “two in one” CL detector capable of simultaneous panchromatic and color live imaging.

Figure 1 illustrates the advantage of such simultaneous imaging. It shows a single scan image of a rhyolite sample. In the color image, two types of grains are clearly distinguishable – blue quartz and pinkish topaz. On the other hand, in the panchromatic image the zoning of quartz and topaz grains is more distinct, as the panchromatic channel of the detector collects the signal for the whole spectral range and maintains higher signal to noise ratio.

Versatility of the detector lies in its unique collection optics. Both sub-micrometer high resolution images (see Figure 2) and images with extra large field of view (FOV) can be done. FOV up to 35 mm can be achieved with a single scan, no stage scanning is needed. The collection efficiency of the detector is uniform over the full FOV even when topographic samples are imaged.

The new compact CL detector can be easily integrated into a modular multi-analytical SEM-based system. Simultaneous acquisition of CL together with other signals is straightforward, see Figure 3 for simultaneous CL and BSE. There is no need for complex sample preparation or precise setting of working distance and thus high quality and fast CL analyzes of various samples can be done easily.


The authors acknowledge funding from the European Union Seventh Framework Program [FP7/2007-2013] under grant agreement No. 280566, project UnivSEM.

Fig. 1: Color (left) and panchromatic (right) image of a rhyolite sample with quartz (blue) and topaz grains. Zoning typical for volcanic quartz is more contrasting in the panchromatic image.

Fig. 2: Sub-micrometer high resolution color image of GaN wires covered with InGaN quantum wells (view from above). Quantum wells (bluish) are deposited close to the sidewall surfaces. Sample courtesy MPI for the Science of Light, Erlangen.

Fig. 3: Color CL (left) and BSE (right) images acquired simultaneously on a ruby containing rock with baddeleyite (blue) grains.

Type of presentation: Poster

IT-5-P-2719 Development of an Analytical TEM with a Transition-Edge Sensor type Microcalorimeter EDS detector

Hara T.1, Tanaka K.2, Maehata K.3, Mitsuda K.4, Yamasaki Y. Y.4, Yamanaka Y.5
1National Institute for Materials Science, Tsukuba, Ibaraki, Japan, 2Hitachi High-tech Science, Corp., Hitachi High-Tech Science, Corp., Oyama-cho, Shizuoka, Japan, 3Kyushu University, Fukuoka, Fukuoka, Japan, 4Japan Aerospace Exploration Agency, Sagamihara, Kanagawa, Japan, 5Taiyo Nippon Sanso Corp., Tsukuba, Ibaraki, Japan
HARA.Toru@nims.go.jp

X-ray spectroscopy is widely used for compositional analysis in a TEM. However, the accuracy and sensitivity of this method has not been realized to the required level from recent advanced materials research. One of the main reasons preventing accurate analysis is the low energy resolution of the detector itself. The energy resolution of a standard SSD(Si(Li) type) detector is approximately 130eV, which results in considerable peak overlap. To solve this problem, we have attempted to use a superconductor transition-edge sensor(TES) type microcalorimeter with a TEM as an EDS detector to improve the quality of compositional analysis(1).

Figure 1 shows an outlook of the first prototype of the TES-EDS mounted on a TEM. The characteristic points of this system are as follows: (i) Cryogen-free cooling system, based on a combination of a mechanical (GM type) and a dilution refrigerator, is newly developed(2). (ii) An X-ray polycapillary is applied to increase detecting solid angle. Figure 2 is an example spectrum from silicon device (Si+W) taken for system confirmation. It is well-known that a standard Si(Li) detector (dotted line) cannot separate adjacent peaks; i.e., the Si Ka and W Ma lines are overlapped to each other. As shown in the figure, the developed TES detector (solid line) can separate them clearly. The FWHM of the silicon Ka peak is 7.8eV, which is more than tenfold higher than that obtained by the standard Si(Li) detector.

The spectrum shown in Fig.2 was taken with a single-pixel TES detector mounted on the TEM (not STEM) with LaB6 thermal emitter (Fig. 1). The acceptable count rate of this detector is very low, approximately 100 cps., From these reasons, an EDS map couldn’t be obtained. In order to obtain EDS map with sufficient count rate, we are now developing a multiple-pixels detector system mounted on a STEM. Figure 3 is a current situation of the developing new system; we have succeeded to obtain an X-ray map with a single pixel TES detector and confirmed the mapping function can correctly work. Multipixel detector system is now under developing to increase count-rate in order to obtain an EDS map effectively.

References:

(1) T.Hara, et al.; “Microcalorimeter-type energy dispersive X-ray spectrometer for a transmission electron microscope”, J. Electron Microsc., 59(1),(2010),17-26

(2) K.Maehata, et al.; “A dry 3He-4He dilution refrigerator for a transition edge sensor microcalorimeter spectrometer system mounted on a transmission electron microscope”, Cryogenics,(2014), in press.


This work has been financially supported by the MEXT Leading Project and JST-Sentan program. The authors acknowledge JEOL Ltd. and Hitachi High-Tech corp. for their cooperation.

Fig. 1: Figure 1. Transition-Edge Sensor type microcalorimeter EDS mounted on a TEM.

Fig. 2: Figure 2. Comparison between the TES (solid line) and the SSD(Si(Li)) (dotted line).

Fig. 3: Figure 1. Transition-Edge Sensor type microcalorimeter EDS mounted on a TEM.

Type of presentation: Poster

IT-5-P-2735 Quantitative Position-Averaged Core-Loss Scattering in STEM

Zhu Y.1, Dwyer C.2
1Monash Centre for Electron Microscopy (MCEM) and Department of Materials Engineering, Monash University, VIC 3800, Australia, 2Ernst Ruska-Centre for Microscopy and Spectroscopy with Electrons, and Peter Grunberg Institute, Forschungszentrum Julich, D-52425 Julich, Germany
ye.zhu@monash.edu

With the rapid development in core-loss spectroscopic mapping in the scanning transmission electron microscope (STEM), it raises the need for the quantitative interpretation of core-loss intensity and map contrast. However, a quantitative interpretation of atomic-resolution chemical maps is not straightforward due to the dynamical scattering of the electron probe. The effect of dynamical scattering on core-loss maps bears similarities to its effect on annular-dark-field (ADF) images. Dynamical-scattering calculations are thus required to interpret both the core-loss map and ADF image intensities on a quantitative level.


Based on recent progresses on quantitative STEM imaging and inelastic multislice simulations, we have performed a quantitative comparison between experimental position-averaged core-loss scattering from K-, L- and M-shells of various elements and simulations based on a single-particle description of the core-loss process. The materials we studied include single-crystal Si, LaB6, SrTiO3, and LaAlO3. To facilitate a direct comparison free of adjustable or compensating parameters, we compare absolute scattering cross-sections for zone-axis-aligned crystals whose thicknesses have been measured independently using convergent electron beam diffraction (CBED). Our study of the position-averaged scattering avoids the complexity and any errors associated with evaluating the effects of aberrations and source size. Experimental results are compared with simulations that include an accurate description of multiple elastic and thermal-diffuse scattering both prior and subsequent to the core-loss events (double-channelling). In order to exclude any pronounced solid-state effects, which are not included in our simulations, we have considered discrete-continuum transitions that are at least 30 eV above edge onsets. The results show that the double-channelling simulations based on a single-particle model quantitatively predict the position-averaged scattering from K-shells, as well as that from L-shells in some cases (Si-L2,3). On the other hand, limitations of the single-particle picture are clearly revealed by the discrepancies in the case of M-shells (La-M4,5). Our results represent a critical step towards quantitatively predicting the absolute intensity and contrast in core-loss chemical maps with nano- or even atomic-resolution.


This work was supported by the Australian Research Council (ARC) grant DP110104734. The FEI Titan at Monash Centre for Electron Microscopy was funded by the ARC Grant LE0454166.

Fig. 1: (a) Experimental (left) and Bloch-wave simulated (right) PA-CBED pattern from 108 nm [110] Si. (b) PA-CBED determined thickness versus t/λ on [110] Si. (c-d) Experimental and simulated (c) ADF and (d) BF average intensities as a function of PA-CBED determined thickness on [110] Si.

Fig. 2: (a) Raw EEL spectra showing Si-L edge at different thickness. (b) Background-subtracted Si-L2,3 edge. The average intensity in the region highlighted in gray was compared to simulation. (c) Experimental (scattered) and simulated double-channelling (line) Si-L2,3 edge intensity at 45 eV above the edge onset.

Type of presentation: Poster

IT-5-P-2785 Implementation of the Zeta-factor method for quantitative EDS

Falke M.1, Kaeppel A.1, Nemeth I.1, Terborg R.1
1Bruker Nano GmbH, Berlin, Germany
meiken.falke@bruker-nano.de

Energy dispersive X-ray spectroscopy is well-established for composition analysis of electron transparent samples in STEM, TEM and SEM. This contribution reports on the implementation of the Zeta-factor method [1, 2] as an absolute EDS quantification method and opposed to the widely used relative Cliff-Lorimer method. The latter can provide quantitative data on the accuracy level of a few at% already and if using large solid and take-off angles even ppm. The quantitative results from the Cliff-Lorimer method are only valid relative to a standard though. Additionally, it has to be assumed that absorption and fluorescence effects can be neglected or it has to be realized that the thickness and composition of the sample of interest are close to the thickness and composition of the standard used, so that absorption and fluorescence cancel out.

An alternative quantification procedure, the Zeta-factor method, has been suggested and developed by M. Watanabe. It includes information on the beam current, sample thickness and density for the standard and can thus provide the absolute quantification of sample compositions while accounting for absorption and fluorescence effects. To obtain all this data, the just mentioned parameters must be well known for a standard sample and it must be possible to measure the beam current during the experiment with the sample of unknown composition and unknown thickness as well.

The Zeta-factor method is currently being implemented and tested in the Bruker ESPRIT software using various standards. For an initial test procedure a 30nm Si3N4 foil (commercially available from Agar) was used as a standard. The foil was punctured by the electron beam to produce folds of known thickness and composition. The spectrum from one of these areas (Fig 1) was processed to determine the net count number for individual X-ray lines and then to compute the respective ζ-factors for Si-K and N-K. Those ζ-factors were then tested on a Si3N4 sample region of a different well known thickness and vice versa.

The experimentally determined ζ-values can be used to calculate a proportionality factor to the respective Cliff-Lorimer factors, theoretically obtained from available atomic data and considering the detection geometry. Based on the experimentally specified Zeta/Cliff-Lorimer factor ratio the Zeta-factors for all element K-lines can then be calculated (Fig.2). Further tests on more complicated material systems including absorption and fluorescence effects are necessary.

[1] Watanabe M, Horita Z, and Nemoto M, Ultramicroscopy 65 (1996) 187–198
[2] Watanabe M. & Williams D.B, J. of Micr. Vol. 221. (2006) 89-109.


We gratefully acknowledge helpful discussions with W. Grogger and S. Fladischer from the ZELMI in Graz, Austria.

Fig. 1: Two areas of the Si3N4 foil and the respective spectra used for testing the Zeta-factor method implementation. The 30nm Si3N4 foil was folded after rupture in the electron beam, so that well-known sample parts of 30nm and 60nm thickness were available for test measurements.

Fig. 2: ζ-factors for K lines calculated based on the theoretically determined Cliff-Lorimer-factors and the N- and Si-ζ-factors obtained experimentally using Si3N4 as the standard.

Type of presentation: Poster

IT-5-P-2909 Advantages of Combining EDS and Energy Filtered STEM Diffraction at Atomic Level

Longo P.1, Aitouchen A.1, Rice P.2, Topuria T.2, Twesten R. D.1
1Gatan Inc., Pleasanton CA, USA, 2IBM Research Division, San Jose CA, USA
plongo@gatan.com

Very recently a technique called STEM Diffraction has been used to collect the entire diffraction pattern (DP) in STEM mode at each probe position and store it in a data-cube [1]. The advantage is that every feature in the DP can now be recorded at each probe position. In this way it is possible to take advantage of the wealth of information present in the DP and have the spatial resolution offered by the STEM probe. In addition, virtual STEM detectors can be created and images can be generated by integrating the intensity of the DP over an appropriate angular distribution that depends on the particular element being imaged. For instance, light elements show high contrast within a narrow angular distribution at low angle in the DP and this is analogous to medium annular bright field (MaBF) imaging (2).

 

Here we propose an alternative approach where EDS spectra and DPs are acquired simultaneously at atomic level. DPs are acquired using the camera attached to an energy filter in EFTEM mode with a 10eV slit inserted in order to remove the inelastic scattering effects that blur the contrast in the DP. EDS can be used to generate elemental distribution maps. However, even with the use of the latest generation of detectors, EDS is not sensitive towards light elements especially when imaged at atomic level. STEM Diffraction can be used to generate and deliver all the additional structural information present in the DP and also images of atomic columns containing only light elements by integrating over low angular distributions in the DP.

 

As example, a EDS/(energy filtered) STEM Diffraction dataset was taken at atomic level across a SrTiO3/LaFeO3 (STO)/(LFO) interface as shown in Figure 1. Figures 2 show DPs extracted from the Ti +O, Sr and pure O columns respectively. Features in the DP at every angular distribution are different across each atomic column and can be used to generate additional structural information. Figure 3 shows color maps of Sr in red, Ti in green, La in amber, Fe in light blue acquired using EDS while pure O columns in blue obtained by integrating the signal over 0 – 8mrad angular range in the DP. It is quite interesting to notice how the position of pure O columns is slightly asymmetric in the LFO region compared to the STO. This might be due to the presence of strain at the interface whose effects can create some crystal distortion in the LFO region. This paper will show other examples and discuss the advantages of this combined EDS/STEM Diffraction approach.

 

References:

1) Okunishi E. et al., Micron 43 (2012) 538-544

2)Findlay S.D. et al., Ultramicroscopy 136 (2014) 31-41


The authors would like to thank IBM for providing the TEM samples and access to their microscope installation 

Fig. 1: ADF STEM survey image. The green box is the area where the beam was scanned for the acquisition of the EDS/STEM Diffraction dataset. According to the image the interface STO/LFO seems quite nice and abrupt.  

Fig. 2: Series of Diffraction patterns from the STO region and extrcated from the Ti + O, Sr and pure O atomic columns. It is interesting to notice how features in the diffraction pattern at low and high angular range are different in each atomic column

Fig. 3: Elemental color maps of Sr in red, Ti in green, La in amber and Fe in light blue obtained using EDS and pure O atomic columns in blue obtained selecting a 0 - 8mrad angular distribution in the diffraction pattern. The position of the pure O atomic column appears to be slightly asymmetric in the LFO area compared to the STO.

Type of presentation: Poster

IT-5-P-2927 Evaluation of valence state in manganese oxide by transition edge x-ray sensor.

Tanaka K.1,5, Ohgaki M.2,5, Miyayama M.3,5, Matsumura S.4
1Hitachi High-Tech Science Corporation, Shizuoka, Japan, 2Hitachi High-Tech Science Corporation, Tokyo, Japan, 3The University of Tokyo, Tokyo, Japan, 4Kyushu-University, Fukuoka, Japan, 5Japan Science and Technology Agency, CREST, Tokyo, Japan
tanaka-keiichi@hhs.hitachi-hitec.com

Electrochemical capacitors are promising candidates for future energy storage devices because of their high power density, long cycle life, and relatively high energy density. There has been considerable interest in MnO2 as a cathode material for such capacitors because of its low toxicity, environmental safety, cost effectiveness, and large capacitance. In particular, two-electron redox reactions involving Mn2+ and Mn4+ are expected to yield a high energy density. Valence states in transition metals are often studied by determining the branching ratio of the L3 and L2 absorption edges using transmission electron microscopy together with electron energy loss spectroscopy (TEM-EELS). In the case of Mn, EELS can distinguish the L3 (640 eV) and L2 (651 eV) absorption edges, and an adequate signal can be obtained.
In the present study, an attempt was made to evaluate the valence state for manganese oxide particles using scanning electron microscope with a transition edge sensor (SEM-TES). A TES is a kind of energy dispersive X-ray spectrometer, but it has a very high energy resolution (typically 10 to 15 eV) and can separate extremely close X-ray peaks. The valence state was determined based on the branching ratio for the Lb and La X-ray emission lines. It was found to be possible to determine the valence state even in stacked structures with layer thicknesses of about 100 nm. Positive electrode powder samples with a composition Hx(Ni1/3Co1/3Mn1/3)O2 were evaluated using both SEM-TES and TEM-EELS. Three types of samples were examined: charged (as-prepared), discharged and recharged. Using SEM-TES, the branching ratios for the three samples were determined to be 0.76, 0.699 and 0.73, respectively. Using TEM-EELS, the values obtained were 2.3, 5.2 and 2.6, respectively. The discharge sample changed valence state from as-prepared one and discharged one. Thus, the SEM-TES shows the capability of identifying a different valance state in the case of the discharged sample.


This work was supported in part by the CREST program, JST, MEXT, Japan. The SEM-TES measurements were performed as part as a research program (A-12-KU-0018) of the Nanotechnology Platform Project conducted by MEXT.

Type of presentation: Poster

IT-5-P-2960 Advanced EELS Spectrum Imaging of RRAM Devices: Chemical State and Three-Dimensional Element Mapping

Chang M. T.1, Lo S. C.2, Hsieh C. Y.3
1Dept. of Electron Microscopy Development and Application, Material and Chemical Research Laboratories, Industrial Technology Research Institute (ITRI)
mtchang@itri.org.tw

The full text of the abstract is not available. Please contact the presenting author.


Type of presentation: Poster

IT-5-P-2963 TEM-EELS/SXES studies on electronic structures of p-type CaB6

Terauchi M.1, Saito T.1, Sato Y.1, Inayoshi K.2, Takeda M.2
1IMRAM, Tohoku University, Sendai, Japan., 2Nagaoka University of Technology, Nagaoka, Japan
terauchi@tagen.tohoku.ac.jp

Metal hexaboride MB6 is based on a network of B6-clusters located on each corner of cubic unit cell. M atom occupies at the body center position of the unit cell. When an M atom can supply two electrons to B6-network, the valence bands (VB) of this material is fully occupied and becomes a semiconductor. Those semiconductor materials have been investigated as a candidate for a high-temperature thermoelectric-power material. Seebeck coefficients of MB6 (M=Ba, Sr, Ca) synthesized by solid-state reaction method were reported to be negative, indicating those are n-type materials [1]. Recently, the p-type character for CaB6 synthesized from the mixture of CaCl2 and NaBH4 in eutectic LiCl-KCl molten salt was reported [2]. Thus, the electronic structure of this new material has been studied by using electron energy-loss spectroscopy (EELS) [3] and soft-X-ray emission spectroscopy (SXES) [4], which are methods for probing over and below the Fermi energy level, respectively.

EDS analysis of the p-type CaB6 showed an inclusion of a few % of Na. Electron diffraction patterns showed a good crystalline order. As one Na atom can transfer one electron to B6-network, Na-doping can be a hole-doping based on the rigid band structure scheme when doped Na atoms occupy Ca site. Valence electron excitation (from VB to conduction bands: CB) EELS spectra showed smaller bandgap energy of 1.5 eV than 2.5 eV of n-type CaB6 synthesized by solid-state reaction method. B K-shell excitation EELS spectra of p- and n-type materials showed almost the same onset energy, which energy position corresponds to the bottom of CB.

Figure 1 shows B K-emission SXES spectra of p-type and n-type CaB6. Those spectra show different intensity distribution especially at the top region of VB, which correspond to the right hand side end of the intensity distribution. The intensity distribution of p-type material apparently extends into the bandgap region of n-type CaB6. Since the bottom of CBs of the two materials were the same, this higher energy position of the top of VB of p-type should be the origin of the smaller bandgap energy of p-type material. This is consistent with the result of valence excitation EELS experiment stated above. This indicates that the doping of Na atoms into the Ca site of CaB6 causes not only the creation of holes in VB but also a change the energy state at the top region of VB, not a simple rigid band structure scheme.

[1] M.Takeda et al., J. Solid State Chemistry, 179, 2823-2826 (2006).

[2] K.Inayoshi and M. Takeda, IUMRS-ICEM (2012).

[3] Y.Sato et al., Ultramicroscopy 111, 1381-1387 (2011)

[4] M.Terauchi et al, Journal of Electron Microscopy 61, 1-8 (2012).


Fig. 1: B K-emission SXES spectra of p-type and n-type CaB6.

Type of presentation: Poster

IT-5-P-2969 Valence Electron States of Carbon Materials studied by TEM-SXES

Terauchi M.1
1IMRAM, Tohoku University, Sendai, Japan.
terauchi@tagen.tohoku.ac.jp

X-ray emission spectroscopy is widely used as a practical tool for compositional analysis of local specimen area and elemental mapping analysis in electron microscopes. X-rays originate form electronic transitions from valence bands (VB, bonding electron states) to inner-shell electron levels inform us energy states of bonding electrons. This X-ray energy ranges in ultrasoft or soft X-ray region form about 0.1 to a few keV. Thus, soft X-ray emission spectroscopy (SXES) based on electron microscopy (EM) can be a sensitive tool for elemental and chemical identifications. For that purpose, we have developed and tested SXES instruments by applying to TEM, EPMA, and SEM [1,2,3]. This SXES spectrometer informs us energy states of VB from specified specimen areas in electron microscopy, which is hardly obtained by EELS and EDS.

Figure 1 shows carbon K-emission (VB→K-shell) spectra of zeolite-Y template carbon (ZTC) [4]. Spectra of graphite and momomer-C60 (Mono.-C60) are also shown for comparison. As this ZTC has a huge surface area of 4000 m2/g, it is a candidate material for applying to fuel cell and electrode of rechargeable batteries. Electron diffraction pattern of ZTC shows amorphous like broad rings. However, C K-emission spectrum shows apparent structures. Those structure positions similar to those of Mono.-C60 than those of graphite. This result suggests that ZTC is mainly composed of curved grapheme, sp2, network with a similar curvature with that of a C60 cluster. Additional structures at the top end and on the lower energy part of VB as indicated by arrows suggest a presence of a certain amount of sp3 component in carbon network of ZTC examined.

When a crystal has anisotropic bonding nature, emission intensities originate from VB electrons should be anisotropic. TEM based SXES experiment can examine this anisotropic emission intensity by changing the crystal orientation. Analyses of anisotropic intensity of C K-emission of graphite by using TEM-SXES have already demonstrated [5]. However, this analysis did not include the effect of polarization on a reflectance of grating used. When the polarization effect was included, the resultant was improved (not shown here). This indicates that the polarization correction is presumably necessary for an accurate analysis of VB of anisotropic crystalline materials by using a grating spectrometer.

[1] M Terauchi et al, Journal of Electron Microscopy 61 (2012), 1.

[2] H Takahashi et al, Microscopy and Microanalysis 19(Suppl.2) (2013), 1258.

[3] M Terauchi et al, Microscopy and Microanalysis, accepted.

[4] K Nueangnoraj et al., CARBON 62 (2013), 455.

[5] M Terauchi in “Transmission Electron Microscopy Characterization of Nanomaterials”, ed. CSSR Kumar, (Springer-Verlag, Berlin Heidelberg) 284.


Fig. 1: C K-emission spectra of zeolite-Y template carbon (ZTC), monomer-C60 (Mono.-C60) and graphite. ZTC shows a similar structure with those of Mono-C60 than those of graphite.

Type of presentation: Poster

IT-5-P-2998 STEM EELS Analysis of 2D Layered Inorganic Materials at Atomic Resolution

Nerl H. C.1, McGuire E. K.1, Backes C.2, Seral-Ascaso A.2, Ramasse Q.3, Houben L.4, Nicolosi V.1
1Centre for Research on Adaptive Nanostructures and Nanodevices (CRANN) and Advanced Materials and Bio-Engineering Research (AMBER), Trinity College, Dublin, D2 Dublin, Ireland. , 2School of Physics, Trinity College, Dublin, D2 Dublin, Ireland. , 3SuperSTEM Laboratory, STFC Daresbury, United Kingdom., 4Ernst Ruska-Centre for Microscopy and Spectroscopy with Electrons, Forschungszentrum Jülich GmbH, Jülich, Germany., 5Advanced Materials and Bio-Engineering Research (AMBER), Trinity College, Dublin, D2 Dublin, Ireland.
nerlh@tcd.ie

In recent years, methods for dispersion and exfoliation of 2D nanostructures of a range of nanomaterials have been successfully developed [1-8], opening up numerous possibilities for a range of innovative technologies [4, 6-10]. Chemical and physical properties of materials can however change when going from bulk material to the 2D state. To make real applications feasible there is a need to fully characterize these nanostructures at an atomic scale. Due to the recent advances in transmission electron microscopy (TEM), scanning TEM (STEM) imaging and STEM electron energy-loss spectroscopy (EELS) can now be used to study the structure and composition of nanomaterials, atom by atom [11]. The focus of the study presented will be on characterization of inorganic 2D layered materials produced by liquid phase exfoliation [3,4], a high-yield method for producing sheets of few atomic layers thickness for a range of materials. Atomic resolution STEM EELS analysis of these sheets allows the determination of the atomic structure, structural defects as well as electronic properties of the material, giving insight into their fundamental physical and chemical properties.

[1] AK Geim and KS Novoselov. The rise of graphene. Nature Materials 6 183 (2007)

[2] SD Bergin et al. Towards Solutions of Single-Walled Carbon Nanotubes in Common Solvents. Advanced Materials 20, 10 1876 (2008)

[3] Y Hernandez et al. High-yield production of graphene by liquid-phase exfoliation of graphite. Nat. Nanotechnol. 3, 563 (2008)

[4] JN Coleman et al. Two-dimensional nanosheets produced by liquid exfoliation of layered materials. Science 331, 568–571 (2011)

[5] M Chhowalla et al. The chemistry of two-dimensional layered transition metal dichalcogenide nanosheets. Nat. Chem. 5, 263-275 (2013)

[6] V Nicolosi et al. Liquid Exfoliation of Layered Materials. Science 340, 1226419 (2013)

[7] AK Geim. Graphene: Status and Prospects. Science 324, 1530-1534 (2009)

[8] KS Novoselov et al. A roadmap for graphene. Nature 490, 192-200 (2012)

[9] QH Wang, K Kalantar-Zadeh, A Kis, JN Coleman & MS Strano. Electronics and optoelectronics of two-dimensional transition metal dichalcogenides. Nat. Nanotechnol. 7, 699-712, (2012)

[10] M Osada & T Sasaki. Exfoliated oxide nanosheets: New solution to nanoelectronics. J. Mater. Chem. 19, 2503 (2009)

[11] OL Krivanek et al. Atom-by-atom structural and chemical analysis by annular dark-field electron microscopy. Nature 464, 571–574 (2010)


Support from the Advanced Microscopy Laboratory (AML), Science Foundation Ireland, Enterprise Ireland and the European Research Council (ERC) is gratefully acknowledged.

Type of presentation: Poster

IT-5-P-3086 On the usability of electron vortices as probes for atomic resolution EMCD experiments

Pohl D.1, Schneider S.1, 2, Rusz J.3, Schultz L.1,2, Rellinghaus B.1
1IFW Dresden, P.O. Box 270116, D-01171 Dresden, Germany, 2TU Dresden, Institute for Solid State Physics, D-01062 Dresden, Germany, 3Uppsala University, Department of Physics and Astronomy, SE-752 37 Uppsala, Sweden
d.pohl@ifw-dresden.de

Recently discovered electron vortex beams, which carry a discrete orbital angular momentum (OAM) L, are predicted to reveal dichroic signals comparable to classical electron magnetic circular dichroism experiments (EMCD) [1]. Since electron beams can be easily focused down to sub-nanometer diameters, this novel technique provides the possibility to quantitatively determine local magnetic properties with unrivalled lateral resolution. For this purpose, specially designed apertures are needed to generate such non-planar electron waves [2]. Dichroic signals on the L2- and L3- edges are expected to be of the order of around 5% [3,4].

We have prepared and successfully implemented a spiral aperture into the condenser lens system of a FEI Titan3 80-300 transmission electron microscope (TEM) equipped with an image CS corrector (cf. fig. 1a). This setup allows for the generation of focused electron vortex beams with user-selectable OAM that can be used as probes in scanning TEM (STEM). Since for such spiral apertures, the different OAM are dispersed along the beam direction (z direction), the selection of the OAM is obtained by defocussing the beam.

First investigations aimed at probing the presence of an EMCD signal with such vortex beams were conducted on a 20 nm thin polycrystalline Ni film prepared by RF sputtering. Fig. 1b) shows the resulting EEL spectra subsequently acquired with L = +1 and L = -1 vortex states, respectively. In order to improve the signal-to-noise ratio, the binned-gain acquisition technique was used [5].
As can be seen from fig. 1b), these first experiments do not provide any evidence for differences in the absorption edges in the two EELS spectra.
In addition, the generation and propagation of the vortex wave functions and the spatial distributions of the OAM were simulated. The results of these simulations show that the orbital momenta and the beam intensity are largely localized (in all three dimensions) symmetrically around the geometrical focal points which are paraxial to the vortex cores (cf. fig. 2). Despite this localization, the superposition of contributions of vortex states (e.g., L = 0 and L = -1) adjacent on the one selected by appropriate defocussing (e.g., L = +1) are large enough that the average OAM is close to zero h, if the defocused portions of the wave are not properly prevented from interacting with the sample (cf. fig. 2b) which explains the absence of a dichroic signal in the experiment.

[1] J. Verbeeck et al., Nature 467 (2010), p. 301-304.
[2] J. Verbeeck et al., Ultramicroscopy 113 (2012), p. 83-87.
[3] P. Schattschneider et al., Ultramicroscopy 136 (2013), p. 81-85.
[4] J. Rusz and S. Bhowmick, Phys. Rev. Lett. 111 (2013), 105504.
[5] M. Bosman and V. J. Keast, Ultramicroscopy 108 (2008), p. 837-846.


Fig. 1: SEM image of the spiral aperture installed in the C2 aperture plane of the electron microscope.

Fig. 2: Ni L3/L2 edges in the normalized EEL spectra acquired with electron probes of vortex states with orbital angular momenta L = +1 and L = -1, respectively. No significant EMCD signal (= difference between the two spectra) is observed.

Fig. 3: Intensity profile of the simulated electron vortex beam with L = +1 generated with a spiral aperture. Arrows indicate radial positions where contributions of the L = 0 and L = -1 states become dominant.

Fig. 4: Normalized expectation value of the OAM as obtained from radial integration from the vortex core to a given radius (1000 a.u. correspond to roughly 130nm).

Type of presentation: Poster

IT-5-P-3124 Core/shell structure in magnetic nanoparticles from HRTEM and EELS

Bertoni G.1,2, López-Ortega A.3, Lottini E.3, Sangregorio C.3,4, Turner S.5, de Julián Fernández C.1,3, Salviati G.1
1CNR-IMEM, Parco Area delle Scienze 37/A, 43124 Parma, Italy, 2Istituto Italiano di Tecnologia, Via Morego 30, 16163 Genova, Italy, 3INSTM and Dipartimento di Chimica , 4CNR-ICCOM and INSTM Via Madonna del Piano 10, Sesto Fiorentino 50019 Firenze, Italy, 5EMAT, University of Antwerp, Groenenborgerlaan 171, 2020 Antwerp, Belgium
giovanni.bertoni@imem.cnr.it

Magnetic properties on nanoscale materials are strongly affected by the possible different composition of the surface with respect to the bulk material.

In this study we investigate the composition and morphology of spherical d ≈ 10 nm magnetic particles, presenting a core/shell structure, synthesized through thermal decomposition of Fe-Co oleate under high boiling point solvents and posteriorly controlled surface oxidation. Core-shell structure has been assessed by X-ray diffraction analysis and corroborates the formation of two differentiated crystallographic phases: Co doped iron oxide spine (magnetite-like) and Co doped iron monoxide (wustite-like). From a careful inspection of HRTEM images (acquired on an aberration-corrected JEOL JEM-2200FS), an extra reflection from the shell region is indeed visible. This is compatible with a (220) reflection from magnetite. The (400) magnetite and (200) wustite reflections are indeed close (about 0.21 nm), indicating that some solubility or epitaxy of the two structures is possible.

HAADF images and spatially resolved EELS maps are acquired an aberration-corrected FEI Titan “cubed” microscope equipped with an electron monochromator and Gatan Enfinium. The core/shell structure of the particles is evident in HAADF. The reflection from magnetite is visible in the shell region. Moreover EELS maps, obtained by model-based fitting, reveals that there is a depletion of cobalt in the magnetite shell, together with the expected increasing in oxygen with respect to wustite. This is accompanied by the change in valence state of the transition metals from +2 to +3, as verified on Fe-L2,3 edge by fitting reference spectra. The energy position of the transition metal edge onset is indeed a valid parameter for determining its valence state.(1)

By further oxidation of the particles the core/shell structure seems to fade, the EELS maps revealing a more uniform distribution of Fe2+ and Fe3+ ions, indicating the magnetite phase extends towards the core. The final particles show higher policrystallinity with respect to the pristine particles. We can conclude that after further oxidation, the two phases are mixed in the whole volume. Further simulations from EELS profiles and/or EDS are in progress for quantifying the shell extension.

[1] H. Tan et al., Ultramicroscopy 116 (2012) 24–33


European Union FP7 Grant Agreement 312483 ESTEEM2 (Integrated Infrastructure Initiative–I3) and the European Union FP7 project 310516 NANOPYME

Fig. 1: a) HRTEM image from a Fe-Co oxide particle, showing an orientation close to [001] of a cubic structure. The FFT from the core can be addressed with wustite (W), while the shell region has clear reflections from (220) of magnetite (M). b) The core/shell structure as revealed in HAADF

Fig. 2: a) Color map from model-based quantification; b) Linescan profile of the particle, showing the concentration of Fe, Co, and O, respectively; c) Valence map for Fe after fitting of reference spectra for Fe2+ and Fe3+ (explained in e); d) After further oxidation, there is no evidence of a core/shell structure

Type of presentation: Poster

IT-5-P-3210 Mapping SPR of Au metallic nano-objects with complex morphologies and environments

Florea I.1,2, Arenal R.3,4, Tréguer-Delapierre M.5, Ihiawakrim D.1, Hirlimann C.1, Ersen O.1
1IPCMS,CNRS/UdS, 23 rue du Lœss, 67034, Strasbourg Cedex , 2LPICM,Ecole Polytechnique, Route de Saclay, 91128 Palaiseau Cedex, France, 3Fundacion ARAID, 50018 Zaragoza, Spain. , 4LMA, INA, Universidad de Zaragoza, 50018 Zaragoza, Spain., 5ICMCB, CNRS, Bordeaux University, 87 av. Dr. A. Schweitzer, Pessac F-33608, France
lenuta-ileana.florea@polytechnique.edu

The remarkable optical properties of the metal nanoparticles(NPs), governed by the excitation of localized surface plasmon resonances (LSPRs), whose character and resonant energy depend on their size, shape, composition and environment, take nowadays a prominent position in research and applications. Imaging the LSPRs can be possible using various techniques such as scanning near optical microscopy(SNOM), photon electron emission microscopy(PEEM), electron energy loss spectroscopy(EELS)and photochemical imaging. Specifically,the EELS technique developed in the scanning(STEM) imaging mode of an electron microscope with a nanometer spatial resolution was often used for probing the SPR of metallic nanoparticles presenting classical morphologies such as spheres, cubes, rods, triangles.[1-3] Along this line the key-issue addressed by this work is the assessment of the optical response of the Au metallic NPs presenting more complex morphologies in presence of particular environments. The most important aspect on which we focused over the entire study relates on accessing information regarding, first the presence and the localization of the “hot spots” or areas as well as coupling effects between the LSPRs modes, and second the influence of the environment of Au NPs on the SPR modes. Regarding the effect of the NPs environment two aspects were closely investigated: the circular nature of the surface supporting the NPs and its dielectric environment.For the STEM-EELS experiments different systems with complex morphologies(see Fig.1) were considered: Au individual BP, Au assembled-BPs forming 2D clustered-structures; silica coated Au BPs;Au NPs presenting a patch morphology and silica beads casted inside the Au patches.At first, a close analysis of the monochromated EELS spectra recorded on a single Au NPs(see Fig.2) allowed identifying two main SPR centered at 1.5 and 2eV.This result helped us later in the analysis of the EELS spectra recorded when the more complex morphologies were studied. More exactly for the systems consisting of enchained Au BP we found that the presence of a neighbor NP induces slight modifications in the SP mode with respect to their position. For the other systems, the silica coated Au BP, Au patch and silica bead casted inside the Au patch the analysis of the plasmon spectra taken at different areas on the NP enabled us observing that the presence of the circular support as well as the dielectric environment are inducing energy shifts of some of the LPSR modes.

[1] J. Nelayah; et al., Nat. Phys.(2007)

[2] R. Arenal, Microsc. and Microanal.(2011)

[3] S. Mazzucco et al., Nano Lett.(2012)


The work was supported by the ANR under Grant no. ANR-BLANSIMI10-LS-100617-15-01. The research leading to these results has received funding from the European Union Seventh Framework Programme under Grant Agreement 312483 - ESTEEM2-I3.

Fig. 1: Au complex morphologies: (a) HAADF image of an agglomerate of Au bipyramids ;(b) BF-TEM image of a Silica coated Au bypiramid;(c) HAADF image of a Au NPs presenting a patch morphology; (d) HAADF image of a silica bead casted inside the Au patch NPs.

Fig. 2: STEM-EELS analysis of the SPRs of an Au BP: (a) HAADF image of a Au BP; (b) Map of the energy of the detected surface plasmon resonances; (c) EELS spectra extracted from the SI registered on the whole area of the Au NPs after ZLP subtraction.

Fig. 3: STEM-EELS analysis of the SPRs of Au patches NPs: HAADF images of: (a) pure Au and (b) with a silica bead casted inside the NP; (c) EELS spectra extracted from the SI registered on the analyzed area after ZLP subtraction.

Type of presentation: Poster

IT-5-P-3251 Parallel acquisition Auger electron spectroscopy

Walker C. G.1, Zha X.2, El-Gomati M. M.1
11-Department of Electronics, University of York, Heslington, York, YO10 5DD., 2York Probe Sources Ltd, York, YO10 5DD
MOHAMED.ELGOMATI@YORK.AC.UK

Auger electron spectroscopy (AES), a surface analysis technique, has traditionally required the use of Ultra High Vacuum (UHV) conditions on account of its high surface sensitivity and the rapidity with which surface become coated with contaminants under High Vacuum (HV) conditions. Energy Dispersive X-ray Spectroscopy (EDS) has limited resolution due to its large excitation volume. Imaging AES, thus achieves much higher spatial resolution than EDS even when the recently developed Silicon Drift Detector technology is employed. The introduction of techniques that acquire an electron spectrum in parallel will allow a much faster acquisition of AES spectra and thus relax the vacuum conditions required.
Two such parallel acquisition electron energy analyser is the Hyperbolic Field Analyser (HFA) [1] and the Magnetic Electron Energy Spectrometer (MEES) [2]. Figure 1 shows a schematic of the experimental setup used for acquiring Electron Energy Loss (EELS) data and AES data using the MEES analyser. Figure 2 shows an image acquired on an Active Pixel Sensor CMOS detector of an elastic peak using the MEES. The image shows the output of the two dimensional detector when the incident electron beam energy is 900 eV and the magnetic field is 80 Gauss. The x and y axes corresponds to an energy interval of ~700 eV to~1000 eV. The image data can be converted into a spectrum by integrating the image data over specific regions (here between the 2 straight lines from top left to bottom right on Figure 2). The resultant spectrum is shown in Figure 3.
The HFA has been used in an SEM (operating at HV; 10-6mbar) to acquire AES spectra. The samples are cleaned using argon ion bombardment as practiced in surface analysis and then rapidly analysed by AES a few seconds after the ion cleaning is ceased. An example of an AES spectrum from Indium is given in Figure 4. The increase in the carbon Auger signal on the samples can also be monitored over the next minutes as the surface is contaminated. This provides a simple demonstration of how parallel acquisition can monitor rapid changes in surface composition in a way that no other technique can. In this article, we also explore in greater detail the theoretical basis of the MEES and its potential as a device for use in Scanning Electron Microscopes (SEMs) for the high speed inspection of objects on the nanometre scale and show further spectra collected using images acquired with the Active Pixel Sensor.
References
[1] M. Jacka, M. Kirk, M.M. El Gomati, M. Prutton, “A fast, parallel acquisition, electron energy analyzer: the hyperbolic field analyzer”, Rev. Sci. Instrum. 70, (1999), 2282-2287.
[2] X. Zha, “Magnetic Electron Energy Spectrometer”, Ph.D. Thesis University of York, York, UK, (2009).


Fig. 1: Schematic representation of the magnetic analyser, MEES. (a) Helmholtz coils shown semi-transparently (b) sample (c) slit (d) sensor (e) board containing electronics for sensor (f) metal plate (g) electron column.

Fig. 2: The locus of the elastic peak is shown as detected by the CMOS sensor. Curve A is the elastic peak and lines B and C mark the limits of integration to create the spectrum in Figure 3.

Fig. 3: A spectrum determined from Figure 2 after image processing.

Fig. 4: The Auger spectrum of indium acquired using the HFA in HV vacuum conditions of an ordinary SEM (JEOL 6400F) and shows the In MNN Auger peak taken immediately after ion cleaning. A carbon contamination layer took about 10 minutes to build up after Ar ion cleaning and the acquisition of the AES data.

Type of presentation: Poster

IT-5-P-3289 Ultra-Fast, High-Resolution Silicon Drift Detectors for Accurate EDS Microanalysis in Electron Microscopes

Niculae A.1, Bornschlegl M.1, Eckhardt R.1, Herrmann J.1, Jeschke S.1, Krenz G.1, Liebel A.1, Lutz G.2, Soltau H.1, Strüder L.2
1PNDetector GmbH, 2PNSensor GmbH
adrian.niculae@pndetector.de

In the recent years significant advances have been done in electron microscopy instrumentation with respect to electron beam intensity and spot size, pushing for higher energy resolution and faster EDS detectors. High-resolution, ultra-fast EDS microanalysis applications require detectors with extremely low input capacitance, insuring optimum detector operation at very short processing times. A substantial development work has been done in the past years in this direction at PNDetector by remodeling the geometry of the anode and of the integrated FET with the goal of reducing all the parasitic capacitances related to the detector anode. This led to a new generation of Silicon Drift Detectors  – the so-called SDDplus series.
The low capacitance anode/FET can be adopted for all SDD types (round or droplet shape) and sizes (from 5 and 10 mm2 up to 100 mm2 or multichannel devices). Fig.1a and 1b show spectroscopic performances measured with the 30 mm2 and the 60 mm2 SDDplus detectors. Whereas energy resolution values of 126 eV are achieved with the round-shape SDDplus devices, when applied to the droplet-shaped SD3 devices, the low capacitance FET drives the energy resolution below 122 eV at shaping times as short as 1 us. With the detector operated at 0.5 us shaping time (maximum input count rate of 400 kcps) the energy resolution is still below 125eV. Further measurements with SDDplus devices of various sizes and shapes will follow.
The improved spectroscopic performance of the SDDplus devices becomes much more visible when it comes to detection of light elements. Combined with a high-performance, loss-free detector entrance window, the SDDplus devices demonstrate their excellent light element detection capabilities. Measured spectra from carbon samples in SEM are shown in Fig. 2a with an achieved energy resolution of 37 eV FWHM for a 10 mm2 SD3plus detector. Even energy lines well below 100 eV (Si-L, Al-l or Li-K) can still be well distinguished from the noise peak (see Fig2b).
When analyzing thin samples or biological probes with a low photon yield the measurement time is directly related to the detector collection angle. Another important development direction is moving toward smaller, more compact detector packages and therefore increasing the solid angle coverage of the detector with respect to the analyzed sample. An example here is the new large area 100 mm2 SDD detector which has been mounted onto a very compact package of 18.5 mm diameter only (see Fig. 3a). The spectroscopic performance is similar to that obtained with smaller size detectors (Fig. 3b) and this at moderate cooling temperature of -30°C. Selected measurements will be presented and the results will be discussed.


Fig. 1: Energy resolution vs. shaping time measured at -30°C with: (a) 30 mm2 SDDplus/SD3plus detectors; (b) 60 mm2 SDDplus/SDD detector

Fig. 2: Light element spectra of SDDplus devices: (a) C-K line (277 eV) and (b) Al-L line (70 eV)

Fig. 3: (a) 100 mm2 SDD in ultra-slim line package (b) energy resolution vs. shaping time at -30°C for the 100 mm2 SDDplus/SDD detectors

Type of presentation: Poster

IT-5-P-3319 Improvement of EFTEM acquisition and data processing using prior knowledge of camera DQE

Lucas G.1, Hébert C.1
1Interdisciplinary Center for Electron Microscopy (CIME), Ecole Polytechnique Fédérale de Lausanne (EPFL), Lausanne, Switzerland
guillaume.lucas@epfl.ch

In transmission electron microscopy (TEM), electrons are traditionally detected using a camera with either indirect or indirect detection. In both cases the “true” image is degraded by the non-ideal point spread function (PSF) of the detector leading to a blurring of the signal and the addition of stochastic noise components such as dark-current noise or readout-noise.

Detector performances can be assessed by the measurement of the modulation transfer function (MTF), the noise power spectrum (NPS) and the detective quantum efficiency (DQE) [1,2]. First of all, it allows us to verify the specifications provided by the manufacturers, to compare the relative performance of different detectors and to estimate their degradation over time. Secondly it can help to optimize the acquisition strategy for a given problem. Finally this information can be used as prior knowledge for data processing algorithms.

Energy filtered transmission electron microscopy (EFTEM) has been used to illustrate the benefits of the knowledge of the characteristics of the detection system. Images corresponding to different energy losses are sequentially recorded on the camera device, resulting in a 3D dataset for which each image plane is convolved with the PSF of the camera and each spectrum with the resolution response function of the spectrometer.

This works aims in a first step to measure the DQE of the Gatan US1000 camera used in our JEOL 2200FS microscope in order to improve our EFTEM acquisitions. The recently developed silhouette method [2] is used for the determination of the MTF. In a seconds step this works tries to apply principal component analysis [3,4] in order to perform the denoising of the data as well as the improvement of its spatial or spectral resolution by deconvolution techniques. The prior knowledge of the noise model and the MTF of the camera are embedded in the deconvolution algorithms in order to perform the regularization of the solutions in a realistic way.

The data processing procedure is demonstrated on a simulated dataset providing a ground truth for exploring the applications, limits and eventual pitfalls of the algorithms under known noise levels and MTF. After the measurement of the camera characteristics, acquisition parameters required for a good signal-to-noise ratio are optimized. The algorithms are then applied to the real datasets.

References:

[1] M. Vulovic et al, Acta Crystallographica Section D: Biological Crystallography. 66 (2010) p. 97-109.
[2] W. Van den Broek et al, Microscopy and Microanalysis 18 (2011) p. 336-342.
[3] G. Lucas et al, Micron 52-53 (2013) p. 49-56.
[4] Multivariate Statistical Analysis plugin for Digital Micrograph™, lsme.epfl.ch/msa.


Type of presentation: Poster

IT-5-P-3376 Comparison of the silicon/phosphorus ratio in natural and synthetic nagelschmidtite for possible use as standard for microanalysis based on X-ray lines of Si and P

Walther T.1
1University of Sheffield
t.walther@sheffield.ac.uk

Quantitative chemical microanalysis by energy-dispersive X-ray spectroscopy (EDXS) in a (scanning) transmission electron microscope (STEM) relies on the use of accurate k-factors. The most commonly used reference line is Si K.

For semiconductor research, standards for elemental semiconductors and for III/V compound semiconductors including elements from groups III and V of the periodic table are required. For the arsenides we have published results on X-ray quantification based on standards of InGaAs [1]. InAs and InP can be used to link arsenides and phosphides. However, the nominal k-factor for the P K-line in the ISIS software of k=1.000 indicates that this has probably not been measured at all.

Here, we use a natural and a synthetic sample of the mineral nagelschmidtite, a calcium silico-phosphate (ideal formula Ca7(SiO4)2(PO4)2 [2]), to evaluate the Si/P ratio from EDXS.

The natural mineral stems from the Hatrurim formation [3] and was cut from a thin section by a focused ion beam to produce an electron transparent specimen for TEM. Electron probe microanalysis (EPMA) of a larger inclusion of nagelschmidtite yielded an atomic ratio of Si/P=3.15. Results from TEM-EDXS are displayed in green.

The synthetic mineral was prepared in the laboratory of C Wu, Shanghai Institute of Ceramics [4]. Its chemical analysis using a Spectro Cirus Vision ICP-OES spectrometer gave a Si/P ratio of 0.36 (by at%), i.e. an almost inverted ratio. (S)TEM-EDXS results from these particles are displayed in dark blue.

Figure 1 shows that, for the detector setting used the deadtime of the detector is linearly related to the count rate up to a max of ~2500 counts/second or 50%, above which the detector runs into saturation.

Figures 2 and 3 plot atomic ratios as obtained from ISIS without absorption correction. The synthetic compound (blue) clearly reveals a higher P/Si ratio than the natural mineral (green) in Fig.2.

If the chemical concentration xn of an element n is proportional to the product of X-ray intensity In, k-factor kn,Si (for weight%) and absorption factor an, divided by the atomic weight An, then we can calculate an effective k-factor [5]:

keffP,Si = kP,Si aP,Si = (ISi xP AP) / (IP xSi ASi)

This is plotted in Fig.4. While the data scatter is rather large, a linear fit to the spectra that gave reasonable densities (≤3.5 gcm-3) as determined by ISIS allows us to determine the thin-film k-factor by extrapolation to zero count rate. The result is kP,Si=1.16±0.45 (R2=0.287). 

[1] T Walther, Proc EMAG2009, J Phys Conf Ser 241 (2010) 012016

[2] G Nagelschmidt, J Chem Soc 1 (1937) 865

[3] M Fleischer, LJ Cabri, GY Chao, A Pabst, Am Mineral 63 (1978) 424

[4] Y Zhou, C Wu, Y Xiao, Acta Biomaterialia 8 (2012) 2307

[5] Y Qiu et al, Proc EMC2008, 2 (2008) 643


Fig. 1:      

Type of presentation: Poster

IT-5-P-3430 The “equivalent sphere” approach to fitting surface plasmon energy loss spectra

Ostasevicius T.1, de la Peña F.1, Collins S. M.1, Ducati C.1, Midgley P. A.1
1Department of Materials Science and Metallurgy, University of Cambridge, Cambridge, UK
to266@cam.ac.uk

With the advent of modern cold field emission guns and commercial monochromators on Schottky sources, the remarkably high energy resolution available in the low loss part of the electron energy loss spectrum (EELS) has enabled detailed characterisation of low energy excitations such as localised surface plasmon resonances (LSPRs) in metallic nanoparticles [1]⁠. Analysis of the LSPRs has conventionally involved fitting characteristic peak shapes using a series of Gaussian or Lorentzian functions. Whilst this has proved to be successful in some cases, inherently the LSPR peaks are asymmetric, most evidently with larger retardation effects, and the number and energy of each Gaussian/Lorentzian is not always easy to determine.
Here we propose a new approach to fit the LSPRs based on the line shapes of plasmon modes excited by an electron passing close to a metal sphere. EELS for some impact parameter and sphere radius can be calculated analytically using a linear sum of basis functions and associated coefficients [2]⁠. Allowing the line-shape to scale in amplitude and shift in energy axis enables fitting such functions to simulated or experimental data.
As an example, we consider fitting EELS of a silver nanocube (with rounded corners and edges) in vacuo. Fig 1 is an e-DDA [3]⁠ simulated loss spectrum from a cube with 10 nm edge. The spectrum has been fitted using three separate line-shapes, each calculated from Mie theory for LSPRs. For spheres of this size, the dipolar mode is dominant. Fig 2 shows an equivalent spectrum to Fig 1, but for a silver cube of 100 nm size. The spectrum again has been fitted using this method, but with different radii for the three spheres. Fig 3 shows how the “equivalent sphere” diameter changes with cube size for the lowest order (dipolar) excitations. The smaller (10nm) cube can be fitted using the usual collection of Lorentzians, but it does not work well with the 100nm cube due to asymmetric peaks.
Work is ongoing to apply this approach in analysing experimental data and determine the limitations of it in terms of nanoparticle size and geometry. Whilst empirically the fits are very promising, we continue to work on the underlying basis for why the spectra can be described in terms of a series of spectra from “equivalent spheres” but note that, for example, in the case of cubes, cubic harmonic functions describing the lattice harmonics of a cubic crystal can be written in terms of spherical harmonics [4].
[1] J. Nelayah et al., Nat. Phys. 3, 348 (2007).
[2] F. J. Garcia de Abajo, Phys. Rev. B 59, (1999).
[3] N. W. Bigelow et al., ACS Nano 6, 7497 (2012).
[4] S. Altmann and A. Cracknell, Rev. Mod. Phys. 37, 19 (1965).


We acknowledge the support received from the European Union Seventh Framework Program under Grant Agreement 312483 – ESTEEM2 (Integrated Infrastructure Initiative – I3) and under Grant Agreement 291522-3DIMAGE. FDP and CD acknowledge funding from the ERC under grant number 259619 PHOTO EM

Fig. 1: Fitted EELS spectrum of 10nm silver cube with the electron crossing between opposite midpoints of edges 6 nm above the surface with a beam energy of 300 keV. Components from separate “effective spheres” have been highlighted.

Fig. 2: Fitted EELS spectrum of 100nm silver cube (same trajectory and beam energy as Fig 1). Components from separate “effective spheres” have been highlighted. Particle is big enough to see retardation effects: lowest energy component can only be approximated with at least two spherical modes and is now asymmetric.

Fig. 3: Diameter of “effective sphere” for lowest energy component versus the simulated rounded cube edge length. A straight line with gradient 1 and no offset is plotted to guide the eye.

Type of presentation: Poster

IT-5-P-3513 Fabrication of High Energy Resolution Silicon Drift Detector for Energy Dispersive X-ray Spectrometer

Hsu C. C.1, Tseng F. G.1, Chen F. R.1, Lee C. H.1, Chuang Y. J.2
1Department of Engineering and System Science, National Tsing Hua University, Hsinchu, Taiwan, 2Department of Biomedical Engineering, Ming Chuan University, Taipei, Taiwan
ful60406@hotmail.com

Energy Dispersive X-ray Spectrometer (EDS) is the most common analytical equipment in SEM or TEM used for the elemental analysis or chemical composition of a sample. The most critical part in EDS system is the x-ray detector. The most common detectors are made of Si(Li) crystal. The major drawback of Si(Li) detector is that they require hours to cool down before use, and cannot be allowed to warm up during use. Besides, the increase of Si(Li) detection area will increase the capacitance that cause the increase of electronic noise. There is a trend towards a newer EDS detector over past decade, called the silicon drift detector (SDD). The key advantage of the SDD is the very lower anode capacitance compared with conventional silicon detectors of the same area. This unique feature reduces electronic noise and shortens processing shaping time to achieve higher energy resolution and counting rate. Due to the small anode in the SDD the leakage current is so low that the SDD can be operated with moderate cooling by Peltier cooler.
High resolution SDD have been designed, fabricated and tested.(Fig.2) The SDDs were fabricated on n-type<111> and a resistivity of more than 4kΩ.cm silicon substrates with a thickness of 400μm. The SDD consists of fully depleted silicon, in which an electric field with a strong component parallel to the surface drives electrons generated by the absorption of x-ray towards a small sized collecting anode. The electric field is generated by a number of increasingly reverse biased cathodes on one side surface of the device. The radiation entrance window on the opposite side is made by a non-structured shallow implanted junction giving a homogeneous sensitivity over the whole detection area.
SDD comprise homogeneous radiation entrance window and the anode guard ring for improving the energy resolution of detector.(Fig.1) First, if the p-n junction is located at deep depth, the thickness of dead layer will to thick. Therefore the homogeneous window is important to control the loss of entrance X-ray. Detailed studies showed that a 1100-1200 Å thick layer of aluminum sufficiently attenuates visible light so that it has a negligible impact on the SDD leakage current. Second, in the detector all electrons generated at the Si/SiO2 interface are collected on the anode guard ring rather than contributing to the detector leakage current.
SDD were characterized to extract critical I-V performance parameter like total leakage current at anode. The value of leakage current of ring_2 is 248 nA without anode guard ring, and reduces to 4 nA with anode guard ring.(Fig.3) And will test the response for detector exposed to the X-ray source in the future. The goal of the testing have shown a FWHM at MnKα line of a radioactive 55Fe source of 170 eV at -20℃.


Fig. 1: Cross section and operation scheme of a SDD detector with anode guard rings and non-structured homogeneous x-ray entrance window.

Fig. 2: Photo image of a fabricated silicon drift detectors.

Fig. 3: Leakage current of SDD with testing anode guard ring.

Type of presentation: Poster

IT-5-P-5719 Major update of the EELS database: eelsdb.eu

Lajaunie L.1, Ewels P.2, Sikora T.3, Serin V.4
1Institut des Matériaux Jean Rouxel, (IMN) – Université de Nantes, CNRS, 2 rue de la Houssinère - BP 32229, 44322 Nantes Cedex 3, France, 2Babraham Institute, Cambridge, UK, 3SAVANTIC AB, Rosenlundsgatan 50, 118 63 Stockholm, Sweden, 4CEMES, Université Toulouse 29 rue Jeanne Marvig BP 94347 31055 Toulouse, France
luc.lajaunie@cnrs-imn.fr

Since its creation at the end of the 1990’s, the EELS database has gathered more than 200 spectra covering 35 elements of the periodic table, becoming the largest open-access electronic repository of spectra from Electron Energy-Loss Spectroscopy and X-ray Absorption Spectroscopy experiments.1 The EELS database is now a common tool used by spectroscopists, theoreticians, students and private firms as a reference catalog for fine structures and data-treatment analyzes2-4 and has been referenced by more than 30 papers.

Much of this success is due to the open-access nature of the database. The database depends on voluntary user contributions; to encourage these contributions, we have performed a major update of the website which is now accessible at http://eelsdb.eu/. The design of the website has been improved (Figure 1) and several new functions have been implemented, including a plotting function (Figure 2) which allows the online comparison of spectra before downloading. The new design gives greater emphasis on the original work of the contributors by strongly highlighting their papers. In addition of the database itself, users can post and manage job adverts and read the latest news and events regarding the EELS community. All these improvements will be discussed further in the poster details.

1. T. Sikora and V. Serin, EMC 2008 14th European Microscopy Congress, pp-439-440, Springer-Verlag Berlin (2008)
2. N. Bernier et al., Materials Characterization, 86, pp-116-126 (2013)
3. L. Zhang et al., Physical Review B, 81, 035102 (2010)
4. R. Núñez-González et al., Computational Materials Science, 49, pp-15-20 (2010)


The authors would like to thank the IMN and CEMES laboratories, the European microscopy network ESTEEM 2, the French microscopy network METSA and the French microscopy society Sfµ, for the funding. The authors warmly acknowledge everyone who has contributed to the database.

Fig. 1: Homepage of the EELS spectra database: http://eelsdb.eu/.

Fig. 2: The plotting page of the website allows the online comparison of spectra before downloading thanks to zoom-in and normalization functions.

Type of presentation: Poster

IT-5-P-5813 Spectrum-based phase mapping of apatite and zoned monazite grains using principal component analysis

Seddio S. M.1
1Thermo Fisher Scientific, Madison, WI, USA.
stephen.seddio@thermofisher.com

EDS X-ray mapping requires trade-offs between interaction volume, collecting enough above-background counts, selecting appropriate elements, and avoiding sample damage. These trade-offs may produce confusing results, especially in samples containing multiple phases with similar compositions. Applying contrast enhancements and filters to X-ray maps fails to eliminate the confusion of interfering X-ray lines and phases with similar compositions. However, acquiring an image cube with an EDS spectrum at every pixel and comparing the mapped spectra using principal component analysis (PCA), phases can be readily distinguished.
A rock sample containing accessory monazite ([La,Ce,Pr,Nd,Th]PO4) was polished, carbon coated, and examined in an FESEM. EDS spectral imaging was done at 5 and 15 kV. Phases were identified using COMPASSTM spectral phase mapping, which identifies phases based on PCA of the EDS spectrum at each pixel [1,2].
In Figs. 1 and 2, an apatite (Ca5[PO4]3[F,Cl,OH]) grain is partially included in a monazite grain. In 15 kV BSE imaging (Fig. 1a), transmission through thin phases (e.g., “Silica;” Fig. 1a) is evident. 5 kV imaging (Fig. 2a) produces images and X-ray maps more representative of the sample surface. In the 5 kV O and P Kα maps (e.g., Fig. 2c), apatite and monazite are indistinguishable (2.5 hour acquisition). If this sample was mapped without a light REE in the setup, monazite could be misidentified as apatite. However, after 7.5 minutes of acquisition time, COMPASS distinguishes the phases (Fig. 2c). Additionally, spectral imaging of the monazite grain in Fig. 1b (15 kV) reveals a partial rim, < 1 μm wide, that contains higher Th. PCA is able to distinguish the Th-rich rim at 5 kV as well (Fig. 2b).
PCA is an important tool for clarifying confusing X-ray maps. Using spectral imaging with PCA can provide higher confidence identification of phases in less time than traditional elemental mapping.

References
[1] Keenan et al., Method of Multivariate Spectral Analysis. Patent 6,675,106 B1. 06 Jan. 2004.
[2] Keenan et al., Apparatus and System for Multivariate Spectral Analysis. Patent 6,675,106 B1. 06 Jan. 2004.


Fig. 1: Fig. 3. EDS spectra from apatite (red), high-Th monazite (blue), and monazite (green). The vertical axis is a square root scale.

Type of presentation: Poster

IT-5-P-5816 Overcoming Quantitative Challenges Presented By X-Ray Line Interferences in EDS and WDS.

Seddio S. M.1
1Thermo Fisher Scientific
stephen.seddio@thermofisher.com

Quantitative analysis using EDS or WDS of phases containing elements with interfering X-ray lines presents challenges to the microanalyst. To illustrate some of these challenges, quantitative analysis of a two phase Ti-V-Al-Fe sample was done.
A sample of Ti-V-Al-Fe metal was examined in an FESEM. EDS and WDS data were collected using a Thermo Scientific™ UltraDry™ EDS detector and the Thermo Scientific™ MagnaRay™ WDS Spectrometer. EDS and WDS data were processed using the Thermo Scientific™ NORAN™ System 7. Quantitative analysis was done at 15 kV. EDS spectral imaging was done at 10 kV. Two phases (Figs. 1, 2) were identified using COMPASS™ spectral phase mapping, which identifies phases based on the principle component analysis of the EDS spectrum at each pixel [1].
 ~5 µm, V-rich (~13 wt% V) grains occur along the boundaries of larger, ~10 µm, V-poor (~3 wt% V) grains. Quantitative results are in Table 1.
Ti Kβ line is only separated from V Kα by 17 eV; these X-ray lines are indistinguishable by EDS and are poorly resolved by WDS (Fig. 3). The effect of this interference in the WDS quantitative analyses of these phases is the over-estimate of V.
There are three methods by which this shortcoming may be overcome. First, the V Kβ line is an appealing peak on which to count because there is no interfering energy line in this sample. However, greatly (>10×) extended acquisition times are required for counting the V Kβ line in the V-rich grains. In the V-poor grains, the V concentrations are low (~3 wt%) and the V Kβ line cannot be distinguished from the background. Second, a difference method can be utilized. This method subtracts the wt% of the other elements from 100% with the remainder representing the V concentration. This method requires that only one line is confounding and that the measurement of the remaining elements is done perfectly. The third method is to perform EDS quantitative analysis with standards. It is typically assumed that WDS is more accurate than EDS. However, EDS has peak deconvolution methodologies with both standards-based and standardless quantitative analysis, providing more accurate results than WDS in this case.
WDS is necessary technique for confirming the presence or absence of interfering elements, but unless the WDS spectrometer is able to completely resolve interfering X-ray lines, it cannot be used for accurate quantitative analysis. In addition, interfering energy lines confound the phase mapping. The peak deconvolution methods involved in modern EDS quantitative analysis provide accurate results when WDS is unable to do so. In addition, the utilization of EDS-based COMPASS discriminates phases with only subtle compositional differences.
References
[1] P. Camus, Thermo Scientific (2009) White Paper 51782.


Fig. 1: .

Type of presentation: Poster

IT-5-P-5821 Quantitative measurement of site-specific spin and orbital magnetic moments by electron energy-loss chiral magnetic dichroism

Xiaoyan Zhong 1, Ziqiang Wang 1, Jing Zhu 1, Rong Yu 1, Zhiying Cheng 1
Beijing National Center for Electron Microscopy, School of Materials Science and Engineering, The State Key Laboratory of New Ceramics and Fine Processing, Laboratory of Advanced Materials (MOE), Tsinghua University, Beijing 100084, People’s Republic of China 1
xyzhong@mail.tsinghua.edu.cn

In this work, we have developed the site-specific electron energy-loss magnetic chiral dichroism (EMCD) method for local magnetic moment determination in magnetic materials with non-equivalent crystallographic sites at a nanometer scale. It’s the first work to experimentally demonstrate that the fast electron as a new source can be used to determine magnetic structure, including quantitative magnetic moments, for a wide range of materials, which is generally considered to be accomplished by neutron diffraction. Compared with previous EMCD works in which EMCD was just used for detecting the ferromagnetic signals of materials, we fundamentally raise the EMCD technique to the new level of magnetic structure determination.

In the example of NiFe2O4, we achieve comprehensive magnetic structure information using the site-specific EMCD method under the assumption of no magnetic information known previously. The magnetic structure information we obtain includes site-specific total magnetic moment, site-specific orbital to spin magnetic moment (mL/mS) ratio and total magnetic moment of a unit cell. Our method is testified to be valid by comparing our results with those obtained by theoretical calculations and other experimental techniques such as X-ray magnetic circular dichroism and neutron diffraction. Using transmitted electron in site-specific EMCD method, we can reach a high spatial resolution, and get site-specific and element-specific magnetic information, as well as distinguish the orbital and spin magnetic moments.

In the technical aspects, the extremely strong EMCD signals have been achieved by using site-specific EMCD method, which allow us to do quantitative analysis. We first did the quantitative works on EMCD spectra to obtain total magnetic moments (the sum of spin and orbital magnetic moments) for atoms in different sites. For example, our work first reports the experimentally determined mL/mS ratios of Fe atoms in octahedral and tetrahedral sites.

In sum, our work opens the door of using fast electrons to determine magnetic structures for a wide range of magnetic materials in a nanometer scale. Site-specific EMCD may benefit much not only to the fundamental research of magnetic states and behavior in complex magnetic materials, but also to revealing the magnetic structure in nanostructures or interface of the composite magnetic films.


This work is financially supported by National 973 Project of China and Chinese National Nature Science Foundation. This work made use of the resources of the Beijing National Center for Electron Microscopy. The authors are grateful to Profs. Z.H. Zhang, J. Yuan, X.Q. Pan, D.S. Wang, Dr. L. Xie, Mr. Y. Xia, D.S. Song, Z.Y. Wang, Profs. S.P. Crane, R. Ramash, Q. Zhan and Dr. S. Löffler.

Fig. 1: Schematic image of site-specific electron energy-loss chiral magnetic dichroism

Type of presentation: Poster

IT-5-P-5868 Trends in X-ray Nano-Analysis by TEM/STEM

von Harrach H. S.1
1SvH Microscopy, East Sussex, UK
svhmicroscopy@gmail.com

The improvements of X-ray detectors gained momentum in recent years with the introduction of silicon drift detectors (SDD), and the increase in collection solid angle (SA) and speed of acquisition. The ability to produce larger detectors up to 100 mm2 without impairing the energy resolution means that detectors with 1 srad collection angle are now available (1). Since the SDDs are capable of higher throughput they have also improved the minimum detection limit of an element in a given analysis time.
Improvements in TEM/STEM resolution were brought about in recent years by the correction of spherical aberration (Cs) and the introduction of higher brightness sources, both Schottky and cold field-emission sources. This has resulted in sub-Angstrom image resolution for thin specimens, but only sub-nanometer probe sizes at the currents (~1nA) required for acquiring X-ray data with Si(Li) detectors at reasonable signal-to-noise ratios (S/N).
The combination of these developments has resulted in the ability to use lower probe currents to produce good quality analytical data at the atomic level within a few minutes (2). This is crucial for many materials that are damaged by exposure to high-energy electrons at high intensity.
So how should the technique be improved further in the future? On the TEM side, the need to operate at lower kV to minimise specimen damage calls for the use of higher brightness sources with lower energy spread (dE) or for chromatic aberration correction. Cold field-emitters are the brightest sources currently available (reduced brightness Br~1e8 A/m2/sr/V) and, either in combination with a monochromator or Cc corrector, could result in sub-Angstrom resolution at 40 -80kV (see Fig.1). For monochromated (dE=0.1eV) and Cs/Cc corrected instruments with cold field-emitters the probe size (d50) at 10pA is below 0.1nm at 40-50 kV and above.
On the detector side, there is room for improvement of collection angles approaching the theoretical limit of 4π steradians. This would reduce the dose required for good quality nano-analysis (S/N>5) by a factor of 10 or more. Fig.2 shows the relative dose at constant X-ray counts detected as a function of acceleration voltage (kV). As the kV is reduced the X-ray yield increases roughly as inverse square root of kV (3). Consequently, with a 10 sr detector the electron dose could be reduced by a factor of almost 20 by operating at 60 kV and sub-Angstrom probe size, compared to the current best instruments at 200kV with 1 sr detectors.
References
1. H.S. von Harrach, P. Dona, B. Freitag, H. Soltau, A. Niculae & M. Rohde (2009) Microsc.Microanal. 15 (Suppl.2), 208-9
2. A. J. D’Alfonso, B. Freitag, D. Klenov, and L. J. Allen (2010). Phys. Rev. B 81, 100101(R).
3. C.J. Powell (1976) Rev.Mod.Phys. 48, 33


The author was working at FEI Electron Optics B.V, The Netherlands until 2013.

Fig. 1: Fig.1 Probe size (d50) at 50% of constant 10pA probe current vs. acceleration voltage for cold FEG source (reduced brightness Br=1e8 A/m2/sr/V) with Cs corrector (Cs<2um), mono-chromated CFEG with dE=0.1eV, Br=3.3e7 and Cc corrected CFEG system with Br=1e8, Cc<0.1mm. [ref. P.Kruit et al. 2006 J. Appl. Phys. 99, 024315]

Fig. 2: Fig.2 Relative electron dose vs acceleration voltage at constant number of X-rays detected for X-ray detectors of collection solid angles SA = 1 and 10 sr.

Type of presentation: Poster

IT-5-P-5871 Single atom Electron Energy Loss Spectroscopy at Low Primary Electron Energy in the Electron Microscope

G Tizei L. H.1, Iizumi Y.1, Okazaki T.1, Nakanishi R.2, Kitaura R.2, Shinohara H.2, Suenaga K.1
1Nanotube Research Center, National Institute of Advanced Industrial Science and Technology (AIST), Tsukuba 305-8565, Japan, 2Department of Chemistry, Nagoya University, Nagoya 468-8602, Japan
Luiz.tizei@aist.go.jp

In this poster we discuss single atom electron energy loss spectroscopy (EELS) of lanthanides (La, Ce, Er and Eu). In particular we analyze the different possibilities of high spatial resolution spectroscopy at low primary electron energy (30 keV and 60 keV) and high electron energy losses (above 800 eV). Atomically resolved EELS experiments were performed in the JEOL-CREST double corrected microscope operated at 30 keV and 60 keV. The samples analyzed were lanthanide atoms (La, Ce and Er) encaged in fullerenes stored in carbon nanotubes and Eu atomic chains inside carbon nanotubes.
The use of low primary electron energies is beneficial due to the decrease in energy loss delocalization and the minimization of one possible sample damage mechanism (knock-on). Examples of the experiments performed at 30 keV and 60 keV are shown in Figure 1. Evidently, one can see that the fullerenes structure is maintained in the experiment at 30 keV, while it is modified at 60 keV. All experiments have been performed with acquisition times ranging from 20 ms to 200 ms.
To observe the effect of energy loss on delocalization, we have performed the parallel acquisition of the loss signal of the N45 and M45 edges of La and Ce at 60 keV. These edges can be observed at 117 eV (121 eV) and 832 eV (881 eV) for La (Ce). To estimate delocalization, we have measured the full width at 50% intensity (L) of profiles for the annular dark field image (ADF), N45 and M45 edges. For a single La atom we have observed LADF = (0.15±0.02) nm, LN = (0.32±0.02) nm and L = (0.20±0.02) nm. The significant decrease is expected. However, the absolute values are smaller than those predicted by simple considerations.
Some effects of atomic movement on spectroscopy experiments will also be discussed.


This work is partially supported by a JST Research Acceleration programme.

Fig. 1: 2D EELS maps identifying the position of La (blue), Ce (green) and Er (red) at 60 kV (a-c) and 30 kV (d-f). The maps presented are not the first on a series of acquisitions. For this reason, electron beam damage can be observed at 60 kV and not at 30 kV.

Type of presentation: Poster

IT-5-P-5918 Identifying suboxide grains at the metal-oxide interface of a corroded Zr-1.0Nb alloy using (S)TEM, transmission-EBSD and EELS

Hu J.1, Garner A.2, Ni N.3, Gholinia A.2, Nicholls R.1, Lozano-Perez S.1, Frankel P.1, Preuss M.2, Grovenor C.1
1Department of Materials, Oxford University, Parks Road, Oxford, UK, 2Materials Performance Centre, School of Materials, University of Manchester, Manchester, UK, 3Department of Materials, Imperial College London , Royal School of Mines, London, UK
jing.hu@materials.ox.ac.uk

Here we report a methodology combining Transmission Electron Microscopy (TEM), Scanning Transmission Electron Microscopy (STEM), Transmission-EBSD (t-EBSD) and Electron Energy Loss Spectroscopy (EELS) to analyse the structural and chemical properties of the metal-oxide interface of corroded Zr-1.0Nb alloys in unprecedented detail. The sample which has been under autoclave condition for 360 days shows no sign of transition, suggesting its excellent corrosion resistance1. TEM and STEM results reveal the complexity of the distribution of suboxide grains at the metal-oxide interface. Convergent beam electron diffraction (CBED) patterns were acquired from a region close to the metal/oxide interface which matches with the [3 2 -4] zone axis of the hexagonal ZrO phase with P-62m symmetry and lattice parameters a=5.31 Å and c=3.20 Å predicted by Nicholls et al2. EELS provided accurate quantitative analysis of the oxygen concentration across the interface, identifying the existence of local regions of stoichiometric ZrO and Zr3O2, with significant local variations in thicknesses from 20 nm to 326nm, much thicker than observed previously in other oxidised zirconium alloys3. T-EBSD confirmed that the suboxide grains can be indexed with the hexagonal ZrO structure. The t-EBSD analysis has also allowed us to map a relatively large region (~7μm) of the metal-oxide interface, revealing the location and size distribution of the suboxide grains. These observations will be compared to previous reports of less corrosion-resistant alloys studied by the same techniques.
1. Wei, J. et al. Autoclave study of zirconium alloys with and without hydride rim. Corros. Eng. Sci. Technol. 47, 516–528 (2012).
2. RJ Nicholls, N Ni, S Lozano-Perez, A London, DW McComb, PD Nellist, CRM Grovenor, CJ Pickard, J. Y. Crystal structure of the ZrO phase at zirconium / zirconium oxide interfaces. Adv. Eng. Mater.Accepted
3. Ni, N. et al. How the crystallography and nanoscale chemistry of the metal/oxide interface develops during the aqueous oxidation of zirconium cladding alloys. Acta Mater. 60, 7132–7149 (2012).



This research was funded by the MUZIC2 consortium and JH is supported by the China Scholarship Council.

Fig. 1: (a): Bright field image of the TEM sample. Figure 1(b): HAADF-STEM image of the TEM sample. The area used for EELS analysis is highlighted. The inset shows higher magnification images of this same area after further FIB thinning (which has created the hole under the left hand crack)

Fig. 2: (a) High magnification bright field image the metal-oxide interface chosen for EELS analysis. The suboxide grain which was diffracting is highlighted using arrows. (b) STEM dark field image of the suboxide grain area, the area where CBED pattern was taken is highlighted. (c) CBED pattern of the suboxide grain which matched with hexagonal ZrO phase.

Fig. 3: (a) Band contrast map, (b) Phase map from t-EBSD analysis of the TEM sample. The yellow part near the metal-oxide interface is matched with hexagonal ZrO phase with P-62m symmetry and lattice parameters a=5.31 Å and c=3.20 Å2. The area for the EELS and CBED analysis is also highlighted.

Fig. 4: Positions of EELS line scans from the suboxide region relative to (a) the t-EBSD map. (b) the HADDF image. Lines started from the oxide towards metal. (c) Zirconium and oxygen concentration of the EELS line scan showing both stoichiometric ZrO and Zr3O2, with significant local variations in thicknesses from 20 nm to 326nm.

Type of presentation: Poster

IT-5-P-5962 Quantitative electron probe microanalysis of Ga-doped BiFeO3 and (Ca,Zr)-doped BaTiO3 thin films

LONGUET J. L.1, JABER N.2, DAUMONT C.2, WOLFMANN J.2, NEGULESCU B.2, RUYTER A.2, FEUILLARD G.2, BAVENCOFFE M.2, FORTINEAU J.2, SAUVAGE T.3, COURTOIS B.3, BOUYANIFIF H.4, AUTRET-LAMBERT C.2, GERVAIS F.2
1CEA, DAM, Le Ripault, BP16, F-37260 Monts, France, 2Laboratoire GREMAN, UMR7347 CNRS Université François Rabelais, faculté de sciences et techniques 37200 Tours, France, 3Laboratoire CEMHTI, UPR3079 CNRS, Site Cyclotron 45071 Orléans cedex 2, France, 4Laboratoire LPMC, Université Jules Vernes Picardie - Amiens, France
jean-louis.longuet@cea.fr

The most efficient multifonctional piezoelectric materials are lead-based, sush as Lead Zirconate Titanate LZT for example. Considering the need to use lead-free materials, an interest in bismuth ferrite (BiFeO3, also commonly referred to as BFO in Materials Science) has grown due to its ferroelectricity. The substitution of Bismuth in BiFeO3 modifies its properties and enhances the piezoelectric activity. Doped BaTiO3 perovskite has also great interest in this domain.

Linear lateral composition gradients were created by combinatorial Pulse Laser Deposition (PLD) using targets of doping material and host perovskite : Gallium is the doping element in the BGFO material (BiFeO3 + GaFeO3). Calcium and Zirconium are the doping elements in the BCTZ material (BaTiO3 + (Ba,Ca)(Ti,Zr)O3).

Here, we report quantitative thin films analysis (TFA) performed by wavelength dispersive spectrometry (WDX) with an electron microprobe (EPMA). Experimental procedure is described to achieve acceptable quantitative results regarding some analytical issues such as small doping content (less than 3 %wt), small top layer thickness (less than 100 µg/cm2) and lines interferences occurring in energy dispersive spectrometry (EDS) but not in WDS analysis (like Ba Lα – Ti Kα).

Mass thickness of BGFO, determined by TFA-EPMA, was confirmed by focus ion beam (FIB) cross section imaging. A short link is made with optimum doping concentration found on BGFO according to local piezoresponse measurements using an atomic force microscope (AFM) in piezoelectric mode (PFM) ) and by dual beam laser interferometry.


Fig. 1: BGFO sample description (10x10 mm²)BGFO (~100nm)/LSM(~40nm)/STO(substrate)

Fig. 2: Ga-doping content along X-Line profile n°2 (round circles are coincidence points with Y-line profiles n°3-4-5-6)

Fig. 3: BGFO mass thickness along X-Line profile n°2 (round circles are coincidence points with Y-line profiles n°3-4-5-6)

Fig. 4: FIB Cross Section

Type of presentation: Poster

IT-5-P-5964 On application of the Multivariate Statistical Analysis in spectrum-imaging

Potapov P.1
1temDM, Dresden, Germany
info@temdm.com

Availability of TEM instruments with fast spectrum-imaging EDX and EELS facilities made it possible to map the composition and structural properties with a high resolution. A typical spectrum-imaging data cube now routinely exceeds the size of 100x100x1000 pixels. The extraction of the chemical/structural information from such huge arrays of data can be significantly improved by using well established techniques of Multivariate Statistical Analysis.

Among the multivariate statistical methods, the most attention is paid to the Principal Component Analysis (PCA) which decomposes the observation set into the set of linearly uncorrelated variables. The components with the highest variance are assumed to have the highest significance and to correlate with the variation of the material parameters such as composition or structure features while the lower-variance components might be associated with the statistical noise and therefore ignored. As PCA is closely related to the eigenvector decomposition in linear algebra, the very efficient algorithms for its implementation are available. However, a caution should be taken when treating a spectrum-image from a system of several objects, for instance, an agglomeration of the particles of different nature. Fig.1 shows the score plot of the PCA components for such a system clearly indicating the separation of the data onto the two distinct clusters. In this situation, the PCA results represent the average eigenvectors for the two unrelated data sets and cannot bear any physical meaning. A more efficient strategy is to segment first the data onto the appropriated clusters and then apply PCA for each cluster individually.

Another approach is the reconstruction of a spectrum-image using a small number of the highest-variance components while cutting off the rest “noise” components [1]. Here PCA is used as a kind of noise filter and the physical meaning of the PCA components is unimportant. The problem appears when the variance of the minor components is comparable or beneath the typical variance due to noise. In this case the useful signal might partially “leak” to the “noise” components and be lost during the subsequent reconstruction [2]. The possible solution is to retrieve a relatively large number of the PCA components and then apply to them the Independent Component Analysis (ICA). Similar to PCA, ICA can be thought of as a rotation in the variable coordinates that maximizes the Curtosis of a given component. This way, the truly independent not just statistically uncorrelated components can be retrieved while the noise can be cut off.

References:

[1] M. Watanabe, E. Okunishi, K. Ishizuka, Microscopy and Analysis 23 (2009) 5-7.

[2] S. Lichtert, J. Verbeeck, Ultramicroscopy, 125 (2013) 35-42.


Fig. 1: Scatterplots of the first three PCA components for the system composed of two objects.PCA retrieves the “average” eigenvectors that cannot unmask the nature of theobjects. The fragmentation of the data and the application of PCA to eachcluster individually make the maximal variance in each object coinciding withthe direction of the eigenvectors.

Type of presentation: Poster

IT-5-P-5988 Distributions of cations and inversion parameter in nonstoichiometric magnesium aluminate spinel characterized by electron energy-loss spectroscopy

Halabi M.1,2, Ezersky V.2, Kohn A.1,2, Hayun S.1,2
1Department of Materials Engineering, Ben-Gurion University of the Negev, 2Ilse Katz Institute for Nanoscale Science and Technology, Ben-Gurion University of the Negev
akohn@bgu.ac.il

The effect of composition and heat treatments on the distribution of cations and on the inversion parameter in magnesium aluminate spinel was studied using Electron Energy Loss Spectroscopy (EELS). Powders of MgO•nAl2O3 (0.95<n<1.07) with nano-sized grains were synthesized by solution combustion and heat treated using Spark Plasma Sintering (SPS) and Pressure-less Sintering (PS) methods. EEL spectra were collected at varying distances perpendicular to grain boundaries from which the Mg to Al cation ratio and the inversion parameter, which is the fraction of tetrahedral sites occupied by Al cations, were calculated. The Mg to Al cation ratio was calculated from their core loss K-edges. To estimate the fraction of tetrahedral sites occupied by Al cations, the inversion parameter was calculated from the Al L-edge using the methodology suggested by Bruley et al., [1] which is based on the integral ratio between L3 to L2 peaks.
We report that cations which segregate to the grain boundaries are the excess component relative to the stoichiometric composition of the spinel (Fig. 1). Heat treatment does not change the type of segregate cation but does affect the degree of order of spinel. We find that spinel powders with nano-sized grains subjected to SPS treatment results in higher order compared to PS, namely lower inversion parameter. Finally, we discuss the experimental requirements for measuring reliable EEL spectra from these materials which are sensitive to damage from the electron beam.

[1] J. Bruley, M.-W. Tseng and D. B. Williams "Spectrum-Line Profile Analysis of a Magnesium Aluminate Spinel Sapphire Interface," Microsc. Microanal. Microstruct, pp. 1-18, 1995.


Fig. 1: (a) Schematic representation of the collection of EEL spectra perpendicular to grain boundaries. (b) Mg to Al cation ratio as a function of relative position perpendicular to the grain powders of spinel samples heat treated by SPS. (SCZ represents space charge zone)

Type of presentation: Poster

IT-5-P-6028 Addressing challenges in Electron Energy Loss Spectroscopy on individual atoms

March K.1, Brun N.1, Gloter A.1, Tencé M.1, Mory C.1, Stéphan O.1, Colliex C.1
1Laboratoire de Physique des Solides Université Paris-Sud - CNRS UMR-8502, Orsay, France
katia.march@u-psud.fr

The latest generation of STEM microscopes based on many instrumental developments (Cs corrector, lower primary voltages, EELS and EDX detector improvements...) offers the ability to track in a spectrum-image mode several signals generated simultaneously by individual atoms [1,2,3] and to rekindle the STEM-EELS spectro-microscopy of single atoms.

A Nion UltraSTEM microscope equipped with a Cs corrector and with a home-made fast EELS detector has been used to record a few typical cases illustrating the present situation in individual atom spectroscopy. With the new spectroscopic hardware, we can acquire EELS spectrum images of typically 100x100 pixels and covering a range of 1600 channels at an acquisition rate of 2300 spectra/s. Furthermore, the representation, exploitation and analysis of such data require some specific algorithms. The most widely used technique is the Principal Component Analysis (PCA) [4][5] and, as a filtering technique, offers an improvement of the signal to noise ratio. However, for high noise levels, a bias is introduced by PCA as signal bearing components are discarded with the removal of components considered as noise [6]. We have tested some algorithms based on non-local methods for denoising by exploiting the natural redundancy of patterns inside an image.

The first case is the determination of the position of Sm interstitial/substitutional dopants in ceria nanoparticles together with their valence changes in accordance with the variation of the ferromagnetic properties measured as a function of the nominal doping level [7]. The spectrum image has a high noise level and Sm doping could not be identified with usual PCA denoising. We have therefore tested Non-Local Sparse PCA [8] which produces interesting results: the filtered spectra display fine structures of edges and both spatial and spectral resolutions are preserved. The second example addresses the challenge of identifying the characteristic EELS signals from heavy (Tb, Th) atoms in rapid motion on a thin carbon layer which imposes a compromise between time acquisition and detection limit (see Figure).

This contribution emphasizes the possibilities currently offered by a tiny electron probe, a suitable efficient detector strategy and a well chosen signal analysis tool for single atom spectroscopy.

[1] K. Suenaga et al. Nature Chemistry 1 (2009) 415.

[2] O.L. Krivanek et al. Nature 464 (2010) 571.

[3] C. Colliex et al. Ultramicroscopy 121 (2012) 80.

[4] N. Bonnet et al. Ultramicroscopy 77 (1999) 97.

[5] F. de la Peña et al. Ultramicroscopy 111, 2 (2011) 169.

[6] S. Lichtert, J. Verbeeck, Ultramicroscopy 125 (2013) 35.

[7] S.-Y. Chen et al. Phys.Chem.Chem.Phys. 16 (2014) 3274.

[8] J. Salmon et al. J.Mathematical Imaging and Vision 48 (2014), 279.


Thanks are due to E. Delain et S. Baconnais (IGR, Villejuif, France) for tricky specimen preparation.

Fig. 1: Imaging and spectroscopy of Th and Tb atoms in rapid motion on a thin carbon foil under the electron beam (60 kV). HAADF images at 2 µs per pixel (a) and at 100 µs per pixel (b). Raw EELS spectra extracted from the SI at two different positions (blue and red) – acquisition time: 100 µs per spectrum (c).

IT-6. Environmental electron microscopy

Type of presentation: Invited

IT-6-IN-1652 Advances in Atomic Resolution-Environmental (Scanning) Transmission Electron Microscopy

Gai P. L.1,2, Yoshida K.3, Lari L.1, Ward M. R.1, Martin T.1, Boyes E. D.1,4
1The York Nanocentre and Department of Physics, University of York, UK, 2Department of Chemistry, University of York , York, UK , 3Institute for Advanced Research, Nagoya University, Japan and York Nanocentre, University of York, 4Department of Electronics, University of York, UK
pratibha.gai@york.ac.uk

Dynamic chemical reactions catalysed by solid surfaces in heterogeneous catalysis play a major role in the development of energy sources, healthcare, environmental controls and industrial chemicals. Visualisation of the evolution of structural changes in catalysts in their working state under controlled gas and temperature conditions at the atomic level in real time is crucial in the development of efficient catalysts and processes but is extremely challenging.

Previously we reported the development of the first atomic resolution-Environmental transmission electron microscope (atomic resolution-ETEM) for in-situ studies of dynamic gas-solid reactions at operating temperatures under controlled conditions at the atomic level [1,2]. Highlights of this development include a novel ETEM design with the objective lens polepiece incorporating radial holes for differential pumping and the regular EM sample chamber as the controlled reaction environmental cell (reactor) [2]. This atomic resolution-ETEM development is now widely used.

Recently we have developed a double aberration corrected E(S)TEM (AC E(S)TEM) at York incorporating a large gap objective lens polepiece for in-situ studies under controlled gas and temperature reaction environments with single atom sensitivity, using a JEOL 2200 FS [3,4,5]. The new E(S)TEM capability enables the visualisation of single atom dynamics in real time (Fig. 1) [4,5]. Here we present E(S)TEM studies of working catalysts at the single atom level. Supported Au nanoparticles are of interest in hydrogenation, water-gas-shift and low temperature oxidation of carbon monoxide. Supported Pt nanocatalysts are used in fuel cells where reactions in hydrogen, oxygen, CO and water are important and in vehicle exhaust emission control measures [4-6]. The E(S)TEM with single atom sensitivity is playing a key role in the development of a heterogeneous process for sustainable biofuels from biomass (including vegetable plants, weeds and grass) and environmentally benign heterogeneous processes to produce medicines for human healthcare.

References
[1] Gai P.L. et al, Science 267 (1995) 661.
[2] Boyes E.D. and Gai P.L., Ultramicr 67 (1997) 219.
[3] Gai P.L. and Boyes E .D., Micros Res Tech. 72 (2009) 153.
[4] Boyes E.D., Ward M., Lari L. and Gai P.L., Ann. Phys. (Berlin) 525 (2013) 423.
[5] Gai P. L., Lari L., Ward M. and Boyes E. D., Chem.Phys.Lett. 592 (2014) 355.
[6] Yoshida K. et al, Nanotechnology 2014 (submitted).  


 We thank the EPSRC (UK) for Critical Mass grant EP/J018058/1.

Fig. 1: Top: single atoms and raft-like clusters of platinum on carbon using our in-house development of E(S)TEM at York. Single atoms are single white dots in the image. (Scale bar=2nm); bottom: intensity profile of a single atom in (a).

Type of presentation: Invited

IT-6-IN-2314 Atomic structure and reactivity in catalysis studied by electron microscopy

Helveg S.1
1Haldor Topsøe A/S, Nymøllevej 55, DK-2800 Kgs. Lyngby, Denmark
sth@topsoe.dk

Developing efficient technologies for the production of fuel and chemicals as well as for reducing environmental harmful emissions are among the largest challenges for our modern society. As their solutions depend on catalysis, research and innovation in this field is mandatory to realize the vision of a clean and sustainable society. In recent years, new opportunities for catalysis research have opened up with remarkable progress in transmission electron microscopy (TEM). On one hand, advancements in aberration-corrected electron optics and data acquisition schemes enable TEM delivering images of catalysts with sub-angstrom resolution and single-atom sensitivity [1,2]. On the other hand, parallel developments of differentially pumped electron microscopes and of gas cells enable time-resolved observations of catalysts in situ during the exposure to reactive gas environments at pressures of up to the one-atmosphere level and temperatures of up to several hundred centigrade [3-5]. In this contribution, I will outline how such instrumentation and methodologies can advance in situ studies of surface structures and reactivity in catalysis. Specifically, the concept of using low electron dose-rates in TEM, in conjunction with in-line holography and aberration-correction, is introduced to allow maintaining atomic resolution and sensitivity during non-invasive in situ observations of catalysts [3,6]. Moreover, a novel nanoreactor concept is demonstrated for directly correlating time-resolved, high-resolution TEM images of catalysts with concurrent measurements of their catalytic functionality under reaction conditions at the ambient pressure level [4-5,7]. These competences expand the applicability of TEM in catalysis and build a foundation for “live” observations of structure-sensitive functional behavior at the single–atom level and in catalytically meaningful environments. Extraordinary benefits are illustrated by in situ studies in e.g. water splitting, hydrotreating and automotive emission abatement catalysis [1-10].

References

[1] C.F. Kisielowski et al, Angew. Chemie. Int. Ed. 49, 2708 (2010)

[2] L.P. Hansen et al, Angew. Chem. Int. Ed. 50, 10153 (2011)

[3] J.R. Jinschek, S. Helveg, Micron 43, 1156 (2012)

[4] J.F. Creemer et al, Ultramicroscopy 108, 993 (2008)

[5] S.B. Vendelbo et al, Ultramicroscopy 133, 72 (2013)

[6] S.Helveg, C.F. Kisielowski, J.R. Jinschek, P. Specht, G. Yuan, H. Frei (2014)

[7] S.B. Vendelbo, C.F. Elkjær, H. Falsig, I. Puspitasari, P. Dona, L. Mele, B. Morana, B.J. Nelissen, R. van Rijn, J.F. Creemer, P.J. Kooyman, S. Helveg (2014)

[8] Z. Peng et al, J. Catal. 286, 22 (2012)

[9] S.B. Simonsen et al, J. Am. Chem. Soc. 132, 7968 (2010); J. Catal. 281, 147 (2011)

[10] L.P. Hansen, M. Brorson, E. Johnson, S. Helveg (2014)


Type of presentation: Oral

IT-6-O-1684 in situ Nanoscale Hyperspectral XEDS Elemental Mapping in Liquids

Lewis E. A.1, Haigh S. J.1, Kulzick M. A.2, Burke M. G.1, Zaluzec N. J.1,3
1Materials Performance Centre and Electron Microscopy Centre, School of Materials, University of Manchester, Manchester, U.K., 2BP Corporate Research Center, Naperville, Illinois, USA., 3Electron Microscopy Center, Argonne National Laboratory, Argonne, IL 60439 USA.
edward.lewis@postgrad.manchester.ac.uk

Recent years have seen an explosion of interest in in situ (scanning) transmission electron microscope (S/TEM) studies of solution-phase processes. Work employing silicon nitride windowed environmental cells (e-cells) to study nanostructures in liquid has yielded important insights into mechanisms of nanoparticle growth.[1,2] One of the great strengths of the S/TEM platform is the potential to combine high resolution imaging with local analytical information obtained using electron energy loss spectroscopy (EELS) and X-ray energy dispersive spectroscopy (XEDS). However, both EELS and XEDS face challenges when applied to specimens in liquid e-cells and elemental mapping has proved impossible until now.[3,4] In this work we show that by rational redesign of an e-cell holder we are able to dramatically increase the collection efficiency of characteristic X-rays,[5] in order to achieve elemental mapping of nanostructures in liquid. Improved X-ray detection is obtained using an analytical XEDS version of the Protochips Poseidon 200 holder in a FEI Titan ChemiSTEM operated at 200 kV. Wet specimens were encapsulated between a pair of 50 nm thick SiNx windows in a Si e-cell with 150 nm spacers separating the windows.
As a proof of principle we have studied a sample consisting of a mixture of pre-synthesised nanostructures immersed in a Cu containing aqueous solution. Beam-induced interactions with the solution result in dynamic Cu nanoparticle growth processes (Fig. 1).[2] Simultaneous XEDS spectrum imaging of nanostructures in liquid facilitates interpretation of the dynamic processes occurring in this complex multicomponent system (Fig. 2). A beam-induced copper plating reaction occurs in the liquid-phase, (fig 2a and 2b) while similar growth is not seen in dry reference samples. The resulting spectrum image (Fig. 2c) reveals that Cu ions from the surrounding liquid are plating the pre-synthesised silver nanowires and gold nanoparticles producing bimetallic, core-shell, structures.
We have shown that it is possible to use XEDS to simultaneously map multiple elements in liquid with a spatial resolution approaching 10 nm. This new technique allows direct observation of nanoscale changes in composition and elemental distribution during solution-phase processes and has great potential in the field of nanoscience, to provide insights to aid the synthesis of mixed metallic nanostructures, as well for corrosion and biological studies.

1. Zheng H et al 2009 Science 324 1309-1312.
2. Liao H et al 2013 Chem. Commun. 49 11720-11727.
3. Jungjohann K L et al 2012 Microsc. Microanal. 18 621-627
4. Holtz M E et al 2013 Microsc. Microanal. 19 1027-1035
5. Zaluzec N J et al 2014 Microsc. Microanal. (in press) 20


This work was supported by multiple grants including: EPSRC Grants # EP/G035954/1 and EP/J021172/1, DTRA grant HDTRA1-12-1-0013, and the BP 2013 DRL Innovation Fund.

Fig. 1: Selected HAADF STEM images from a video sequence showing the beam-induced growth of Cu nanoparticles from an aqueous solution containing Cu ions. Images taken at time = 0s (a), 6.3s (b), 13.1s (c), 19.9s (d), 26.2s (e), and 31.4s (f). Scale bar = 80 nm.

Fig. 2: Beam-induced growth of Cu nanostructures occurs during extended spectrum imaging, as observed by comparison of the HAADF images before and after (a and b). XEDS data facilitates simultaneous mapping of multiple elements in liquid (c) with a spatial resolution of the order of 10 nm.

Type of presentation: Oral

IT-6-O-2397 Raman Spectroscopy coupled with environmental scanning transmission electron microscope

Picher M.1,2, Lin P. A.1,2, Blakenship S.1, Winterstein J.1, Sharma R.1
11Center for Nanoscale Science and Technology, National Institute of Standards and Technology, Gaithersburg, MD 20899-6203, 22Institute for Research in Electronics and Applied Physics, University of Maryland, College park, MD 20740
renu.sharma@nist.gov

In recent years the environmental transmission scanning electron microscope (ESTEM), has been successfully employed to reveal and understand the structural and chemical changes occurring in the nanoparticles under reactive environments [1,2]. The lack of statistical information available from TEM measurements is generally balanced by using other, ensemble measurement techniques such as x-ray or neutron diffraction, x-ray photoelectron spectroscopy, infrared spectroscopy, Raman spectroscopy etc.  However, it is almost impossible to create identical experimental conditions in two separate instruments to make measurements that can be directly compared. Moreover, ambiguities in ESTEM studies may arise from the unknown effects of the incident electron beam and uncertainty of the sample temperature.  Here, we present a unique platform that allows us to concurrently measure atomic-scale and micro-scale changes occurring in samples subjected to same reactive environmental conditions by incorporating a Raman Spectrometer on the ESTEM. 

We use a parabolic mirror, attached at the end of a hollow rod that can be inserted between the sample holder and the lower pole piece of the microscope (Fig. 1-2a). The mirror focuses the incoming laser on the sample and collects the scattered Raman photons. A set of optics then carries the Raman signal up to the spectrometer. Fig. 2.b,c show the Raman D and G band as well as the radial breathing modes of single walled carbon nanotubes (SWCNT) formed in the ETEM and an atomic-resolution still image extracted from a video sequence, respectively. We can monitor the growth rates using the G-band intensity under different growth conditions (Fig. 2d). This versatile optical setup can also be used i) to measure the temperature using Raman shifts, ii) to investigate light/matter interactions iii) as a heating source, iii) for general spectroscopy such as cathodoluminescence. Details of the design, function, and capabilities will be illustrated with results obtained from experiments on the in situ synthesis of carbon nanotubes.

Reference:

[1] Sharma, R., J. Mat. Res. 2005, 20, 1695

[2] Hansen et al., Science 2001, 294, 1508


Fig. 1: Schematic representation of the Raman data collection system: the laser passes through the hollow parabola holder, and is then focused on the sample by the parabola. The parabola collects the Raman signal and directs is back to the spectrometer.

Fig. 2: (a) Location of the parabolic mirror that collects the Raman signal, is located between the sample holder and the lower pole piece. (b) Raman spectrum collected from SWCNTs grown in the ESTEM. (c) Atomic Resolution image showing as grown SWCNTs (d) Time resolved evolution of the G band intensity (SWNT growth rate) at 650 °C under two C2H2 pressure.

Type of presentation: Oral

IT-6-O-2933 Operando TEM of CO Oxidation Catalyst by Quantification of Gaseous Reaction Products

Miller B. K.1, Crozier P. A.1
1Arizona State University, Tempe AZ, USA
benmiller002@aol.com

In-situ transmission electron microscopy allows materials to be observed at the atomic scale while they are simultaneously subjected to stimuli relevant to some application. Operando TEM goes one step further, and additionally measures some performance metric of the material during the in-situ observation. We have developed a technique for performing operando TEM of a catalyst for CO oxidation, which allows us to quantitatively monitor the gas composition leaving the environmental cell in an FEI Tecnai F20 ETEM [1]. We have done this by two complimentary methods [2]. Electron energy loss spectroscopy (EELS) was used to quantitatively probe the gas composition directly in the sample chamber of the microscope at discreet times in the course of an experiment. Mass spectrometry was simultaneously used to measure the gas composition continuously via a residual gas analyzer attached to the vacuum system near the main turbo-pump which pumps the environmental cell. High resolution images can then be linked to the precise conditions in the cell at all times as seen in Figure 1. All of this was made possible by the introduction of a novel sample preparation technique in which a pellet with a 0.5mm hole in the center was formed from glass-wool fibers, and impregnated with a silica-sphere supported catalyst. This pellet was then placed into a Gatan heating holder along with a metallic grid, which was also covered in the silica-sphere supported catalyst. This approach is currently being adapted to an aberration corrected FEI Titan ETEM.
Some of the initial results of this technique are shown in Figure 3. High resolution images of individual particles are placed into context within the operando experiment using a plot of the mol fraction of CO2 determined using EELS. A hysteresis of the CO conversion is clearly seen as the temperature is ramped up and then back down over a period of several hours. The EELS data which was acquired in the core loss region of the spectrum was quantified using a linear combination method to fit the carbon k edge π* peaks from CO and CO2, as shown in Figure 2. Changes in the structure of the Ru-RuO2 catalyst are clearly seen, including transitions from an oxidized to a reduced state at temperatures above 400°C.

References:

[1] Chenna, S. and Crozier, P. A. ACS Catalysis 2, 2395-2402. (2012).
[2] Miller, B. K. and Crozier, P. A. Microscopy and Microanalysis, 2014 (in press)


Financial support from National Science Foundation CBET-1134464 and the Fulton Schools of Engineering at ASU, and the use of ETEM at John M. Cowley Center for HR Microscopy at Arizona State is gratefully acknowledged.

Fig. 1: Operando data from CO oxidation experiment. a) Mass spectrometry data from residual gas analyzer (RGA), showing a sudden increase in CO2 as the temperature is increased to 230°C. b) Core loss EELS data with peaks from both CO and CO2. c) High resolution image acquired at the same condition as the EELS data.

Fig. 2: Method for quantifying the CO2 conversion from core loss EELS data by a linear combination of reference spectra. After background subtraction of both the reference and mixture measured spectra, a least squares fitting method is implemented in MATLAB allowing precise determination of the ratio of CO to CO2 in the gas phase.

Fig. 3: Mole fraction of CO2 as a function of temperature determined by analysis of core-loss EELS spectra. The data show a clear hysteresis as the temperature is increased and then decreased. This is attributed, in part, to the reduction of the initially oxidized Ru. Images at three temperatures illustrate the changes in the catalyst structure.

Type of presentation: Oral

IT-6-O-3407 Correlative Electron Microscopy and Photon Science Characterization of Working Catalysts

Stach E. A.1, Li Y.2, Zhao S.3, Zakharov D.1, Nuzzo R.3, Frenkel A.2
1Center for Functional Nanomaterials, Brookhaven National Laboratory, Upton, NY 11733 , 2Department of Physics, Yeshiva University, New York, NY 10016, 3Department of Chemistry, University of Illinois, Urbana-Champaign, Champaign, IL, 61801
estach@bnl.gov

Characterization of catalytic reactions is often hindered by the fact that the behavior the system is mesoscopic, while the materials involved are nanoscale, with features that can span a broad range of temporal and spatial scales and which involve a broad range of competitive interactions. As a result, the description of a catalytic system requires interrogation with a variety of techniques – involving imaging, diffraction and spectroscopy – to describe the dynamic changes in structure that can occur during reactions. Commonly, this is done by simple use of standard techniques, and inference of how the results relate to the working condition of the system. It is, however, preferable that multiple probes are used to characterize physical and electronic structure of the catalyst during reaction, over multiple time and length scales. To date it has not been possible to directly link the observations across these techniques in such a way as to confirm that the data (imaging, diffraction, spectroscopy) is obtained from the system in the exact same “working” state. Here we report an experimental approach that allows: (1) characterization – via x-ray absorption spectroscopy, extended x-ray absorption fine structure, x-ray fluorescence, Raman spectroscopy, transmission electron microscopy, scanning transmission electron microscopy, electron energy loss spectroscopy and energy dispersive electron microscopy – from the same sample, (2) characterization at atmospheric pressures in reactive environments, and (3) simultaneous, real-time and on-line analysis of the reaction products – i.e. “operando” experimentation.

We take advantage of recent developments in sample holders for transmission electron microscopy that allow catalysts to be confined between two, thin nitride membrane supports that are separated by a narrow gap, and that allow continuous flow of liquid or gas through the system. We exploit the simplicity of this system in such a way as to allow utilization in both synchrotron x-ray beamlines and transmission electron microscopes. We have chosen a simple, model catalyst reaction for the demonstration phase of this work, the catalyzed conversion of ethylene to ethane, though the use of Pd/SiO2 and Pt/SiO¬2 heterogenous catalysts. Extension to high-temperature experimentation will be reported, thereby demonstrating the extension of this approach to the full class of catalytic systems. 


Research carried out at the Center for Functional Nanomaterials, Brookhaven National Laboratory, supported by the U.S. Department of Energy, Office of Basic Energy Sciences, under Contract No. DE-AC02-98CH10886. Y.L and A.F.F acknowledge additional support through the Synchrotron Catalysis Consortium, U.S. Department of Energy, Office of Basic Energy Sciences, under Contract No. DE-FG02-03ER15476.

Fig. 1: Compendium of data. RGA data demonstrating that the catalysts are in the same operating environment in both the TEM and the XAS experiment. XANES, EXAFS, HR-TEM and EELS data from the Pt/SiO2 catalysts during the dehydrogenation reaction.

Type of presentation: Poster

IT-6-P-1552 In situ and related TEM techniques for the characterization of chemically synthesized nanomaterials

Kamino T.1,2, Shimizu T.2, Yaguchi T.3
1Yamanashi University, Kofu, Japan , 2Japan Automobile Research Institute, Ibaraki, Japan, 3Hitachi High-Technologies, Ibaraki, Japan
tkamino@yamanashi.ac.jp

The wide range of techniques associated with an environmental transmission electron microscope (ETEM) are now applied to the studies of the nanomaterials used in the batteries and fuel cells. We have developed a specimen heating holder for in situ TEM observation of gas-solid reaction at elevated temperatures. The specimen heating holder has a nozzle for gas injection adjacent to the specimen. This unique design of the holder, together with the differential pumping system of the TEM, has made it possible to achieve a maximum gas pressure of about 100 Pa, at the specimen area, without affecting the vacuum of the electron gun chamber. The short gas path length, of only 50 cm, from a gas bottle to the nozzle, enabled us to replace the gas atmosphere in the specimen area, within several seconds. In recent years, the technique has been applied to the characterization of the degradation mechanism of the electrocatalysts of polymer electrolyte fuel cells as reported elsewhere. Most of the chemically synthesized metal-matrix nanoparticles contain organic elements as a residue, and if electron irradiation exceeds critical levels, the residue may be sputtered out from the particle, and the particle may crystallized. Therefore, careful control of illumination conditions and irradiation time are essential for the analysis of those materials. The results of our recent study revealed that the classic selected area diffraction (SAD) is one of the preferable techniques in the structural analysis of the composite. Figure 1 shows TEM images and SAD pattern of a Nafion coated Pt/GC electrocatalyst observed at 200kV with
the electron beam density of 25 mA/cm2 on the specimen. Except for a slight deformation of the GC layer, the morphology of the specimen remained unchanged during the electron beam irradiation, for 14 minutes. The microscope used for the observation was a Hitachi H-9500 environmental TEM equipped with a standard high resolution objective lens pole-piece. The size of the field limiting aperture, used for the SAD pattern observation, was 1.0 micrometer and the size of the selected area was smaller than 15nm. The result reveals that an intermediate voltage TEM can be applied to the study of beam sensitive composites if the observation condition is carefully controlled. However, the influence of the electron beam, in the phenomena observed, in situ, can not be completely avoided. From this point of view, we have developed a gas reaction device for ex situ TEM study ( Fig.2 ) . The chamber is designed for use with the TEM specimen heating holders so that comparison of the results of in situ experiments, with that of an ex situ experiments, or a combination of both experimental techniques, using same specimen heating holder, are possible.


Fig. 1:  TEM images(a,c) and SAD patterns (b,d) of a Nafion coated Pt/GC electrocatalyst observed at 200kV             with the electron beam density of  25 mA/cm2 . a,b : 0min, c,d : 14min.

Fig. 2: Desktopgas reaction device equipped with a TEM specimen heating holder

Type of presentation: Poster

IT-6-P-1598 Quasi in situ TEM characterization of Ni reduction in regenerated Ni/alumina catalysts

GAY A. S.1, DUBREUIL A. C.1, BROURI D.2, MASSIANI P.2
1IFP Energies Nouvelles - Rond point de l'échangeur de Solaize - BP 3 - 69360 Solaize (France), 2Laboratoire de Réactivité de Surface, CNRS-UMR 7197, UPMC - Site d’Ivry, 3 rue Galilée - 94200 Ivry-sur-Seine (France)
anne-sophie.gay@ifpen.fr

Ni/alumina are active and selective catalysts for the selective hydrogenation of pyrolyse gasoline produced by steam cracker. This reaction allows removing alkadiene and alkenyl aromatics from C5+ fraction, without hydrogenating the aromatic rings and forming saturated hydrocarbons. In order to extend catalyst life, regeneration of spent catalysts followed by a reduction step can be applied to obtain a metallic and redispersed catalyst with activity as close as possible to the one of the fresh catalyst.
In this work, such reduction of a regenerated Ni/alumina catalyst containing reoxidized Ni particles was studied. The morphological evolution of the nanoparticles during reduction was followed by “quasi in situ” TEM, using a special Gatan HHST4004 “Heating Environmental Cell Holder” equipped with an ex-situ reactor and allowing the observation of the same zone before and after thermal treatment, in controlled atmospheric conditions [1].
After regeneration, all Ni in the sample was in oxidic form, as shown by FFT and SAED analyses. Two kinds of Ni oxide particles were observed, namely (i) small well-dispersed particles with sizes centered around 11 nm and (ii) large hollow particles (up to 30 nm in diameter), probably formed by Kirkendall effect from the initial metallic particles, as previously reported for Co Fischer-Tropsch catalysts [2]. After 2 hours of reduction at 410°C in Ar-5% H2 (treatment performed in the sample holder reactor), a modification of the morphology of the large hollow particles was observed. Thus, the hollow spheres broke during the reduction into a group of smaller Ni° particles forming a ring-like aggregate whose size is the same as for the initial hollow particle (figures 1 and 2). Besides, small nanoparticles with sizes nearly as above were still present but smaller nanometric particles also appeared. Moreover, all nickel was fully reduced, whatever the type of the nanoparticle.
These observations are in good accordance with previous results obtained for Ni catalysts used in the partial oxidation of methane [3] and for cobalt FT catalyst [2]. In summary, the present study shows that regeneration and reduction of a spent Ni/alumina catalyst used in selective hydrogenation of pyrolyse gasoline leads to a reduced well-dispersed catalyst. Besides, this study highlights the interest of the "quasi in situ" TEM technique to follow morphological evolutions during activation and/or regeneration treatments of supported catalysts.

[1] E. Sayah, D. Brouri, P. Massiani Catalysis Today 218– 219 (2013) 10–17
[2] C. J. Weststrate and al., Top Catal (2011) 54: 811-816
[3] S. Chenna and al., ChemCatChem 3 (2011) 1051-1059


Fig. 1: Hollow NiO particle in the regenerated catalyst

Fig. 2: Same zone - Ring-like aggregate of Ni° particles after 2h of reduction at 410°C under Ar/5%H2

Type of presentation: Poster

IT-6-P-1815 Development of a TEM specimen holder system for catalytic materials

Hashimoto A.1, Takeguchi M.1
1National Institute for Materials Science, Tsukuba, Japan
Hashimoto.Ayako@nims.go.jp

   Catalytic materials are often used as particles dispersed on a support material. Reduction of their size to nanoparticles, clusters and atoms is crucial for improving their performance. Transmission electron microscopy (TEM) is an indispensable tool for characterization of catalyst nanoparticles. However, the observation environment in general TEM differs significantly from those in actual applications. In this study, we developed a TEM specimen holder system for in situ observation of catalytic materials.

   Figure 1 shows a schematic of the developed system. The specimen holder includes a heater, a gas nozzle for introducing gases to the specimen, and a gauge for measuring pressure near the specimen. Orifice plates are arranged above and below the specimen to create differential vacuum.

   Figure 2 shows a TEM image and a fast Fourier transform (FFT) pattern of Pt nanoparticles on an amorphous carbon film taken in 0.5 Pa vacuum at room temperature, i.e., under gas-phase TEM conditions. The used microscope was JEM-2100 (JEOL, Japan), operating at an accelerating voltage of 200 kV. Lattice fringes of the Pt nanoparticle were observed, as indicated by the arrow in the TEM image and by the satellite spots in the FFT pattern (inset). These results demonstrate that the developed specimen holder system is compatible with high-resolution imaging in low vacuum. We have applied this holder system to study the movement of Pt nanoparticles on carbon materials in a gas atmosphere at high temperatures.


This work was partly supported by the Japan Society for the Promotion of Science, Grant-in-Aid for Scientific Research and Center of Materials Research for Low Carbon Emission.

Fig. 1: A schematic of the developed specimen holder system.

Fig. 2: TEM image and FFT pattern of Pt nanoparticles on an amorphous carbon film taken under 0.5 Pa air at room temperature.

Type of presentation: Poster

IT-6-P-1924 Dynamic Observation of Gold Particles in Water by Environmental-cell TEM

Kawasaki T.1, 2, 4, Imaeda N.1, Murase H.1, Yamasaki K.3, Matsutani T.3, Tanji T.1, 4
1Nagoya University, Nagoya, Japan, 2Japan Fine Ceramics Center, Nagoya, Japan, 3Kinki University, Osaka, Japan, 4Global Research Center for Environment and Energy based on Nanomaterials Science, Japan
kawasaki@nuee.nagoya-u.ac.jp

An environmental transmission electron microscope (ETEM) is very powerful technique enabling to in-situ observation of specimens immersed in gases, electric/magnetic fields, and even in liquids by using a closed-type ETEM. In the closed system, diaphragms which are set at the top and the bottom of a space for the specimen to isolate from the vacuum are one of the most important components. Although silicon nitride (SiN) films are generally used for this purpose [1], charging effect on the diaphragms is inevitable because of the insulating SiN. In order to solve this problem, we have developed a carbon/SiN hybrid diaphragm which consists of amorphous carbon film as conductive base material and amorphous SiN thin layer coated on it [2]. In this study, the diaphragms were applied to in-situ TEM observation of motion of gold colloidal particles in water.
Figure 1 shows a photograph of our developed environmental liquid-cell. This consists of three parts; upper/lower grids and a spacer ring. Each of the grids had seven holes of ?0.1mm in diameter as windows for the electron beam, which were covered by the carbon/SiN hybrid membrane. The fluid specimen was sandwiched by these two membranes with controlling distance between them by 240 nm with the spacer deposited on the upper grid. The liquid-cell containing the fluid specimen was set in an environmental-cell specimen holder. In the experiment, the fluid specimen was gold colloidal particles having the size of 80±5.7nm in diameter dispersed in water in 1.1×107/?l concentration, which were observed by using a conventional TEM H-8000 (200kV; Hitachi High Technologies).
Figures 2 show results of dynamic observations of motions of the gold particles in water. A gold particle in FIG. 2(a) migrated slowly at a speed of 10 ~ 30 nm/s, as shown in a quantitative analysis of variation of the moving speed in FIG. 2(b). This is not a Brownian motion because the gold particle drifted in a single direction. In contrast, a gold particle in FIG. 2(c) moved quickly just after starting the electron beam illumination at high speed of more than 1200 nm/s. Although a next short movement occurred after the first quick motion, the gold particle subsequently stopped, as shown in FIG. 2(d). These different types of the motion mean driving force of the movement of the gold particles is not unique. In the present cases, it is considered that an effect of water flow and coulomb force by charging due to the electron beam irradiation caused the motions of the particles in FIG. 2(a) and (c), respectively.


References
[1] U.Mirsaidov, C.D.Ohl, and P.Matsudaira, Soft Matter 8, 7108-7111(2012)
[2] T. Kawasaki et al., Proc. M&M2011. (2011) 465.


The authors acknowledge the financial support of this work by the Grant-in-Aid for Scientific Research (C) from the Japan Society for the Promotion of Science (#25390078).

Fig. 1: Photograph of our developed environmental liquid-cell which consists of upper/lower grids and a spacer ring.

Fig. 2: TEM images of motion of gold particles ((a) slow and (c) quick movement) (b), (d) variation of speed of motions in (a) and (c), respectively.

Type of presentation: Poster

IT-6-P-1943 The Study of Ice Impurities Using the Environmental Scanning Electron Microscopy at Higher Pressures and Temperatures.

Neděla V.1, Runštuk J.1, Klán P.2, Heger D.2
1Institute of Scientific Instruments of ASCR, Královopolská 147, 61264 Brno, Czech Republic, 2Department of Chemistry, Faculty of Science, Masaryk University, Kamenice 5, 62500 Brno, Czech Republic
vilem@isibrno.cz

Natural ice and snow accumulate and concentrate significant amounts of impurities that can be stored or chemically transformed, and eventually released to the environment. The location of impurities and their interactions with the water molecules of ice have not yet been sufficiently clarified. The aim of this work is to observe an uranyl-salt brine layer on the ice surface using a back scattered electron detection and the ice surface morphology using a secondary electron detection under equilibrium conditions in a specimen chamber of environmental scanning electron microscope (ESEM).

Our specially modified ESEM AQUASEM II equipped with the YAG:Ce3+ backscattered electron detector, an ionization detector of secondary electrons, a special hydration system and a Peltier cooled stage was used. The pressures between 400-700 Pa, 50% water-vapor saturation, and the temperatures above 250 K were utilized in our experiments. At these conditions, the phenomena of etching and subsequent stripping of impurities are largely suppressed.

Our samples were frozen under atmospheric pressure on a silicon plate cooled by the Peltier cooled stage. The initial sample holder temperature was above –1°C. A droplet of pure water or the uranyl nitrate solution was exposed to freezing. The uranyl nitrate solution (0.01 M) acidified by perchloric acid to pH = 1 were used in our second experiment because the hydrolysis of UO22+ is suppressed and only a single species (i.e., a hydrated uranyl ion) is present under these conditions.

Figure 1A shows an ESEM image of the ice sample prepared by freezing of pure water under atmospheric pressure inside the specimen chamber. Different shapes and sizes (30–200 µm) of the ice grains can be distinguished. Due to the detection of secondary electrons (SE), which are sensitive mostly to the surface topography, the ice grain boundaries are visible as black lines with a bright halo. At this temperature, the ice crystals are covered with a disordered interface (also called quasi liquid layer), however it is too thin to be identified by ESEM. Since the amount of backscattered electrons (BSE) is related to the atomic number of the present elements, 92U-rich regions appear brighter, whereas the regions consisting of water molecules remain dark, see Figure 1B. The difference between pure ice and the frozen uranyl solution is largely manifested in the channels and pools of concentrated UO22+ solutions (bright) along with the individual ice grains (black). Pools are usually the largest at the triple junctions, although some may also be present on the ice surface. A liquid layer containing UO22+ was expected to be considerably more concentrated than the parent solution due to the freezing concentration effect.


This work was supported by the Grant Agency of the Czech Republic: grant No. GA 14-22777S.

Fig. 1: Figure 1: An ESEM image of ice prepared by freezing of a droplet inside the specimen chamber: (a) frozen pure water; a SE detector mode; 270 K, 695 Pa (5.2 torr); (b) the frozen uranyl salt solution; c = 10–2 M; a BSE detector mode; 267 K, 525 Pa (3.9 torr). Bar 100 um

Type of presentation: Poster

IT-6-P-1959 The Atmospheric Scanning Electron Microscope (ASEM) Observes Axonal Segmentation and Platelet Generation in Solution.

Kinoshita T.1, Motohashi H.2, Hirano K.1, Maruyama Y.3, Kawata M.3, Ebihara T.3, Sato M.3, Nishiyama H.4, Suga M.4, Yamamoto M.5, Nishihara S.1, Sato C.3
1Faculty of Engineering, Soka University, Tokyo, Japan, 2Institute of Development, Aging and Cancer, Tohoku University, Sendai, Japan, 3National Institute of Advanced Industrial Science and Technology (AIST), Tsukuba, Japan, 4JEOL Ltd., Tokyo, Japan, 5Tohoku University Graduate School of Medicine, Sendai, Japan
ti-sato@aist.go.jp

The new Atmospheric Scanning Electron Microscope (ASEM) is a Correlative Light-Electron Microscope (CLEM) [1]. Cell cultures of a few milliliters can be grown and differentiated directly in the removable ASEM dish, in a CO2 atmosphere if required. After fixation, the cells are imaged in situ, immersed in liquid and at atmospheric pressure, by optical microscopy (OM) and SEM in a fully correlative manner. Various dish coatings have been developed to increase the range of culturable cell types [2] allowing axonal segmentation and platelet generation to be investigated.

Axonal partitioning of neurons was correlated with specific cytoskeletal structures including microtubules. For this, isolated Drosophila primary neurons were grown on a poly-DL-ornithine-coated ASEM dish and immunolabeled for neuron markers HRP and BP102. Fluorescence microscopy demonstrated the localization of HRP in the whole axial fiber (Figure 1, B) and the specific localization of BP102 in the proximal region (A). ASEM revealed a hexagonal frame-like structure of BP102 at the boundaries of the most intra-axonal segments (E, arrow) [4], which has not been observed by OM. Two tubulin bundles running alongside one another make contact at the intra-axonal boundary and possibly elsewhere. Eight of the ten axons examined showed such contacts. In the other two axons, the immunolabeling was disconnected at the intra-axonal boundary [3]. This may mean that the two tubulin bundles are separated, although the possibility that labeling is prevented by proteins bound to the α-tubulins in this region cannot be excluded.

Mature megakaryocytes (MKs) generate beaded cell projections named proplatelets, and further shed off platelets, which are indispensable cellular components of blood for hemostasis. We cultured MKs on ASEM dish and fixed the cells at an appropriate timing captured with OM. The cells were stained with heavy metal solution and observed at high resolution with the inverted SEM. The pseudopodia extended beaded strings (Figure 2A-B), including vesicles (C). These vesicles are necessary for blood clot formation, which is related to cerebral or myocardial infarction under pathological conditions [4]. Immunolabeling of P-selectin indicates that the vesicles could be a-granules. Additional labeling of α-tubulin indicates their transportation on microtubules (D-E).

References

[1] H. Nishiyama et al, J Struct Biol 169 (2010), p. 438-449.

[2] Y. Maruyama et al, J Struct Biol 180 (2012), p. 259-270.

[3] Microsc. Microanal. in press, doi:10.1017/S1431927614000178

[4] Ultramicroscopy in press, (http://dx.doi.org/10.1016/j.ultramic.2013.10.010)


We thank Dr. Toshihiko Ogura at AIST for valuable discussions.

Fig. 1: Figure 1. Axonal segmentation. (A) Localization of HRP. (B) BP102 (red). (C)Merge. (D-E) ASEM. (D) BP102 accumulates forming a special hexagonal frame-like structure at the intra-axonal boundary(arrows). (Ultramicroscopy. in press).

Fig. 2: ASEM of platelet generation from MKs. Primary MKs with proplatelet formation cultured on an ASEM dish were fixed, stained with Ti-blue (A-B) or gold-tagged for P-selectin (C), and further for microtubule (D-E). (B) Arrowheads indicate beaded proplatelets. (C-E) Arrows indicate putative alpha-granules. (Ultramicroscopy in press).

Type of presentation: Poster

IT-6-P-1960 Time-resolved Observations of Single Protein's Motions Using Diffracted Electron Tracking (DET) with Wet Cell SEM

Ishikawa A.1, Ogawa N.1,2, Hirohata Y.1, Yohda M.2, Sekiguchi H.3, Sasaki Y. C.4
1Nihon University, Tokyo, JAPAN, 2Tokyo University of Agriculture and Technology, Tokyo, JAPAN, 3SPring-8/JASRI, Hyogo, JAPAN, 4The University of Tokyo, Chiba, JAPAN
ishikawa@phys.chs.nihon-u.ac.jp

Diffracted electron tracking (DET) method has been developed for obtaining the information about the dynamics of a single protein molecule[1,2]. DET can be performed using a Scanning Electron Microscope (SEM) equipped with a highly sensitive detector for electron backscattered diffraction (EBSD). DET can trace the rotating motion of individual nanocrystals linked to the specific site in the molecule. Fig.1 shows the principle of the DET and the parameters to be measured. (a) When the electron beam irradiates a nanocrystal, inelastically scattered primary electrons form a band-like EBSD pattern (EBSP) and the 3D motion of nanocrystals can be traced from the shifts of the pattern. (b) shows the rotation angle ω around a single axis and the rotation angles α, β, and γ of the principal lattice vectors , a, b and c of the nanocrystal, respectively, between each time step. For tracing the motion of protein molecule, we have developed the wet cell sealed with the very thin carbon film for SEM observation[1,2]. The EBSP can be obtained from the colloidal gold linked to chaperonin protein in water under the carbon sealing film of the wet cell. In DET, radiation damage of the specimen is the biggest problem. To reduce the damage, specimen supporting was improved, as shown in Fig.2. (a) When the chaperonin protein is fixed to the carbon sealing film, although the motion of the colloidal gold could be traced, no directional motion could be observed in both conditions with and without adenosine tri-phosphate (ATP) which causes the rotation of the chaperonin protein. With this supporting, the chaperonin was irradiated by both the incident electron beam and EBSD electrons, and so damaged it could not move. Therefore the chaperonin supporting system was changed. The molecules are fixed to thin tri-acetyl-cellulose film to the opposite side of the sealing film as shown in (b). With this supporting, the chaperonin is covered with the “shadow” of the colloidal gold and hardly irradiated by the electrons. With this supporting system, the motion of the colloidal gold was traced by DET. The mean square of displacement (MSD) of the rotation angles of the colloidal gold particles, in both conditions with and without ATP, are shown in Fig.3. (a) Without ATP, each MSD of the α, β, and γ is almost same, and no directional motion is observed. (b) On the other hand, with ATP, the magnitude of the γ is clearly decreasing compared to other angles. This means that the chaperonin, linked to the colloidal gold, increases rotational motion around the ND axis as shown in (c). These results correspond with other single protein observations using other techniques.

References [1] N. Ogawa et al., Scientific Reports, 3, 2201 (2013) 1-7 [2] N. Ogawa et al., Ultramicroscopy, 140 (2014) 1-8


This research was supported by the Japan Science and Technology Agency under the Core Research for Evolutional Science and Technology (CREST) program.

Fig. 1: Principle of DET and parameters to be measured. (a) Inelastically scattered electrons in the crystal form a band pattern and crystal motion can be traced from the shifts of the EBSP. (b) The rotation angle ω around a single axis and the rotation angles α, β, and γ of the principal lattice vectors a, b and c are measured.

Fig. 2: Improvement of specimen supporting system for DET to reduce the radiation damage for chaperonin protein. (a) The chaperonin protein is fixed to the carbon sealing film of the wet cell. (b) The chaperonin protein is fixed to thin tri-acetyl-cellulose film opposed to the sealing film and is covered from the electron beam by the colloidal gold.

Fig. 3: MSD of the rotation angles of chaperonin molecules measured by DET. (a) Without ATP, almost no directional motion is observed. (b) With ATP, the g was clearly decreasing compared to other angular. This means that the many motions are the rotations around ND axis (γ = 0) corresponding to the motion of each chaperonin protein (c).

Type of presentation: Poster

IT-6-P-1968 Investigation of hexagon shape nanoparticle growth mechanism using in-situ liquid Cell TEM

Ahn T.1, Kim Y.1, Hong P.1, Nam K.1, Kim Y.1
1Department of materials science and engineering, Seoul National University, Korea
an2027@snu.ac.kr

A number of efforts have been made on the synthesis of monodisperse nanoparticles with various morphologies to take advantage of their physical and chemical properties of nanoparticles attainable from the chemical composition and their dimensions. Growth of nanoparticle can be affected by many factors, such as temperature, surfactants, types of precursor, and its relative concentration. In order to understand the growth and the formation mechanism of nanoparticles, it was proven that the liquid cell TEM is one of the most powerful techniques because of its nanometer-level spatial resolution with transparency of the internal structure. Researchers were able to observe the growth procedures in real time using the liquid cell TEM, which made a great fore-step to observe nanoparticle growth using electron beam induced process.
Growth behaviors of spatially aligned, hexagon single crystals were investigated from the liquid cell with D.I water based solution. Formation sequence was observed from home-built liquid cell TEM stage in JEOL 2010F. No intentional heating was made during the crystallization. Streaming video was recorded from the Gatan ES500W camera. Figure1 shows the boundary regions with and without the nanoparticles formed by the electron beam induced growth. Inhomogeneous distribution of the particles might be come from exposure time difference for the particle formation. Figure2 shows the snapshots from the streaming video taken while shining electron beam onto the liquid layer. Nanoparticles were formed in spherical shape up to 123 second irradiation. However as irradiation time increased, nanoparticles were gradually changed to hexagonal shape. Orientation alignment and the growth rate were measured from the snapshots up to 260 second irradiation. Effect of precursor concentration and the electron current density on the formation and the growth of nanoparticles were examined.


This research was supported by the Nano-Material Technology Development Program(Green Nano Technology Development Program) through the National Research Foundation of Korea funded by the Ministry of Science, ICT & Future Planning (2011-0019984)

Fig. 1: Low magnification image of the field of view. The regions in which electron induced growth occurred or not.

Fig. 2: Snapshots from the streaming video. As time increases hexagon shape nanoparticles growth were confirmed.

Type of presentation: Poster

IT-6-P-2009 Development and application of environmental high voltage electron microscope

Wakasugi T.1, Isobe S.2, Wang Y.2, Hashimoto N.1, Ohnuki S.1
1Graduate School of Engineering, Hokkaido University, 2Creative Research Institution, Hokkaido University
wakatake@eng.hokudai.ac.jp

Introduction:

For the practical use of hydrogen storage materials, the improvement of their hydrogen storage properties has been required. In order to improve the properties, we should understand the mechanisms and the dynamics of the materials in nano scale. Recent research indicated that a strain field introduced in the materials affected on the properties [1]. For understanding the dynamics in nano scale not only for improving the hydrogen storage properties, but also for developing other functional materials, in-situ TEM is the best way. In this research, we developed a high pressure gas Environmental cell for High Voltage Electron Microscope (EHVEM) and applied to in-situ High-Resolution (HR) observation with increasing hydrogen gas pressure.

Experimental:

The sample was a Pd thin film (~10nm-thick) deposited on a Silicon Nitride (SiN) window film. Firstly, this film was observed in a vacuum condition at room temperature. And then, an in-situ observation was carried out with the hydrogen pressure up to 40 kPa. The microscope used in this study was the EHVEM based on the JEOL ARM-1300 and operated at an acceleration voltage of 1250 kV.

Result and discussion:

The EHVEM allowed an in-situ HR observation in a hydrogen pressure up to 40 kPa. As shown in the Figure.1, however, the lattice fringes of PdH0.6 (200) and (111) grain with a distance of 0.20 nm and 0.23 nm were barely visible at the pressure. This result suggested that there was a significant influence of the gas pressure on the image resolution despite the use of high voltage electron beam.

On the other hand, using the corresponding IFFT (Figure.2) allowed us to recognize some dislocation cores introduced around/in a hydride grain clearly. The result told us that the increase in the number and the distribution change of cores with the pressure.

Reference:

[1] N. Hanada et al. J. Phys. Chem. C 113 (2009) 13450-13455


Fig. 1: High resolution images and FFTs (the insets) corresponding to PdH0.6 (200) and (111) in the hydrogen pressure of 10, 20, 40 kPa.

Fig. 2: High resolution images (the inset) and IFFTs (from white square) corresponding to PdH0.6 (111) in the hydrogen pressure of 10, 20, 40 kPa.

Type of presentation: Poster

IT-6-P-2220 TEM Imaging of CO Oxidation Catalyst of Gold Nanoparticle on TiO2 in CO and O2 Environments

Tanaka T.1, 3, Yamamoto N.2, 3, Takayanagi K.2, 3
1Meijo University, 2Tokyo Institute of Technology, 3JST, CREST
ttanaka@phys.titech.ac.jp

Gold nanoparticle on TiO2 (Au/TiO2) is promising for application to a low temperature CO oxidation (2CO+O2→2CO2) catalyst [1-4]. It is reported that the catalytic reaction proceeds at a peripheral region of the Au/TiO2 interface [1, 4]. It is proposed that the catalysis emerges from a negatively charged O2 molecule (O2) [1, 4], which is generated by Au-Ti co-bonding [4] and/or interstitial Ti ion [5]. We have studied the structure and electronic states of Au/TiO2 by using advanced TEM techniques [2, 3].
Interstitial Ti ions in a TiO2 substrate with and without a gold nanoparticle were observed by aberration corrected TEM [3]. Interstitials of Iv sites were observed at TiO2(001) surface as shown in Fig. 1(a) (a). Interstitials were accumulated at a perimeter/interface of Au/TiO2, while interstitials were depressed in a peripheral area of the accumulated region. A specific phase, which seems an expansion of the interstitial-accumulated region, was observed at the edge of the Au/TiO2 interface [2]. The specific phase grew extensively by exposing to O2 gas at 100 Pa into a pillar which has a chemical composition of Ti1-xO2 (x > 0) in Fig.1 (g). We will discuss report the CO oxidation catalysis.

[1] M. Haruta, et al., J. Catal. 144 (1999) 175.
[2] T. Tanaka, et al., Surf. Sci. 604 (2010) L75.
[3] T. Tanaka et al., Surf. Sci. 619 (2014) 39.
[4] Z.-P. Liu, X.-Q. Gong, J. Kohanoff, C. Sanchez, and P. Hu, Phys. Rev. Lett. 91 (2003) 266102.
[5] S. Wendt et al., Science 320 (2008) 1755.


The presentworkwas supported by a Grant-in-Aid for Scientific Research (A) (No. 16201020) and Exploratory Research (No. 24656031) of Japan Society for the Promotion of Science (JSPS). We thank Associate Professor N. Yamamoto (Tokyo Institute of Technology) for his valuable comments and stimulating discussion.

Fig. 1: (a) 

Type of presentation: Poster

IT-6-P-2380 Controlled environment specimen transfer for investigation of catalysts by ETEM

Damsgaard C. D.1,2, Zandbergen H.3, Hansen T. W.1, Chorkendorff I.2, Wagner J. B.1
1DTU Cen, Lyngby, Denmark, 2DTU CINF, Lyngby, Denmark, 3Kavli, TU Delft, Delft, The Netherlands
cdda@cen.dtu.dk

The full text of the abstract is not available. Please contact the presenting author.


Type of presentation: Poster

IT-6-P-2451 Aberration-Corrected, Environmental TEM Studies on Carbon Nanotube Oxidation and the Influence of the Imaging Electron Beam

Koh A. L.1, Gidcumb E.2, Zhou O.2, 3, Sinclair R.1, 4
1Stanford Nanocharacterization Laboratory, Stanford University, Stanford, California 94305, USA, 2Department of Applied Physical Sciences, University of North Carolina at Chapel Hill, Chapel Hill, North Carolina 27599, USA, 3Department of Physics and Astronomy, University of North Carolina at Chapel Hill, Chapel Hill, North Carolina 27599, USA, 4Department of Materials Science and Engineering, Stanford University, Stanford, California 94305, USA
alkoh@stanford.edu

One of the major applications for carbon nanotubes (CNTs) is as field emission electron sources, for example in image displays and high-intensity medical X-ray tubes [1-3]. The emission currents and lifetimes of CNTs are found to decrease under less stringent vacuum conditions [4, 5]. Earlier reports of carbon nanotube oxidation performed in an external laboratory setting, and surveyed a posteriori with a transmission electron microscope (TEM), suggested that the nanotube caps were selectively attacked during the oxidation process [6, 7].

Recently, we reported the direct study on the structural changes in CNTs as they were oxidized in-situ using an aberration-corrected environmental TEM (ETEM) [8]. Nanotubes were identified and tracked for structural changes as they were heated to increasing temperatures in a 1.5 mbar, high purity (99.9999%) oxygen environment. In order to investigate the effect of gaseous oxygen molecules on the nanotubes, rather than ionized gas species, we established a protocol whereby heating and oxidation were performed without an imaging beam, and the changes on identifiable nanotubes were documented after purging the gas from the chamber. Our studies showed that the oxidation of multiwall CNTs proceeds layer by layer, starting with the outermost wall, and not initiating at the nanotube cap, as reported previously. Nanotubes with a larger number of walls (greater than six) were found to be more resistant to oxidation, with all walls remaining intact during the ETEM experiments [8].

To simulate the highly ionized environment, which is expected during field emission, we repeated these observations at room temperature in the ETEM in the presence of the imaging beam. Under such conditions, we found that more rapid attack takes place, even at room temperature, and both the hemispherical cap and side walls of the CNTs are vulnerable. The influence of the imaging electron beam in the observation of this gas-solid reaction in the ETEM will be discussed.

References:

[1] Q. H. Wang et al., Appl. Phys. Lett. 72 (1998) pp. 2912–2913

[2] G. Cao et al., Med. Phys. 37 (2010), pp. 5306–5312

[3] X. Qian et al., Med. Phys. 39 (2012), pp. 2090–2099

[4] K. A. Dean and B. R. Chalamala, Appl. Phys. Lett. 75 (1999), pp. 3017–3019

[5] J.-M. Bonard, et al., Ultramicroscopy 73 (1998), pp. 7–15

[6] P. M. Ajayan et al., Nature 362 (1993), pp. 522–525

[7] S. C. Tsang, P. J. F. Harris and M. L. H. Green, Nature 362 (1993), pp. 520–522

[8] A. L. Koh et al., ACS Nano 7(3) (2013), pp. 2566–2572


Funding from the National Cancer Institute grants CCNE U54CA-119343 (O.Z.), R01CA134598 (O.Z.), and CCNE-T U54CA151459-02 (R.S.) is acknowledged. We thank Dr. Bo Gao of Xintek for providing the CNTs.

Fig. 1: Aberration-corrected TEM images of multiwall carbon nanotubes (MWNT) at (a) 400°C, (b) 400°C after exposure to 1.5 mbar oxygen for 15 min, and (c) 520°C after exposure to 1.5 mbar oxygen for 15 min. The electron beam is blanked during the oxidation process. The MWNT is found to resistant to oxidation under these conditions.

Fig. 2: Aberration-corrected TEM images of multiwall carbon nanotubes (MWNT) acquired at room temperature and 1.0 mbar oxygen, with the imaging electron beam on. The red arrow indicates attack on the MWNT cap and side wall, due to exposure to the ionized oxygen. The time elapsed between (a) and (b) is 32 sec.

Type of presentation: Poster

IT-6-P-2499 Carbon gasification by silver nanoparticles followed in situ at atomic resolution under oxygen partial pressure in an Environmental TEM (ETEM)

EPICIER T.1,2, CADETE SANTOS AIRES F. J.2, AOUINE M.2, LANGLOIS C.1, BLANCHARD N.3
1University of Lyon, MATEIS, umr CNRS 5510, INSA de Lyon/Université Lyon I, 69621 Villeurbanne, FRANCE, 2University of Lyon, IRCELYON, umr CNRS 5256, Université Lyon I, 69626 Villeurbanne, FRANCE, 3University of Lyon, ILM, umr CNRS 5306, Université Lyon I, 69622 Villeurbanne, FRANCE
thierry.epicier@insa-lyon.fr

Generally, the gasification of hydrocarbon-based materials leads to the formation of syngas. Gasification can also be performed, at high temperature in oxidative environments, in presence of metal catalysts supported on solid carbon materials. Indeed, oxygen adsorbs and dissociates on the metal surface then interacts with the carbon at the interface with the metal leading to the formation of carbon dioxide; concurrently carbon is consumed and the metallic nanoparticle advances to maintain the interface with the carbon forming in this way a trench on the surface of the carbon. On structured materials such as graphite or graphene these trenches tend to be rather 2D [1,2] at the surface of the material. In this study we chose to study the gasification of a non-structured material (amorphous carbon) by silver-based nanoparticles in a Cs-corrected TITAN ETEM, 80-300 kV, recently installed at CLYM in Lyon. Samples were prepared according to a synthesis described in [4]; the silver based nanoparticles (NPs) hang on onto the supporting carbon films and gasification of this film is observed between 400 and 500°C under variable oxygen partial pressures (between 10-1 and 5 mbar). We could follow in real-time the dynamics of carbon gasification and catalyst evolution by high resolution imaging (Fig.1) unlike previous studies. In situ EELS yields complementary information (regarding oxygen) necessary to support the proposed mechanism deduced from our in situ study and schematically summarized in Fig.2: (i) at the beginning of the gasification experiment the NP have an hexagonal structure consistent with Ag2O or hexagonal-Ag (known to exist under the form of NPs or nanorods [5]) containing diluted oxygen (Fig.2 a-c); (ii) at a given moment during gasification the NP transforms to fcc-Ag, gasification slows down, the particle begins to shrink (which is consistent with the decomposition of a surface oxide) while a coating forms around it (Fig.2d); (iii) once the particle is completely coated gasification stops and the NP shrinking stops (Fig.2e).


[1] S.K. Shaikhutdinov, F.J. Cadete Santos Aires, Langmuir, 14 (1998) 3501.
[2] T.J. Booth et al., Nano Letters, 11 (2011) 2689.
[3] N. Severin et al., NanoLetters, 9(1) (2009) 457.
[4] S. Li et al., Chem. Comm. 49 (2013) 8507.
[5] I. Chakraborty et al., J. Phys.: Condens. Matter, 23 (2011) 325401.


Thanks are due to the CLYM (Centre Lyonnais de Microscopie, www.clym.fr) for its guidance in the ETEM project which was financially supported by the CNRS, the Région Rhône-Alpes, the ‘GrandLyon’ and the French Ministry of Research and Higher Education. The authors acknowledge S. Li, A. Tuel and D. Farrusseng for providing samples and L. Burel for her assistance in preparing them.

Fig. 1: video frames from an HREM sequence of about 6 minutes recorded in situ at 495°C under 0.6 mbar of oxygen (see text and fig. 2 for details).

Fig. 2: a) starting geometry; b) O2 dissociation at the silver surface; c) gasification of the carbon support and motion of the NP (arrow); d) oxygen diffusion inside the NP and decomposition of the formed oxide; e): reaction of the Ag, O species with the surrounding (irradiation damaged) carbon to form a protective shell around the silver core.

Type of presentation: Poster

IT-6-P-2722 Development Aberration Corrected Wet-ETEM System and Its Application to Pt/Carbon Fuel Cell Catalysts in Moisturized Gases Environments

Yoshida K.1,2, Zhang X.2, Hiroyama T.2, Boyes E. D.3,4, Gai P. L.3,4
1Institute for advanced Research, Nagoya University, Japan, 2Nanostructures Research Laboratory, Japan Fine Ceramics Center, Japan, 3The York Nanocentre, University of York, UK, 4Department of Physics, Electronics and Chemistry, University of York, UK
ky512@esi.nagoya-u.ac.jp

Environmental transmission electron microscopy (ETEM), which was first reported in 1997, has proven to be one of the most efficient tools for in situ visualisation of the deactivation of heterogeneous catalysts in a reactive gas atmosphere at the nanometer scale. Development of wet environmental TEM (wet-ETEM) was also an essential for in situ studies of liquid-catalyst reactions [1].
Here we report a progressive gas injection system (Figure 1) for the latest spherical aberration corrected environmental transmission electron microscope [2-4], which enables real-time/atomic-sacale observation in moisturised gas atmospheres. The newly developed wet-ETEM system [5] is applied to platinum carbon electrode (Pt/Carbon) catalysts in proton exchange membrane fuel cells (PEMFC) to investigate the effect of water molecules on the Platinum/Carbon interface during deactivation processes such as sintering and corrosion.
Pt/Carbon is a typical electrode catalysts in PEMFC. But it is now well established that degradation of the carbon support at the cathode limits the lifetime of Pt/Carbon catalysts and thus the performance of the PEMFC. We evaluated the robustness of the Pt/Carbon electrode catalysts using the new Wet-ETEM system (Figure 1(b)-(d)). Humidity in the E-cell was accurately measured/controlled using the quadrupole mass spectrometer (Figure 1(e)). Sintering and migration of Pt nanoparticles observed in moisturized N2 atmosphere was extremely faster than ones in pure N2 atmosphere as shown in Selected Area Captured (SAC) images of Figure 2. Fig. 2(b) shows connected Pt nanoparticles, which were typically observed in wet condition. The damage and shrinking of carbon is not reason of such connection because granular pattern of carbon support is still surviving. White contrast surrounding the Pt connections (arrowed in Fig. 2(b)) also indicates that thickness of carbon films became thinner at Pt/Carbon interface because of the hydrocarbon desorption. We considered that physical adsorption of water and hydroxylation of the carbon surface is a main reason of the higher mobility of Pt nanoparticles observed in moisturized N2 atmosphere. The present in situ observation suggested we should induce much stronger trapping sites on the carbon supports for use on cathode in the PEMFC. In situ microscopy to show the dynamic behaviour of the fuel cell catalyst is thus very valuable to improve understanding of the degradation mechanisms and thus improve robustness.

[1] Gai P. L. et al., Microsc. Microanal. 8 (2002) 21.
[2] Yoshida K. et al., J Electron Microsc. 61 (2012) 99.
[3] Yoshida K. et al., Nanotech. 24 (2013) 065705.
[4] Yoshida K. et al., Microsc. 62 (2013) 571.
[5] Yoshida K. et al, Invited Paper, MMC 2014 Conf. Proceedings, and Nanotechnology (sub).


The authors thank the EPSRC (UK) for Critical mass grant EP/J018058/1 and The JSPS for a Grant-in-Aid for Young Scientific Researchers (B) (No. 24710110).

Fig. 1: Fig 1. Design schema of the new wet-ETEM system (a), optical micrographs of (b) Humidifier on the hot stirrer, (c) Thermostatic chamber, (d) gas injection and differential pumping line of the ETEM. (e) QMAS spectra corresponding 24% moisturized nitrogen.

Fig. 2: Fig 2. (a) SAC image at 230s of a movie obtained from the Pt/Carbon samples in pure Nitrogen environment. (b) SAC image at 40s of a movie obtained from the Pt/Carbon in 24% moisturized Nitrogen environment. (c), (d) and (f) SAC images from other rigion in 24% moisturized Nitrogen environment at 0, 5.5 and 14 s, respectively.

Type of presentation: Poster

IT-6-P-2774 Experimental evaluation of Environmental Scanning Electron Microscopes at high chamber pressure [200 - 4000 Pascal]

Rattenberger J.1, Fitzek H.1,2, Schroettner H.1,2, Wagner J.1, Hofer F.1,2
1Graz Centre for Electron Microscopy (ZFE), Steyrergasse 17, 8010 Graz, Austria, 2Institute for Electron Microscopy and Nanoanalysis (FELMI), Steyrergasse 17, 8010 Graz, Austria, Graz University of Technology (TU Graz)
johannes.rattenberger@felmi-zfe.at

Environmental scanning electron microscopy (ESEM) is an established method to investigate uncoated insulators, organic or biological samples in their original state. The presence of the imaging gas inside the specimen chamber is responsible for the secondary electron (SE) detection caused by gas amplification and the generated positive gas ions suppress charging artefacts. Water vapour as imaging gas at high chamber pressure (800 Pascal at 4°C or 2809 Pascal at 23°C) enables the opportunity to investigate wet samples or by varying the pressure or temperature to do wetting experiments [1].
Nevertheless, at high chamber pressures (200 - 4000 Pa) the gas amplification of SEs decreases and the scattering of primary beam electrons inside the imaging gas increases, which degrades the signal to noise ratio (SNR) and prevents image acquisition. Especially for low acceleration voltages, which are typically used for biological samples, the increase in scattering strongly limits the area of applications.
To evaluate the high pressure performance of ESEM and to compare different electron microscopes, information about special resolution and detector type is not enough. The contrast in SE images vanishes at high pressure and the big advantages of elaborated and expensive field emission guns are wasted.
Therefore a key feature for ESEM manufactures and users should be the stagnation gas thickness (additional distance the electron beam travels inside the imaging gas above the pole piece) and the SNR in SE detection for high pressure application, a fact which is not taken into account at the moment [2].
By using a special designed faraday cup, the fraction of scattered and unscattered electrons can be determined and the stagnation gas thickness calculated (see figure 1) [3]. The SNR in SE images can be measured by analysing a single image displaying a copper wire on carbon tape.
Results are presented for different types of SE detectors and beam transfer conditions (see figure 2 and 3). All experiments were performed using a FEI ESEM Quanta 200 or 600 (field emission gun).
1. G.D. Danilatos, 1988. Foundations of Environmental Scanning Electron Microscopy. Adv. Electron Electron Phys. 71, 109–250.
2. G. D. Danilatos, J. Rattenberger, V. Dracopoulos, Journal of Microscopy, (2010), DOI: 10.1111/j.1365-2818.2010.03455.x
3. J. Rattenberger, J. Wagner, H. Schröttner, S. Mitsche, A. Zankel, Scanning 31, (2009), p. 107


The author wants to thank Gerry Danilatos (ESEM Laboratory, Sydney) for helpful discussions and the Austrian Research Promotion Agency (FFG) for financial support (PN 839958).

Fig. 1: Stagnation gas thickness (Θ) [mm] as a function of chamber pressure P [Pa] using the gaseous secondary electron detector (GSED) as pressure limiting aperture.

Fig. 2: Test images: copper wire on carbon tape (imaging gas: water vapor)

Fig. 3: Signal to noise ration SNR [dB] as a function of chamber pressure [Pa] using the gaseous secondary electron detector (GSED)

Type of presentation: Poster

IT-6-P-3009 Multi-slice simulations for in-situ HRTEM studies of nanostructured magnesium hydride at elevated hydrogen pressures of 1 bar

Surrey A.1,2, Schultz L.1,2, Rellinghaus B.1
1IFW Dresden, Dresden, Germany, 2TU Dresden, Institut für Festkörperphysik, Dresden, Germany
b.rellinghaus@ifw-dresden.de

Nanostructuring of many hydrides has been shown to reveal improved thermodynamic and kinetic properties, which are needed for both mobile or stationary applications of solid-state hydrogen storage materials. During structural characterization utilizing conventional (HR)TEM, however, hydrides such as MgH2 degrade fast upon the irradiation with the imaging electron beam due to radiolysis in vacuum and as a consequence, the hydride phase cannot be studied at highest resolution. This problem can be overcome using a novel nanoreactor recently developed by H. Zandbergen (TU Delft) that allows for in-situ TEM studies at elevated H2 pressures (up to 4.5 bar) and temperatures (up to 500°C) [1]. A point resolution of 0.18 nm has already been demonstrated experimentally for Cu nanocrystals [2].

We have studied the feasibility of HRTEM investigations of light weight metals such as Mg and its hydride phases with the nanoreactor by means of multi-slice HRTEM contrast simulations. Such a setup provides the general opportunity to fundamentally study the dehydrogenation and hydrogenation reactions at the nanoscale under realistic working conditions. We analyze the dependence of both the spatial resolution and the HRTEM image contrast on parameters such as the defocus, the metal/hydride thickness, the hydrogen pressure and the nanoreactor geometry in order to explore the possibilities and limitations of in-situ experiments with this reactor. Such simulations may be highly valuable to pre-evaluate future experimental studies.

Fig. 1 shows schematically the details of the nanoreactor as it was implemented in a super cell used for the multi-slice simulations. The hydrogen is encapsulated between two 20 nm thin α-Si3N4 windows with the metal/hydryde positioned on top of the the bottom window. First simulations were conducted for a metallic Mg film of varying thickness oriented with its [001] direction parallel to the electron beam. The slicing was chosen to account for the varying density of atomic scatterers along the beam direction. While the slice thickness was reduced to contain only a single layer of scatterers within the Mg layer, it was increased to 1 nm and 100 nm in Si3N4 and in the hydrogen containing volume, respectively. Fig. 2 shows as an example the simulated Weber contrast of a Mg column (averaged over 611 individual columns) with respect to the background due to the Si3N4 windows as a function of the Mg thickness and the defocus. (Simulation conditions: Linear imaging. gmax = 20/nm. Imaging parameters match a FEI Titan microscope at 300 kV for NCSI imaging. Absorption, the MTF of the CCD, and a noise level of 3% were included in the simulations.)

[1] T. Yokosawa, Ultramicroscopy 112 (2012) 47.

[2] J.F. Creemer et al., Ultramicroscopy 108 (2008) 993.


Fig. 1: Schematic illustration of the simulated super cell representing the nanoreactor used for in-situ HRTEM investigations of the (de)hydrogenation of Mg(H2).

Fig. 2: Mean Weber-type image contrast of a column of Mg atoms as obtained from averaging over 611 individual atomic rows along the [001] direction of a Mg film with varying thickness. Despite the two Si3N4 windows and some scattering from the hydrogen atoms the individual Mg columns can be clearly imaged for thicknesses below some 30 nm.

Type of presentation: Poster

IT-6-P-3318 Understanding catalytic properties of nanoalloys by using aberration corrected electron microscopy in gaseous environment

Ricolleau C.1, Nelayah J.1, Nguyen N.1, Prunier H.1, Wang G.1, Piccolo L.2, Alloyeau D.1
1Laboratoire Matériaux et Phénomènes Quantiques, CNRS-UMR 7162, Université Paris Diderot-Paris 7, Case 7021, 75205 Paris Cedex 13, France, 2Institut de Recherche sur la Catalyse et l’Environnement de Lyon, Université Lyon 1 – CNRS, Lyon, France
christian.ricolleau@univ-paris-diderot.fr

Catalysis is involved in most of industrial chemical processes for refining, pollution control and synthesis of chemicals. Heterogeneous catalysis has always used nanoparticles in order to maximize the surface/volume ratio of active particles. Therefore, “particle size effect” is a well-known concept of catalysis. Moreover, combining metals within catalysts can improve catalytic performances with respect to pure metals, e.g., increased selectivity or resistance to poisoning. The “alloying effect” is classically ascribed to either electronic structure or active site geometry. However, this phenomenon is poorly understood and controlled, due to the difficulty to elaborate homogeneous collections of multimetallic nanoparticles with imposed composition, and to the lack of structural characterization.
Our general objective is to get insights into the interplay between the structure of nanoalloys (Pd-Au, Au-Cu and Pd-Ir on oxide supports) with well-controlled size and composition and their catalytic properties. For that purpose, we have synthesized and characterized supported bimetallic nanoparticles, and analyzed their catalytic behavior by using a MEMS-based technology developed by Protochips Inc.. This MEMS gas cell allows to image and to follow the dynamics of nano-objects in an encapsulated gas environment as a function of the temperature. By combining this technology with our JEOL ARM 200F cold FEG aberration correction microscope, we can obtain images of nano-materials with an information limit better than 0.8 nm under 1 bar gas pressure and at 1000°C (Fig. 1).
From the study of the above systems, we want to address, by using this instrumentation, the following fundamental questions:
- How does the structure of supported nanoalloys depend on particle size and bulk phase diagrams (total miscibility for Pd-Au and Au-Cu vs. miscibility gap for Pd-Ir)?
- How does the chemical structure (ordering, random alloying, partial segregation, core-shell, etc.) of the nanoparticles influence their catalytic properties towards the series of selected prototypic catalytic reactions (oxidation, hydrogenation…)?
- How does the nature of the support drive the structure of the nanoalloys? What is the particle-support interface structure?
- How do temperature and gaseous environment affect the structure of the nanoalloys? Can we gain insights into the atomic mechanisms of sintering, redispersion and strong metal support interaction (SMSI effect)?
The concepts are not new but the methodology is novel and promising thanks to the recent development of gas cells technology that allows reproducing the real operando conditions.


Fig. 1: (a) TiO2 substrate classically used in catalytic reactions with nanoalloys imaged under 1 bar pressure of O2 and at 1000°C (with a JEOL ARM 200F cold FEG microscope equipped with an aberration corrector of the objective lens). (b) Enlargement of the rectangular area in dotted line of (a).

Type of presentation: Poster

IT-6-P-3521 Beam skirt resolution in Gaseous Scanning Electron Microscopy

Khouchaf L.1
1Univ-Lille Nord de France, Ecole des Mines de Douai, Douai France
lahcen.khouchaf@mines-douai.fr

The electron beam scattering by gaseous environment is the fundamental parameter limiting the performance of the Gaseous Scanning Electron Microscopy (GSEM). The result is the enlargement of the primary beam characterized by the radius skirt Rs. The scattering phenomena require a much closer re-examination. In fact, depending on the localization of EDX detector and the particles shape to analyze, the collected signal after the beam skirt will be different and Rs also will be different. So, except for homogeneous materials, Rs cannot describe the scattering behavior.
In fact, Danilatos, introduced the radius Rs which represents the radius containing 90% of the incident beam) as below:

RS = (364*Z/E)*(P/T)1/2.GPL3/2

where rs is the skirt radius, Z the gas atomic number, E the incident beam energy, P the pressure, T the temperature and GPL the gas path length.

As given by equation above, the value of Rs depends on the gas introduced, the incident energy, the pressure, the temperature and the working distance but does not depend on the total or individual cross section. In Figure 1 we can notice a nonlinear beahvior of Rs versus the pressure. In order to take into account this approach we introduced a surface of the skirt Ss instead of the Rs. In the case Ss is given by the equation below:

Ss = ∏*Rs2

Ss = α*P

Unlike Rs, expression given above, the equation above shows that Ss is a linear function versus the pressure. Figure 2 shows the evolution of Ss versus the pressure for water vapor and helium.
In this study, the surface skirt Ss instead of the radius skirt is introduced. Unlike Rs, the results show that Ss is a linear function versus pressure. This may help to use Ss in different scattering regimes and for a best interpretation of the consequences of electron scattering beam by gaseous environment. Examples are given with two gases environment: helium and water vapor.
References
G.D Danilatos, Scanning Microscopy 4 (1990) P. 799.
D Stokes in “Royal Microscopical Society”, ed. Mark Rainforth, (Wiley,Chichester) P. 221.
J.F Mansfield, Microchimica Acta 132, (2000), P. 137.
L Khouchaf et al, Vacuum 81, (2007), P. 599.
L Khouchaf et al, J.De Phys. IV France 118, (2004), P. 237.
L Khouchaf et al, Vacuum 86, (2011), P. 62.

L. Khouchaf. (2012). V. Kazmiruk (Ed.),  978-953-51-0092-8, InTech, Croatia (2012)

L Khouchaf et al, Microscopy Research, 2013, 1, 29-32.


Fig. 1: Variation of Rs versus pressure at 20 keV and GPL= 2mm for (a) H2O vapor, (b) He.

Fig. 2: Variation of Ss versus pressure at 20 keV and WD= 2mm for (a) H2O vapor, (b) He.

Type of presentation: Poster

IT-6-P-5740 Cultivation and Observation of HeLa Cells in the Microfluidic Environmental Electron Microscopy

Huang Y. C.1, Ma T. W.1, Huang T. W.1, Liu S. Y.1, Chen F. R.1, Tseng F. G.1, Chuang Y. J.2
1Department of Engineering and System Science, National Tsing Hua University, Taiwan, 2Department of BioMedical Engineering, Ming Chuan University, Taiwan
v800213@gmail.com

  Progress in the processing of wet tissues, without the need of fixation and other complex preparation, may facilitate the microscopic examination of tissues and cells. To solve the challenge that the moisture in the wet sample will be dried out by electron microscope’s vacuum system when observing the living cell, we attempt to develop an advanced MEMS wet-cell device with fluid-exchange to achieve the macromolecular dynamics observation accompanied with in-situ manipulating/monitoring under an electron microscope (EM), and then we report the observation of cells dynamics in solutions.

  In this study, we design a special wet chamber (liquid SEM capsule) for environmental SEM by MEMS technology[1], consisted of one in-frame and one out-frame fitting to each other with controllable gap between for cell incubation and EM observation. In the current SEM application, the environmental wet chamber, composed of a disposable out-frame and a capsule, is inverted in SEM after sealing for the sample to facing up toward the incident electrons for getting stronger signals. Then, we connect the PEEK tubes to the capsule and use a syringe pump to provide liquid circulation (Figure1).

  For living cell incubation inside wet cell, we immersed out-frame into culture dish to contain culture medium (DMEM with 10%FBS and 1%Penicillin/Streptomycin) , and then incubated HeLa cells for 8-12hr. To improve the image quality for the thick cell, we process cell permeabiliztion treatment (100%Methanol), immersing cell in the Milli-Q water in place of the cytoplasm. Comparing the image of different immersing time, we found that the two-hour immersion had a clearer view of cytoskeleton and nucleus (Figure2).Then we observed the living cell with our self-design component by fluid circulation way and recorded the cell division with slow flow rate(0.01ml/hr) under eight long hours OM observation (Figure3), which confirm the practicality of our design. Since the existence of liquid seriously influence the contrast, we replace the culture medium with glycerol, finding that the resolution is improved under long time SEM observation (Figure4). With the unique liquid circulating system incorporated with SEM, we can successfully incubate HeLa cells for a long period of time in the wet micro environment. The image resolution under a wet condition is characterized as 40-50 nm, suitable for observing interaction between virus and cells or subcell organelles.

1. ”Self-aligned wet-cell for hydrated microbiology observation in TEM.” T.W. Huang, S.Y. Liu, Y.J. Chuang, et al., Lab on a Chip.12:340-347(2012).
2. ”A Novel Method for Wet SEM,” Iris Barshack, Juri Kopolovic, Yehuda Chowers, Opher Gileadi, Anya Vainshtein, Ory Zik and Vered Behar, Ultrastructural Pathology, 28:29-31(2004).


This work was supported by National Science Council (NSC102-2321-B-007-007 and NSC 102-2120-M-007-006-CC1).

Fig. 1: Figure1. Liquid-SEM device for dynamical study on HeLa cells

Fig. 2: Figure2. Cell permeabilization and the Milli-Q water immersion for (a) 2 hr, (b) 4 hr. The image resolution under 2 hours immersion, which is 65 nm, is better than that under 4 hours immersion, which is 107nm.

Fig. 3: Figure3. Hela cell division growth under liquid circulation environment.

Fig. 4: Figure4. Dehydrated Hela cell with fluid circulation under environmental SEM. (a) Culture medium, (b) Glycerol.

Type of presentation: Poster

IT-6-P-5803 ETEM observation of degradation of platinum and platinum-cobalt alloy nanoparticle electrocatalysts on carbon black

Nagashima S.1, 6, Kang Y.2, 3, Yoshida K.2, 4, Hiroyama T.2, 4, Liu K.2, 4, Ikai T.5, Kato H.5, Nagami T.5, Kishita K.6, Yoshida S.1
1Materials Research and Development Lab., Japan Fine Ceramics Center, Atsuta-ku, Nagoya, 456-8587, Japan, 2Nanostructures Research Lab., Japan Fine Ceramics Center, Atsuta-ku, Nagoya, 456-8587, Japan, 3Dept. of Frontier Materials, Nagoya Institute of Technology, Showa-ku, Nagoya, 466-8555, Japan, 4Institute for Advanced Research, Nagoya University, Chikusa-ku, Nagoya, 464-8603, Japan, 5Catalyst Design Dept., Material Design Div., Toyota Motor Corporation, Toyota-cho, Toyota, Aichi, 471-8572, Japan, 6Material Analysis Dept., Material Development Div., Toyota Motor Corporation, Toyota-cho, Toyota, Aichi, 471-8572, Japan
s_nagashima@jfcc.or.jp

   The Proton Exchange Fuel Cell (PEFC) is expected as a promising energy source to use of Fuel Cell Vehicle. Platinum nanoparticles on carbon black (Pt/C) are a typical electrode catalyst used in the PEFC. For development of advanced electrode catalysts in the PEFC, reducing Pt usage and enhancement of the durability are still problems. In order to reduce the Pt usage, Pt-metal alloys were investigated in recent few years and previous researches have reported that Pt-Co alloy represented the high Oxygen Reduction Reaction (ORR) activity and improved stability on cathode condition of PEFC. For further design concept of the Pt-Co electrode catalysts, it can be essential to understand the degradation mechanism in real space about both Pt and Pt-Co alloy. Because actual Pt-Co catalysts consist of pure Pt nanoparticle, Co nanoparticle, the ordered Pt-Co alloy (L12 etc.) and disordered Pt-Co alloys. In this report, we evaluated structural changes of the Pt/C and the PtCo/C electrode catalysts during electrochemical degradation by using Environmental Transmission Electron Microscopy (ETEM). In addition to such ex-situ approach, we achieved the dynamic in-situ observation in controlled water (H2O) atmosphere which is known as product molecule of ORR in PEFC.

   For electrochemical degradation tests simulating the start and stop test of PEFC, potentio/galvanostat was used with a potential range from +1.00 V to +1.50 V, a scan rate of 500 mVs-1 in 0.1 M HClO4 at room temperature for 40,000 cycles. In our degradation tests, Electrochemical Surface Area (ECSA) of Pt/C decreased 57% and PtCo/C decreased 27% through 40,000 cycles (Fig. 1(e, f)). The size of nanoparticle and the surface area of nanoparticles were clearly increased in ex-situ TEM images obtained from both Pt/C and PtCo/C (Fig. 1(a-d), Fig. 2(a)). It indicated that Pt and PtCo nanoparticles grow and carbon black particles shrink during the electrochemical degradation test.

   Then, we achieved a dynamic observation to investigate the cause of the coalescence of Pt nanoparticles. In 10 Pa of water vapor (H2O), Pt nanoparticles rapidly diffused on the carbon surface and formed an interconnected structure (Fig. 2(b-e)). We considered that physical adsorption of H2O molecule and hydroxylation of the dangling bond on carbon surface were the main causes of the rapid mobility of Pt nanoparticles. Therefore, we consider that much stronger trapping sites on the carbon is needed to reduce the mobility of Pt and Pt-Co nanoparticles.


Fig. 1: (a-d), Ex-situ TEM images of Pt/C and PtCo/C before and after degradation. Size distribution of nanoparticles as an inset. (e, f), CV curves of Pt/C and PtCo/C obtained before and after degradation test over a potential range from +0.05 V to +1.20 V at a scan rate of 500 mVs-1 in 0.1 M HClO4 at room temperature.

Fig. 2: (a), Specific surface area of Pt and PtCo nanoparticles before and after degradation. It was calculated by presuming that volume of a carbon black particle is volume of a sphere with a diameter of 35 nm (the circles in Fig. 1(a-d)). (b-e), Selected area captured images of a movie obtained from Pt/C in 10 Pa of water atmosphere.

Fig. 3:

IT-7. In-situ microscopic techniques and cryo-microscopy

Type of presentation: Invited

IT-7-IN-2863 The opportunities and challenges of liquid cell electron microscopy

Ross F. M.1
1IBM T. J. Watson Research Center, Yorktown Heights, NY 10598, USA
fmross@us.ibm.com

Liquid samples, particularly samples containing water, have traditionally been difficult to examine using transmission electron microscopy because of the incompatibility between the microscope vacuum and the high vapour pressure liquid. But in recent years, advances in sample design have allowed us to enclose liquids in a form that permits examination by TEM. Microfabricated devices are constructed in which two electron transparent membranes are spaced 100nm-1um apart. A liquid is introduced between the membranes, allowing imaging of structures and processes in situ. The technique of liquid cell electron microscopy has been adopted by many laboratories worldwide, and is of interest to the microscopy, materials and biology communities because it enables data to be obtained at a spatial and temporal resolution not accessible with other techniques.

In this presentation we focus on the use of liquid cell electron microscopy to examine the mechanisms of electrochemical processes in aqueous electrolytes. Liquid cell microscopy is well suited for electrochemistry because electron-transparent electrodes can be included during device fabrication. Images and movies of the transient structures that form during nucleation or dissolution can then be correlated with electrochemical parameters (voltage, current) controlled or measured by a potentiostat. We show measurements made during deposition and stripping of metals (Cu, Zn) on Au or Pt electrodes. After nucleation and coalescence, we measure the propagation of the growth front outwards from the electrode and into the liquid layer. We will show that an initially planar growth front roughens and becomes unstable, forming dendrites or ramified patterns. Such growth instabilities can affect the charging of batteries and the electrodeposition of thin films and multilayers. We will show how the development of diffusion fields works together with kinetic roughening to cause the onset of growth instabilities. Techniques have been developed to control the onset of instability, including pulse plating, electrolyte flow and the use of additives to alter interface parameters. We will examine these approaches using liquid cell microscopy.

In any liquid cell experiment, obtaining quantitative data that is suitable for matching with models involves understanding the pitfalls and artifacts that can occur during liquid cell EM. We therefore discuss electron beam effects, in particular the strong changes in solution chemistry that can be induced by the beam. Beam-induced radiolysis of water can lead to phenomena such as particle and bubble formation. These can be minimised with low-dose techniques, but may also be useful in forming patterned structures and in measuring the properties of nanoscale bubbles.


The results presented here have been funded, in part, by the US National Science Foundation under grants 1129722, 1225104 and 1066573.

Type of presentation: Invited

IT-7-IN-6082 Connectivity between imaging tools under controlled conditions: learning’s from 20 years experience with a variable cryo transfer system for the future

Wepf R.1
1ScopeM/EMEZ, ETH Zürich, Zürich, Switzerland
roger.wepf@emez.ethz.ch

Sample preparation has become more crucial with modern microscopy and compositional analysis. The most obvious requirement, is that the specimen is reduced in size and exposed without relocating, changing or exchanging atoms or part of the sample, so to say with minimal or no alteration. Once prepared the transfer to the microscope is the last potential destructive step prior to the final analysis. To reduce such influences we therefore first established a controlled connection between a high vacuum cryo preparation device and a cryo-SEM (FEG-XL30 & cc-corrected LVSEM) in 1994 to avoid contamination of freshly prepared samples at cryogenic conditions and enhanced sample preservation and throughput for high resolution SEM work.

Soon later this system was brought to market by Bal-Tec under the name VCT, including a high resolution cryo-stage, and adapted to a large number of SEM’s, FIB/SEM, cryo-AFM, cryo-IonTof and others. Later it was also extended to ESEM’s for “inert-gas” or controlled environment sample handling.

Proofing that connectivity under “inert gas” or high vacuum between sample preparation devices and analysis devices combines higher sample quality with higher sample throughput without the risk of loosing samples due to contamination, change of structure by oxidation, amorphisation, cracking or simple loss of sample by remounting.

This kind of connectivity between single devices is well established in semi conductor industry in so-called fabrication plant or “FAB’s”-lines, where the sample (wafer) is handled between production and analytical stations such as LM, Spectrometer, EM, Auger- and SIMS instruments without any remounting and interference of an operator and mostly under high vacuum conditions for on-line quality and process control.

In structure research we often face the problem on non-periodic (non crystalline) samples that several independently and comparative studies do not merge into a common picture. Understanding materials heterogeneity at various order of scale very much depends on imaging and analysis the same area/region of interest (AOI/ROI) without the risk of changing the sample properties and configuration between the investigations. This not only helps to reduce multiple experiments but also allows to zoom-in on pin-point selected ROI by a correlative combination of analytical imaging investigations (EELS, X-ray, SIMS, APT).

If we want to maintain highest possible quality of our carefully prepared samples for multimodal analysis we need to establish versatile transfer devices between the different analytical tools. In addition we need to standardize sample handling and interfaces to be able to investigate “close to native” samples at different resolution and sensitivity scales ideally on the same ROI. This will not only help to save time and number of samples but improving the output of multimodal analysis.


Fig. 1: For connectivity between various analytical modalities a kind of (cryo) innert gas exchange workstation for sample exchange and remounting under controlled environmental conditions is needed. This exchange station should avoid influcences affecting the sample native or virgin composition and ultrastructure (Δ critical interface steps for transfer)

Type of presentation: Oral

IT-7-O-1567 In situ Scanning Transmission Electron Microscopy study of CuO reduction

Martin T. E.1, Lari L.1, Gai P. L.2, Boyes E. D.3
1The York Nanocentre and Department of Physics, University of York, UK, 2The York Nanocentre and Departments of Chemistry and Physics, University of York, UK, 3The York Nanocentre and Departments of Physics and Electronics, University of York, UK
tm526@york.ac.uk

Methanol is one of the most important basic components in the chemical industry (worldwide production approx. 45 million tons 2010). Furthermore, it has potential as an in situ source of hydrogen for fuel cells [1, 2]. Cu is one component used to catalyse methanol synthesis and consequently, understanding of the activation and deactivation mechanisms of Cu is necessary to improve both catalyst activity and lifetime. Due to the scale of methanol production, small improvements in catalytic technology can lead to large economic impacts and make green technologies, such as fuel cells, more financially viable whilst also providing improved function. The activation process (in this case reduction) required to transform the precursor, CuO, into catalytically active Cu is very significant in determining the final size, structure and distribution of catalytic nanoparticles [1]. Subsequent to reduction, deactivation mechanisms (such as sintering) cause the catalyst activity to reduce with time. Single atom imaging under reaction conditions in ESTEM (Environmental Scanning Transmission Electron Microscope) can provide insights into activation and deactivation mechanisms, as well as the basis for improved catalyst designs.

The ESTEM at the York JEOL Nanocentre has recently been modified to provide the unique capability to directly visualise single atoms and the atomic structure of heterogeneous catalysts, such as Cu, in a gas environment [3]. This has allowed the in situ reduction of CuO in H2 (Figure 1) and investigation of the temperature-pressure parameter space to ascertain effects on particle morphology and size. As temperature is increased the Particle Size Distribution becomes bimodal with the particles divided into two distinct categories (facetted and unfacetted, Figure 2). Using ESTEM combined with EDXS, CuO particles and the more facetted Cu particles are seen to coexist. This suggests that reduction is dependent on the characteristics of the particle in question and thus that an atomic scale observation is required to fully understand the reduction process. Subsequent deactivation of the Cu particles is driven by reduction of the surface free energy and is shown to be primarily via the Ostwald Ripening (OR) mechanism (Figure 2). Understanding of the OR mechanism at the atomic level is currently lacking and the single atom resolution of the ESTEM, combined with Kinetic Monte Carlo simulations, provide a unique perspective on the factors affecting sintering such as particle size, temperature, activation energy and particle distribution.


The authors thank the EPSRC for support from critical mass grant
EP/J018058/1

References:

1. Hansen, P.L., et al., Science, 2002. 295(5562): p. 2053-2055.

2. Avgouropoulos, G., et al., Applied Catalysis B: Environmental, 2009. 90(3): p. 628-632.

3. Boyes, E.D., et al., Annalen der Physik, 2013. 525(6): p. 423-429.

Fig. 1: Reduction of CuO particles in situ using ESTEM at 3Pa Hydrogen, 361˚C to Cu. Diffraction patterns observed before and after reduction in UHV TEM.

Fig. 2: (a) Bimodal distribution suggests 2 groups of particles. These can be seen as grey (A) and white (B-more facetted) particles, (b) before reduction (c) after heating at 312˚C at 2Pa, (d) after heating at 361˚C at 3Pa. Particles become more facetted with reduced surface area as sintering process occurs.

Type of presentation: Oral

IT-7-O-1894 In-situ (S)TEM redesigned: Concept and electron-holographic performance

Börrnert F.1,3, Riedel T.2, Müller H.2, Linck M.2, Büchner B.1,3, Lichte H.1
1Technische Universität Dresden, Germany, 2CEOS GmbH, Heidelberg, Germany, 3IFW Dresden, Germany
felix.boerrnert@tu-dresden.de

The progress in (scanning) transmission electron microscopy and electron holography has led to an unprecedented knowledge of the microscopic structure of functional materials at the atomic level. Nevertheless, in-situ studies inside a (scanning) transmission electron microscope ((S)TEM) are extremely challenging. Here, we introduce a concept for a dedicated in-situ (S)TEM with a large sample chamber for flexible multi-stimuli experimental setups.

In conventional (S)TEMs the sample space is restricted by the pole pieces of the objective lens to a few millimeters; additionally, the sample is immersed into a strong magnetic field forbidding the investigation of magnetic phenomena. The solution to this problem is a radical redesign of the sample chamber and thus an adaptation of the electron optical layout. A versatile in-situ sample chamber requires space and access ports to incorporate different devices for applying various stimuli. This can be achieved by the use of a spherical-aberration corrected Lorentz type objective lens [1]. The size of the sample chamber is not anymore restricted by the electron optics and can be easily adapted to emerging experimental demands. Also, for the large-area control of experimental setups in situ a scanning surface imaging mode, i. e. a secondary electron detector, is needed.

A fundamental drawback of TEM is that the imaging process acts like an edge filter, thus no large-area field variations could be detected, and the image contrast is largely non-quantitative. In electron microscopy, the fully quantifiable image wave can be recorded only by an interferometric technique, i. e. off-axis electron holography [2]. Crucial for in-situ experiments is a large field of view while maintaining a high spatial resolution [3].

Here, we report on the state of the conversion of a JEOL JEM-2010F retro-fitted with two Cs correctors [4] from a dedicated low-voltage high-resolution (S)TEM into a large-chamber in-situ microscope. Both correctors are aligned to act as a corrected Lorentz lens in conventional as well as in scanning mode. The complete column section originally housing the pole pieces of the conventional objective lens will be replaced by a sample chamber providing multiple large ports for accessing the sample. Special care has been taken to make the chamber design most flexible.

[1] B Freitag et al., Microscopy and Microanalysis (2009), 184.
[2] H Lichte et al., Ultramicroscopy 134 (2013), 126.
[3] M Linck et al., Microscopy and Microanalysis (2010), 94.
[4] F Börrnert et al., Journal of Microscopy 249 (2013), 87.


The authors acknowledge financial support from the European Union under the Seventh Framework Programme under a contract for an Integrated Infrastructure Initiative (Reference 312483—ESTEEM2). We thank Prof. A. Kirkland (University of Oxford) for providing the SE detector.

Fig. 1: Scheme illustrating the conversion of the (S)TEM sample region. Green – electron beam, red – lens magnetic field, blue − sample.

Type of presentation: Oral

IT-7-O-1998 In situ STEM studies of reversible electromigration in thin palladium–platinum nanobridges

Kozlova T.1, Rudneva M.1, Zandbergen H.1
1Kavli Institute of Nanoscience, Delft University of Technology, Delft, The Netherlands
t.kozlova@tudelft.nl

Electromigration is a process in which a metallic contact line is thinned by passing a current through it, thus gradually displacing atoms and ultimately leading to its destruction [1]. The electromigration process in Pd–Pt nanobridges was investigated by in situ scanning transmission electron microscopy (STEM), using a FEI Titan operating at 300 keV. This technique together with a special electrical sample holder, built in-house, allows the nanobridge morphology transformations to be imaged down to the atomic scale during passage of electrical current [2]. Correspondent I–V curves are also recorded in real time. We focus in particular on the direction of material migration in relation to the electric current direction.
Polycrystalline Pd–Pt nanobridges with different lengths (500–1000 nm) and widths (200–500 nm), and a thickness of 15 nm were produced by e-beam evaporation from a metal alloy source onto a 100-nm-thick freestanding silicon nitride membrane [3] (Fig. 1a). The experiments were conducted in bias-ramping mode, i.e. a uniform increase in voltage from 0 V to a maximum of 350–600 mV (this was chosen in each separate experiment), followed by a decrease back to 0 V, sometimes a subsequent increase into the negative range (−350 to −600 mV) was done, followed by a decrease back to the original starting point of 0 V (Fig.1b).
Electromigration in Pd–Pt alloy [4] is quite different from the pure elements Pt and Pd. Material transport in Pt and Pd is very similar: after a recrystallization (which resembles that of the Pd–Pt alloy) the bridge gradually becomes narrower until a nanogap is formed, whereby grain boundary grooving is not a dominant feature. For the Pd–Pt alloy the dominant change is grain boundary grooving (which occurs near the cathode side), where the outer shape of the nanobridge is maintained. For high current densities (3 – 5×107 A/cm2), material transport in Pd–Pt alloy occurs from the cathode towards the anode side, indicating a negative effective charge. While polarity is changed, the voids formed near cathode side are refilled (Fig. 2). The reversal of material transport upon a change of the electric field direction could be the basis of a memristor.

[1] Ho, P. S.; Kwok, T. Rep Prog Phys 1989, 52, (3), 301-348.

[2] Rudneva, M.; Kozlova, T.; Zandbergen, H. Ultramicroscopy 2013, 134, 155-159.

[3] Gao, B.; Osorio, E. A.; Gaven, K. B.; van der Zant, H. S. J. Nanotechnology 2009, 20, (41), 415207.

[4] Kozlova, T.; Rudneva, M.; Zandbergen, H. Nanotechnology 2013, 24, 505708.


The authors gratefully acknowledge NIMIC and ERC project 267922 for support.

Fig. 1: (a) Typical TEM image of the initial configuration of the bridge. (b) Typical I–V curve for one loop in bias ramping mode.

Fig. 2: Snapshots from the STEM footage. (a) Initial view of the bridge. During electromigration, voids form on the cathode side (shown with wide arrow) and material accumulates on the anode size (b, d–e, g–h). When the current is reversed, the voids are refilled (c, f). White arrows indicate the direction of electrons.

Type of presentation: Oral

IT-7-O-2861 A cryo high-vacuum shuttle for correlative cryogenic investigations

Tacke S.1, Krzyzanek V.2, Reichelt R.1,3, Klingauf J.1
1Institute of Medical Physics and Biophysics, Muenster, Germany, 2Institute of Scientific Instruments of the ASCR, Brno, Czech Republic, 3Rudolf Reichelt initiated the project but unfortunately he passed away too early to see the results
s.tacke@uni-muenster.de

The preservation of the native state is the key element in sample preparation. In the case of hydrated objects, embedding in vitreous (amorphous) ice and subsequent examination under cryogenic (cryo) conditions are the means of choice [1,2]. Over the last years, cryogenic techniques such as cryo-electron microscopy (cryo-EM) or soft X-ray cryo-microscopy have become increasingly popular, as they provide a direct, unaltered view on the specimen [3,4].

However, to provide a snapshot of the pristine architecture of the specimen, cryo techniques require constant cooling below the recrystallization temperature of 138°K [1] and avoidance of any contamination. This has been proven to be particularly challenging in the case of correlative cryo investigations [4,5], since these methods include several transfer steps due to their extensive post-processing [6] and complex workflow [7]. In the past, several transfer concepts were introduced and they are now commercially available. However, these systems are limited either by not offering a high-vacuum environment or constraining the applications to a restricted workflow.

Here, we introduce an improved cryo high-vacuum transfer system (CHVTS) that allows for the first time to combine all kinds of cryogenic experiments. Moreover, we provide a solution that offers the highest degree of freedom in terms of connectivity of experiments (Fig.1). As shown in the detailed scheme of the CHVTS, our system is composed of cartridge, storage unit and cryo high-vacuum shuttle (Fig. 2). Once vitrified and mounted to cryo-holder cartridges (CT3500, Gatan) up to eight samples can be transferred to the storage unit. Thereafter, the cartridges can be transferred to the electron microscope or any other system extended by our docking device. A constant vacuum level of 7 ± 2 x 10-7 mbar and a temperature well below 133°K guarantee a contamination free transfer (see Fig. 3). Taken together, the CHVTS introduced in this work streamlines the handling of the frozen-hydrated specimen while solving for all problems generally associated with cryogenic investigations.

[1] J. Dubochet et al, Q. Rev. Biophys. 21 (1988), p. 129.

[2] L. Fitting Kourkoutis, J. M. Plitzko, and W. Baumeister, Annu. Rev. Mater. Res. 42 (2012), p. 33.

[3] S. G. Wolf, L. Houben and M. Elbaum, Nat. Methods online publication (2014), p. 1.

[4] C. Hagen et al, J. Struct. Biol. 177 (2012), p. 193.

[5] A. Rigort et al, J. Struct. Biol. 172 (2010), p. 169.

[6] S. Rubino et al, J. Struct. Biol. 180 (2012), p. 572

[7] E. Villa et al, Curr. Opin. Struc. Biol. 5 (2013), p. 771.


This research was supported by the DFG Grants RE 782/11-1,-2. Vladislav Krzyzanek acknowledges the support by the grant 14-20012S (GACR). We kindly acknowledge the help of the precision mechanical workshop, especially Martin Wensing. Additionally, we would like to thank Ulrike Keller for providing the EM grids, Harald Nüsse and Roger Wepf for numerous discussions.

Fig. 1: Theoretical concept for the connection of different types of cryogenic experiments. Green: established techniques. Red: possible extensions of the workflow.

Fig. 2: Main parts of the cryo high-vacuum transfer system: a) Cartridges assembly: (*) universal cartridge (**) EM grid (***) Clip for fixing the grid. Scale bar: 3mm. b) Storage device: (*) Cooling stage (**) docking device. Scale bar: 18cm. c) Cryo high-vacuum shuttle and pressure measurement during uncoupling of the shuttle. Scale bar: 21cm.

Fig. 3: Temperature measurement during the transfer of the cartridge.

Type of presentation: Oral

IT-7-O-2696 The Effect of Electron Beam on Aqueous Solution Composition during Liquid Cell Microscopy

Schneider N. M.1, Norton M. M.1, Mendel B. J.1, Grogan J. M.1, Ross F. M.2, Bau H. H.1
1Department of Mechanical Engineering and Applied Mechanics, University of Pennsylvania, Philadelphia, PA 19104, USA, 2IBM T. J. Watson Research Center, Yorktown Heights, NY 10598, USA
schnic@seas.upenn.edu

Liquid cell electron microscopy has emerged as a powerful tool for the real time imaging of objects suspended in liquids, and for characterizing processes that take place in liquids with the nanometer resolution of the electron microscope. However, as with all microscopy experiments, the electron beam interacts with the sample. Energy transferred from the fast-moving electrons to the irradiated medium causes excitation and ionization, resulting in the generation of radical and molecular species, which for water include eh (hydrated electrons), OH, H+, H2, O2, and H2O2. The hydrated electrons, oxidizing agents, and gaseous species can cause, respectively, reduction and precipitation of cations from solution, dissolution of metals, and nucleation and growth of bubbles. A quantitative understanding of electron beam-induced effects is critical to assessing whether the electron beam significantly affects the imaged phenomenon, so that we can correctly interpret experiments carried out with liquid cells; design experiments so as to minimize and mitigate unwanted effects; and take advantage of beam effects. We have developed a mathematical model to estimate radiolysis products during electron microscope imaging. The model includes the production of species by the electron beam, their destruction by reverse reactions, and their continued diffusion and reaction outside the irradiated region. We compute the concentrations of radiolysis products as functions of beam intensity, beam geometry, time, position, and solution initial composition (Fig. 1). We will describe this model and use its predictions to delineate various phenomena observed during liquid cell electron microscopy. For example, we predict that radiation chemistry causes large changes in pH within the irradiated region (Fig. 1a), and localized concentrations of reducing agents (Fig. 1b, c) and oxidizing agents. The pH of neat water can drop from 7 to 3.5 or lower within the imaged region under normal imaging conditions. Changes in pH can have significant effects on the phenomena under observation and may be the cause of aggregation of colloids (Fig. 2) that was observed during liquid cell imaging. We will compare the model with experiments carried out in a liquid cell, the nanoaquarium, at 300 kV in a Hitachi H9000 TEM and at 30 kV in an FEI Quanta FEG ESEM with a transmission detector, in each case imaging at 30 fps. The experiments and simulations suggest that liquid cell microscopy can provide a unique tool for studying radiolysis and for examining the behavior of materials subjected to high doses of radiation. We hope that the modeling tools described here will be useful for interpreting microscopy data obtained with liquid cells and for designing experiments that minimize unwanted effects.


The authors acknowledge funding, in part, from the National Science Foundation, grants 1129722 and 1066573.

Fig. 1: Heterogeneous model predictions of the concentrations of a) H+, b) e-, c) H, d) OH as functions of space and time. The beam (gray region) and liquid cell radii are, respectively, 1 µm and 50 µm. The beam current is 1 nA and the dose rate is 7.5x107 Gy/s. These values are typical for TEM imaging.

Fig. 2: Beam induced aggregation of 5 nm gold nanospheres in water. Dynamic imaging of large cluster-to-cluster aggregation (a-b) and early stage aggregation of small clusters (c-d).

Type of presentation: Oral

IT-7-O-2818 Cold-field emission and charge measurements of a carbon cone nanotip studied by in situ electron holography

de Knoop L.1, Gatel C.1, Houdellier F.1, Masseboeuf A.1, Monthioux M.1, Snoeck E.1, Hÿtch M. J.1
1CEMES-CNRS, Toulouse, France
ludvig.deknoop@cemes.fr

The cold-field emission gun (C-FEG) is the brightest electron source available, and also exhibits the smallest energy spread [1]. This technology has been greatly improved over the years concerning the electron optics and the vacuum, but the same cathode materials are still in use [2]. We have recently developed a new C-FEG source using a carbon cone nanotip (CCnT) mounted on a standard tungsten cathode using a focused ion beam (FIB) [3]. This source exhibits very good spatial coherence properties, which could be useful for electron interferometry applications [4].

Here, we have inserted a CCnT inside an in situ biasing transmission electron microscope (TEM) sample holder (Nanofactory Instruments) incorporating a nanomanipulator, in order to approach the CCnT towards a Au-anode plate (Fig. 1). We then ramped up the voltage between the nanotip and the anode from 0 to 95 V until the electric field around the tip was strong enough to allow the electrons to tunnel through the barrier and a field emission current could be recorded.

We have previously reported of how quantitative information of the local electric field of the CCnT (Eloc = 2.55 V/nm at the onset of field emission at 80 V) could be obtained by using off-axis electron holography and finite element method (FEM) modeling (Fig. 3 b)). By combining this with the Fowler-Nordheim equation [5], also the work function of the CCnT (Φ = 4.8 ± 0.3 eV) could be found [6].

Knowing the local electric field and the work function, the study has been expanded further to focus on the accumulation of charges on the CCnT before, at and after the onset of field emission. This was done with a technique that we recently have developed [7], which quantitatively measures the number of charges by applying the elegance and power of Gauss’s Law to electron holograms (Fig. 2). It provides a direct measurement of the charge inside a contour integral, with a sensitivity of one unit of charge.

The number of accumulated charges and the charge density on different places on the tip has been determined. We will show quantitative charge measurements along the CCnT as a function of applied voltage (Fig. 3 a)). Particularly the charge density at the beginning and during the field emission process provides some remarkable results. We will then discuss the importance of these values.

[1] O. L. Krivanek et al., Advances in Imaging and Electron Physics 153 (2008)
[2] A. V. Crewe et al., Rev. Sci. Instrum. 39 (1968)
[3] F. Houdellier et al., Carbon 50 (2012)
[4] F. Houdellier and M. Monthioux, International Patent Number WO2012035277 (2012)
[5] R. H. Fowler and L. Nordheim, Proceedings of the Royal Society of London 119 (1928)
[6] L. de Knoop et al., Micron, accepted (2014)
[7] C. Gatel et al., Phys. Rev. Lett. 111 (2013)


The authors acknowledge the European Integrated Infrastructure Initiative reference 312483-ESTEEM2 and the French "Investissement d'Avenir" program reference No. ANR-10-EQPX-38-01.

Fig. 1: The front part of the in situ TEM sample holder with a nanomanipulator for coarse and fine motion and biasing functionality. The inset shows the CCnT mounted on a wire in the nanomanipulator opposite the Au anode.

Fig. 2: Unwrapped hologram at field emission onset voltage of 80 V on the anode. The integration contour is indicated by small arrows, and the integration direction by the big, dotted arrow (analogous to the black dotted arrow in Fig. 3 a)).

Fig. 3: a) Number of electrons along the CCnT and in the vacuum for different voltages. b) Profiles of phase shift maps from electron holography and finite element modeling. The bias on the anode was 80 V and the tip-anode separation distance 680 nm.

Type of presentation: Oral

IT-7-O-2947 In Situ Analytical Electron Microscopy: Imaging and Analysis of Steel in Liquid Water

Schilling S.1, Janssen A.1, Burke M. G.1, Zhong X. L.1, Haigh S. J.1, Kulzick M. A.2, Zaluzec N. J.3
1Materials Performance Centre, The University of Manchester, Manchester UK 1, 2BP Corporate Research Center, Naperville, IL USA 2, 3Electron Microscopy Center, Argonne National Laboratory, Argonne, IL USA 3
m.g.burke@manchester.ac.uk

In situ transmission electron microscopy has become an increasingly important and dynamic research area in materials science with the advent of unique microscope platforms and a range of specialized in situ specimen holders. In metals research, the ability to image and perform x-ray energy dispersive spectroscopy (XEDS) analyses of metals in liquids are particularly important for detailed study of the metal-environment interactions with specific microstructural features. We have recently demonstrated that both STEM imaging and XEDS data can be successfully obtained from nanoparticles in liquid in an aberration-corrected FEI Titan G2 S/TEM with Super EDX [1] [2].  Furthermore, a special hybrid specimen preparation technique involving electropolishing and FIB extraction has been developed to enable metal specimens to be studied in the liquid cell TEM specimen holder [3].  We have applied these techniques to examine austenitic stainless steel in distilled H2O.

Conventional Type 304 austenitic stainless steel was prepared for examination in a Protochips Poseidon P200 liquid cell specimen holder with a 500 nm gap between the amorphous SiN windows. This specimen holder had been modified to optimize it for XEDS microanalysis [1]. TEM/STEM examination was performed using an FEI Tecnai T20 analytical electron microscope operated at 200 kV, equipped with an Oxford Instruments Xmax80TLE windowless Silicon Drift Detector (SDD) for XEDS spectrum imaging and analysis.  Fig. 1 shows the Type 304 steel specimen imaged in distilled H2O. Fig. 2a shows several crystalline particles that were observed after 24 hours in H2O. Spectrum images (Fig. 2b,c) obtained from this area revealed that these particles were enriched in Fe and depleted in Cr, and were consistent with the formation of an Fe-rich oxide. An XED spectrum (Fig. 3) obtained from the coarse angular oxide demonstrated that the particle was Fe-rich but also contained low levels of Ni. These Fe-rich oxides can form because the thin Cr2O3 film formed on the Type 304 foil surfaces depletes the matrix of Cr. Thus, any defects in the passive film will enable the local Cr-depleted matrix to oxidise, thereby forming Fe-rich oxides. Further studies on the development of surface oxides and coarse oxide particles can aid in the study of passive film development in steels.

References

1. Zaluzec, N.J. et al., X-ray Energy-Dispersive Spectrometry During In Situ Liquid Cell Studies Using an Analytical Electron Microscope. Microsc. Microanal. 20, in press doi:10.1017/ S1431927614000154 (2014).  

2. Lewis, E.A. et al., Wet Chemistry goes Nano. Submitted to Nanotechnology Letters.

3. Zhong, X.L. et al., Novel Hybrid Sample Preparation Method for In Situ Liquid Cell TEM Analysis. Submitted to Microsc and Microanal 2014.


The authors thank the BP 2013 DRL Innovation Fund, US DoE, Office of Basic Energy Sciences, and Contract No. DE-AC02-06CH11357 at the EM Center of Argonne National Laboratory.

Fig. 1: TEM image of steel sample in distilled H2O; 50 micron wide window.                               

Fig. 2: (a) STEM image and corresponding spectrum images for (b) Fe Kα and (c) Cr Kα obtained after 24 h in H2O.

Fig. 3: XED spectrum obtained from angular Fe-rich oxide (circled) that formed in H2O.

Type of presentation: Oral

IT-7-O-2952 In situ TEM observation of electrochemical deposition process

Oshima Y.1,2, Tsuda T.3, Kuwabata S.3, Yasuda H.1, Takayanagi K.2,4
1Research Center for Ultra-High Voltage Electron Microscopy, Osaka University, Osaka, Japan, 2JST–CREST, Japan, 3Department of Applied Chemistry, Osaka University, Osaka, Japan, 4Department of Condensed Matter Physics, Tokyo Institute of Technology, Japan
oshima@uhvem.osaka-u.ac.jp

Recently, in situ transmission electron microscope (TEM) observations of lithiation and delithiation processes in lithium ion battery have been achieved in order to improve the performance. However, the electrolyte is liquid in conventional lithium ion battery. In order to observe the lithiation and delithiation processes in situ, it is necessary to develop an electrochemical cell, which keeps the liquid electrolyte in vacuum.

In this study, we developed an electrochemical cell with three terminals (working, reference, and counter electrodes) and demonstrated the process of electrochemical copper deposition on gold surfaces in-situ. Figure 1 shows a photograph of our home-made liquid cell. It is composed of two quartz glass pieces which are glued by a heat-curing epoxy with each other. The observation window is covered with a 50 nm silicon nitride film to keep the liquid inside. The advantage of our cell is that it is available to arrange suitable materials as cathode or anode. In this observation, the working, reference and counter electrodes were gold (Au), gold (Au) and copper (Cu), respectively. The electrolyte contained 0.2M CuSO4 and 0.05 M H2SO4. Cyclic voltammetry (CV) was obtained by using a VersaSTAT4 with a scan rate of 25 mV/ s.

Figure 2 shows a series of TEM images taken during CV measurement and the corresponding CV curve. Darker background corresponds to the deposited Au thin film of about 30 nm in thickness. We observed that Cu clusters were nucleated on the Au film when the bias voltage was negative, while they were desorbed when the voltage was positive. And also, during measuring CV repeatedly, we observed that Cu clusters were nucleated at the same position, corresponding to dots of slightly darker contrast as shown in the TEM image of (a). We consider that these dots correspond to the position where gold atoms were alloyed with copper atoms.

In conclusion, we have developed a new electrochemical cell for in situ TEM observation. Using the liquid cell, we have demonstrated electrochemical Cu deposition process on thin Au film simultaneously with measuring cyclic voltammetry.


This research was supported by Japan Science and Technology Agency (JST).

Fig. 1: A photograph of our developed electrochemical cell.

Fig. 2: (a)-(d) A series of TEM images taken during measuring cyclic voltammetry. Graph of cyclic voltammetry curves.

Type of presentation: Oral

IT-7-O-3046 MAGNETIZATION REVERSAL PROCESS OF MAGNETIC SUPERDOMAIN STRUCTURES IN COBALT ANTIDOT ARRAYS

Rodríguez L. A.1,2, Magén C.1, Snoeck E.2, Gatel C.2, Castán-Guerrero C.3, Sesé J.1, García L. M.3, Herrero-Albillos J.4, Bartolomé J.3, Bartolomé F.3, Ibarra M. R.1
1LMA-INA, Universidad de Zaragoza, Zaragoza, Spain, 2CEMES-CNRS, Toulouse, France, 3ICMA, Universidad de Zaragoza-CSIC, Zaragoza, Spain, 4Centro Universitario de la Defensa, Zaragoza, Spain
luisaf85@unizar.es

Geometric confinement of the magnetization in magnetic thin films by patterning regular antidot (hole) arrays has been considered a potential method to fabricate storage media of ultrahigh capacity or magnonic devices for high frequency applications [1]. Reducing the distance between antidots modifies favorably the magnetic properties towards the creation of individual magnetic entities that could be used as magnetic bit of information [2]. For this reason, it is necessary to use magnetic imaging techniques that can provide information of the local magnetic states at submicron scales. In this work, high spatial resolution Lorentz Microscopy (LM) combined with the in situ application of magnetic fields has been used to perform quantitative studies of the magnetic states of cobalt square antidot arrays with periodicities (p) ranging between 524 and 95 nm. Antidot arrays were patterned by Focused Ion Beam (FIB) etching on a 10-nm-thick cobalt film deposited on Si3N4 membrane. The FIB etching process produced holes of 55 nm diameter. At remanence, defocused LM images revealed a periodicity dependence of the magnetic domains and a transition in the domain wall geometry around p ~ 300 nm, changing from 90° and 180° walls to superdomain (SD) walls for small periodicities (see Fig. 1). A Fourier filtered method has been implemented to improve the direct visualization of one-dimension SD (magnetic chains) for arrays with p > 95 nm (see Fig. 2), which has allowed to determine the magnetic configuration inside the antidot cells in both the SD and the SD walls. The magnetization reversal processes by means of hysteresis cycles upon magnetic fields parallel and diagonal to the antidot rows have been studied by in situ LM experiments. As illustrated in Fig. 3, we have found that the reversal magnetization process occurs by simultaneous (parallel hysteresis cycle) or sequential (diagonal hysteresis cycle) nucleation and propagation of horizontal and vertical superdomain walls, respectively.

[1] Xiao Z L, Han C Y, Welp U, Wang H H, Willing G A, Vlasko-Vlasov V K, Kwok W K, Miller D J, Hiller J M, Cook R E and Crabtree G W 2003 Nanotechnology 3 357.

[2] Torres L, Lopéz-Diaz L and Iñiguez J 1998 Appl. Phys. Lett. 73 3766.


This work was supported by the Spanish Ministry of Economy and Innovation (MINECO) through the projects MAT2011-28532-C03-02 and MAT2011-23791 including FEDER funding, by the Aragón Regional Government through Projects E26 (MAGNA), E34 (IMANA) and CTPP4/11, and by the European Union under the Seventh Framework Program under a contract for an Integrated Infrastructure Initiative Ref 312483-ESTEEM2.

Fig. 1: Defocused LM images for square antidot arrays with periodicities of (a) 524 nm, (b) 327 nm, (c) 160 nm and (d) 116 nm. In the latter, magnetic contrast has nearly disappeared.

Fig. 2: (a) Raw and (b) Fourier filtered defocused LM images of the antidot array with p = 160 nm. (c) Color-coded magnetization orientation map obtained by the Transport-of-Intensity Equation reconstruction to a small region marked with a yellow rectangle in (a) and (b).

Fig. 3: Sequence of filtered defocus LM images recorded during the in situ application of in-plane magnetic fields (a) parallel and (b) perpendicular to the antidot rows.

Type of presentation: Oral

IT-7-O-3113 In-situ biasing and switching of electronic devices into a TEM.

Mongillo M.1, Garbin D.1, Navarro G.1, Vianello E.1, Coue M.1, Mayall B.1, Cooper D.1
1CEA-LETI Minatec, 17 rue des Martyrs 38054 Grenoble, FRANCE
massimo.mongillo@cea.fr

In order to understand the physics of new materials that are currently being developed for use in electronic memories it is now necessary to perform in situ switching inside a Transmission Electron Microscope (TEM).
In this talk we will present our approach towards the development of a robust integrated characterization system that enables in-situ biasing and/or switching of electronic devices inside a TEM. As microscope time is valuable, the basic idea is to be able to electrically test a device before and then after specimen preparation outside of a TEM such that the electrical properties are understood before in situ operation in the TEM. The goal is to correlate the electrical properties to modifications in the crystalline structure and composition measured using HAADF STEM and EELS and the dopant/vacancy distribution measured by electron holography [1]. For this task we have been using a dedicated specimen holder featuring six static electrical contacts and a piezo-actuated movable probe tip which can act as a local electrical lead.
Figure 1 shows a TEM image of the movable tip used to switch a SrTiO3 resistive memory [2-3] that has been prepared using focused ion beam milling. The TEM image shows that the probe has introduced stress into the membrane. The poor electrical contact can also cause local heating and can even cause the specimens to explode. Despite these problems, the external electrostatic potential applied to the probe can cause a reversible switching of the active layer between high and low-resistive states, however the experiment is difficult, stressful and time consuming.
Our approach is to use fixed contacts on both simple and complicated devices. An example is provided in Figure 2 which shows a resistive memory cell [4]. A thick slice of the wafer has been sawn and then a Xenon-Ion FIB has been used to remove a large volume of material to provide a site specific region of interest. This region is then thinned to electron transparency using a conventional Ga FIB. Metal deposition in the FIB has been used to rewire the electrical contacts inside the device such that the switching can be performing by wire bonding the top contacts.
In this presentation we will present the two different approaches of switching memory devices in situ in the TEM and compare the advantages of each.

References:
[1]Nature Materials, 8, 271 (2009)
[2]Nature Materials, 5, 312 (2006)
[3]Advanced Materials, 21, 2632 (2009)
[4]Nature Materials,6, 824 (2007)


This work has been performed on the nanocharacterisation platform (PFNC) at Minatec. The authors thank the European Research Council for the Starting Grant “Holoview” and LabEx Minos.

Fig. 1: In-situ biasing using a probe manipulated by a piezo-electric motor. A thin TEM lamella prepared using conventional FIB specimen preparation techniques is mounted onto a TEM grid. The movable metallic probe (a) approaches and makes contact (b)to the top of the specimen.

Fig. 2: Memory cell prepared using Plasma FIB Xenon milling. The milling rate of the Xe-Fib enables us to remove large quantities of material. The electron transparent region containing the memory cell is patterned starting from a “bulky” slab. The two bonding pads can be hard-wired and used in the TEM to allow for in-situ switching of the device.

Type of presentation: Oral

IT-7-O-3133 In-situ Nano-compression tests on Shape Memory Alloys

San Juan J.1, Gómez-Cortés J. F.1, López G. A.2, Hernández J.3, Molina S.3, Nó M. L.2
1Dpt Física de la Materia Condensada, Univ. del País Vasco, Bilbao, Spain, 2Dpt Física Aplicada II, Univ. del País Vasco, Bilbao,Spain, 3Dpt Ciencia de los Materiales, Univ. de Cádiz, Puerto Real, Cádiz, Spain
jose.sanjuan@ehu.es

Recently, there has been growing interest in the potential use of shape memory alloys (SMA) in micro and nano-scale structures and devices, for example as sensors or actuators in micro electromechanical systems (MEMS). With a growing worldwide market in excess of hundred billion dollars, MEMS constitute a new paradigm of technological development for the present century, and smart materials are converging with miniaturization technologies, enabling a new generation of smart MEMS or SMEMS. Among the different smart materials targeted for use in SMEMS, shape memory alloys (SMA) have attracted considerable interest because they offer the highest work output density, about 107 J/m3, and exhibit specific desirable thermo-mechanical effects such as superelasticity and shape memory.

In previous works, completely recoverable superelastic strain and shape memory in micro and nano pillars was first reported for Cu-Al-Ni SMAs [1] showing the competitive advantage of these SMAs over the commercially used of Ti-Ni. In addition several size effects on superelastic behaviour were also demonstrated [2, 3] in Cu-Al-Ni SMAs. For practical applications the superelastic behavior must be reproducible in order to be functionally reliable, and first studies on cycling SMA micropillars by nano compression tests were recently published [4, 5].

In this work we present an In-situ characterization of the nano-compression superelastic behaviour of Cu-Al-Ni micro-pillars at the scanning electron microscope. Micro-pillars were milled by Focused Ion Beam technique on [100] oriented Cu-Al-Ni single crystals. All pillars were tested in an instrumented pico-indenter Hysitron PI-85, introduced inside the chamber of a JEOL-FEG 7500, by using a diamond flat indenter, as can be seen on Figure 1. The nano-compression stage was tilted in order to allow imaging by the SEM. Simultaneous video-image was taken during the nano-compression test acquisition data in order to correlate the mechanical behaviour with microstructure evolution. Fully recoverable and reproducible superelastic behaviour has been obtained and a picture of the screen containing both, image and mechanical test, is shown in Figure 2.

[1] J. San Juan, M. L. Nó, and C. A. Schuh, Advanced Materials 20 (2008), p. 272.

[2] J. San Juan, M. L. Nó, and C. A. Schuh, Nature Nanotechnology 4 (2009), p. 415.

[3] J. San Juan and M. L. Nó, J. Alloys & Compounds 577S (2013), p S25.

[4] J. San Juan, M. L. Nó, and C. A. Schuh, Acta Materialia 60 (2012), p. 4093.

[5] J. San Juan, J. F. Gómez-Cortés, G. A. López, C. Jiao, and M. L. Nó, Appl. Phys. Lett. 104 (2014), p.011901


The authors thank the Spanish Ministry of Economy and Competitiveness, MINECO, project MAT2012-36421 and the CONSOLIDER-INGENIO CSD2009-00013, and the Basque Government for Consolidated Research Group IT-10-310 and ETORTEK-ACTIMAT-2013. J. San Juan and M.L. Nó also thank EOARD Grant FA8655-10-1-3074. J.F. Gómez-Cortés thanks the Ph.D. Grant from MINECO.

Fig. 1: Figure 1. Sub micrometre pillar of Cu-Al-Ni SMA, milled by focused ion beam, just before the in-situ nano-compression test. On the lower side of the image it can be appreciated the flat tip of the diamond indenter in contact with the top of the pillar.

Fig. 2: Figure 2. In-situ superelastic nano-compression test performed on a sub-micrometre pillar of Cu-Al-Ni SMA. The video image allow to correlate the different points of the load-displacement curve with the corresponding images taken at the SEM. The above image corresponds to the final screen image just after finishing the in-situ test.

Type of presentation: Oral

IT-7-O-3169 Temperature-induced sphere-to-tetrapod transformation of CdSe nanocrystals investigated by in-situ transmission electron microscopy

van Huis M. A.1,2, Fan Z.3, Li W. F.1, Yalcin A. O.2, Tichelaar F. D.2, Talgorn E.4, Houtepen A. J.4, van Blaaderen A.1, Vlugt T. J.3, Zandbergen H. W.2
1Condensed Matter and Interfaces, Debye Institute for Nanomaterials Science, Utrecht University, Princetonplein 5, 3584 CC Utrecht, The Netherlands, 2Kavli Institute of Nanoscience, Delft University of Technology, Lorentzweg 1, 2628 CJ Delft, The Netherlands, 3Process and Energy Laboratory, Delft University of Technology, Leeghwaterstraat 39, 2628 CB Delft, The Netherlands, 4Department of Chemical Engineering, Delft University of Technology, Julianalaan 136, 2628 BL Delft, The Netherlands
m.a.vanhuis@uu.nl

Colloidal CdSe nanocrystals (NCs) can be synthesized in a wide variety of (heterogeneous) nanostructures including sphere, rod, tetrapod, and octapod morphologies. Using a low-drift TEM heating holder [1] employing MEMS microheaters with 15 nm thick SiN windows, the thermal evolution of spherical CdSe NCs was followed in real time and with atomic resolution. With increasing temperature, the NCs were found to transform from spheres to multipods to rectangular single crystals. The thermal evolution is shown schematically in Figure 1. 

The as-synthesized CdSe NCs consist of multiple subcrystals, but are spherical in shape. Upon heating to a temperature of 80 °C, most NCs transform into bipods, tripods, or tetrapods, whereby the core exhibits the zinc blende (ZB) crystal structure while the pods have the wurtzite (WZ) crystal structure [2]. These multipods are remarkably stable, up to a temperature of 300 °C, as long as they remain isolated. Multipods that are close together, though, fuse into rectangular single crystals having the ZB structure at temperatures of 170-200 °C. This is an unexpected result, as the stable bulk phase of CdSe is WZ. The ZB NCs undergo multiple crystal fusion events by oriented attachment, which was recorded in real time and with atomic resolution. The fusion is followed by coalescence into larger agglomerates which eventually transform to the WZ crystal structure. The last step in the thermal evolution is sublimation which takes place at temperatures of 360‒400 °C. 

Force-field molecular dynamics (FF-MD) simulations [2] (Figure 2), and density functional theory (DFT) calculations were performed in order to investigate the driving forces inducing these most remarkable transformation phenomena. It is concluded that off-stoichiometry can slightly favor the bulk ZB phase with respect to the bulk WZ phase, but that temperature-dependent interface-related energies are most likely the cause of the rich thermal behavior. Furthermore, it becomes clear that the ZB-WZ transformations are mediated by vacancies on the {111}Cd or {0001}Cd atomic planes.

[1] M.A. van Huis, N.P. Young, G. Pandraud, J.F. Creemer, D. Vanmaekelbergh, A.I. Kirkland, H.W. Zandbergen, ‘Atomic imaging of phase transitions and morphology transformations in nanocrystals’, Adv. Mater. 21 (2009) 4992-4995.
[2] Z. Fan, A.O. Yalcin, F.D. Tichelaar, H.W. Zandbergen, E. Talgorn, A.J. Houtepen, T.J.H. Vlugt, M.A. van Huis, ‘From sphere to multipod: thermally induced transitions of CdSe nanocrystals studied by molecular dynamics simulations’, J. Am. Chem. Soc. 135 (2013) 5869-5876.


Fig. 1: Schematic showing the thermal evolution of the CdSe nanocrystals, as they transform from spheres to tetrapods to rectangular zinc blende nanostructures, whereby the potential energy decreases

Fig. 2: Left: Result of a force-field molecular dynamics (MD) simulation performed at a temperature of 800 K, whereby a spherical CdSe NC has transformed into a tetrapod (details in Ref. [2]). Right: HRTEM images of several CdSe multipods formed during in-situ heating.

Type of presentation: Poster

IT-7-P-1394 An in-situ transmission electron microscopy study on room temperature ductility of TiAl alloys with fully lamellar microstructure

Kim S.1, Na Y.1, Yeom J.1, Kim S.1
1Light Metal Division, Korea Institute of Materials Science, Changwon 642-831, South Korea
mrbass@kims.re.kr

Gamma titanium aluminides (TiAl) have gained great interest for research on high-temperature applications due to their weight saving in combination with excellent high temperature properties such as creep and oxidation resistance. However, their poor room temperature ductility and machinability have hindered their application in areas such as aerospace and automobile products. In this study, mechanical properties of newly-developed TiAl alloys were investigated. The new TiAl alloys contain less aluminum compared with conventional gamma TiAl alloy to improve processibility and machinability. Especially, room temperature ductility of fully lamellar TiAl alloys was acquired without heat-treatment or TMP process.Adding beta stabilizers and lowering Al contents in conventional gamma-based TiAl alloys were found to be beneficial for room temperature ductility of TiAl alloys. An in-situ transmission electron microscopy study was conducted at room temperature in order to understand an underlying mechanism on room temperature ductility of TiAl alloys. From in-situ straining transmission electron microscopy experiments, it was revealed that the crack path is different between the TiAl alloys with/without room temperature ductility. The crack in TiAl alloys having room temperature ductility interacted with lamellae by forming bridging ligaments between the two alpha lamellae and the gamma lamellae (Fig. 1). In contrast, the cracks in TiAl alloys without room temperature ductility propagated along grain (colony) boundaries showing brittle intergranular fracture (Fig. 2). Finally, we proposed the important microstructural factors to have room temperature ductility of TiAl alloys.


This work was supported by the Fundamental R&D Program of the Korea Institute of Materials Science.

Fig. 1: Bright field images of alloy having room temperature ductilitytaken during in-situ TEM experiment.

Fig. 2: Bright field images of alloy having no room temperature ductility taken during in-situ TEM experiment.

Type of presentation: Poster

IT-7-P-1582 Simultaneous in situ SEM/STEM observation of catalyst reaction under an air atmosphere using a Cold-FE environmental TEM

Sato T.1, Matsumoto H.1, Nagaoki I.2, Yaguchi T.2
1Application Development Department, Hitachi High-Technologies Corporation, 2Advanced Microscope System Design Department, Hitachi High-Technologies Corporation
sato-takeshi-3@naka.hitachi-hitec.com

In the development of catalysts and fuel cell materials, there is an increasing demand for fine structural characterization using Environmental TEM (E-TEM). Recently, we have developed an E-TEM based on a conventional analytical TEM combined with a gas injection-specimen heating holder[1-2]. To further clarify the mechanism of the degradation of electrode-catalyst, a simultaneous in situ SEM/STEM study was carried out under the accelerated degradation condition. With the surface information from the SEM detector, we have obtained information in three-dimensional, gained significant new understanding of the behavior of Pt/C catalysts.
Figure 1 shows an overview and a schematic diagram of the specially designed Hitachi HF-3300 Cold-FE in situ TEM equipped with STEM and SEM imaging capabilities. In order to maintain the electron gun area under ultrahigh vacuum of better than 10-8 Pa near the gun, yet introducing gas into the specimen chamber, an additional ion pump (IP3) and an extra orifice have been added between the gun valve and specimen chamber. In situ simultaneous SEM/STEM observation in a gaseous atmosphere is realized by using the gas injection specimen heating holder.
A picture and a schematic diagram of the gas injection specimen heating holder are shown in Figure 2(a) and 2(b), respectively. The reaction gas is introduced to the area around the specimen by means of a gas injection nozzle. Therefore, in situ observation in a gaseous atmosphere can be carried out using this Cold-FE TEM even if up to 10 Pa near the specimen. The specimen was used a commercially available Pt/C catalyst. To simulate an accelerated aging, the specimen was heated to 200˚C. The morphological changes of Pt/C catalyst operated at accelerating voltage of 300 kV.
Figure 3 shows the results of in situ SEM/STEM simultaneous observation. At the beginning of air with a measured specimen chamber pressure of 1 Pa after 270 sec., a cluster of Pt particles on the carbon support have grown and agglomerated, and the number of Pt particles appeared to decrease on the carbon support, as indicated in red circle. After 660 sec., most of the Pt particles have gradually started inserting into the carbon support, and the grain growth and agglomeration of Pt particles have occurred inside the carbon support, as shown in Figure 3. After 1080 sec., the behaviors of migration, coalescence and grain growth of Pt particles inside the carbon support were clearly observed by STEM image.

References
[1] T. Yaguchi et al., J. Electron Microsc., 61(4), 199-206 (2012)
[2] T. Kamino et al., J. Electron Microsc., 54(6), 497-503 (2005)


The authors gratefully acknowledge Professor Kazunari Sasaki and Associate Professor Akari Hayashi of Kyushu University for valuable discussions.

Fig. 1: External view and a schematic diagram of the Hitachi HF-3300 Cold-FE TEM

Fig. 2: External view (a) and schematic diagram (b) of the gas injection specimen heating holder

Fig. 3: The results of in situ SEM/STEM simultaneous observation

Type of presentation: Poster

IT-7-P-1610 Observation method of cross-sectioned cells by cryo-scanning electron microscopy

Nishino Y.1, Ito Y.1, 2, Miyazawa A.1
1Graduate School of Life Science, University of Hyogo, 2Leica Microsystems K. K.
ynishino@sci.u-hyogo.ac.jp

Protein and cellular structures have been visualized in a close-to-native state by cryo-transmission electron microscopy (cryo-TEM). In many cases for cryo-TEM cells are so thick that we have to prepare cryo-ultrathin sections. In such case compression of cryo-sections must be taken into consideration. The compression makes the image complicated because it occurs inhomogeneously. Cells and organelles are compressed whereas small rigid complexes such as ribosomes and microtubules have been reported to resist compression.

On the other hand freeze-fractured cells and tissues have been examined by cryo-scanning electron microscopy (cryo-SEM). However this method is limited because observation objects are only randomly fractured surface.

In order to visualize non-distorted cross-sectioned cells, we focused on the block surface after cryo-sectioning, and imaged it by cryo-SEM. Budding yeastwas pelleted, high-pressure frozen and cryo-sectioned. The sections were imaged by cryo-TEM while the block was imaged by cryo-SEM. As a result, ultrastructure such as ribosomes and invaginated plasma membranes as well as organelles were clearly visualized by cryo-TEM, however vesicle structure such as whole cells, nuclei and vacuoles were ellipsoidal in the same direction (Fig. 1a). They were obviously compressed along the cutting direction. Meanwhile the block was transferred to cryo-SEM and observed. We could image cells and organelles without any staining or coating. In the cryo-SEM images yeast was oval in shape, and nuclei and vacuoles were circle in shape (Fig. 1b). They are consistent with the fluorescently-labeled images by light microscopy.

Furthermore we showed an example of repetitive cryo-sectioning and observation by cryo-SEM. A piece of diaphragm was cryo-sectioned in the direction parallel to sheet-like structure of diaphragm and observed by cryo-SEM. On the sectioned face near the surface of isolated diaphragm connective tissue was clearly observed (Fig. 2a). In order to observe structure beneath the connective tissue, the block observed by cryo-SEM was returned to the cryo-ultra-microtome using a cryo-transfer system. After the block was cryo-sectioned again, sectioned surface was observed by cryo-SEM again. As shown in Fig. 2b, sectioned muscle cells appeared. Repetitive sectioning and observing would be helpful to find objects localized in a limited area.

In this study it was shown that non-compressed coss-sectioned hydrated cellular and tissue architectures are clearly visualized by cryo-SEM.


We would like to thanks Ms. Ishihara A. (Leica Microsystems K. K.) for technical support. This work was partially supported by JSPS KAKENHI Grant Number, 23570196.

Fig. 1: Comparative observation by cryo-SEM and cryo-TEM. (a) Cryo-section of budding yeast observed by cryo-TEM. (b) Cryo-sectioned surface of budding yeast observed by cryo-SEM. Arrows: cutting direction, CW: cell wall, N: nucleus, V: vacuole, Mt: mitochondrion, IM:invaginated pasma membrane, Bars=500 nm.

Fig. 2: Repetitive sectioning and observation of a block surface of diagram. A piece of diagram was cryo-sectioned and observed by cryo-SEM repeatedly. (a) Sectioned diagram near the surface of the frozen block. (b) Sectioned diagram inside the frozen block. CF: collagen fibers, N: nucleus, MF: muscle fibers, Bars=1 μm.

Type of presentation: Poster

IT-7-P-1655 Characterization of Melting and Crystallization Behavior in the Au-Ge Eutectic System Using Au-catalyzed Ge Nanowires

Marshall A. F.1, Thombare S. V.1, Chan G.1, McIntyre P. C.1
1Stanford Nanocharacterization Facility and Materials Science and Engineering Department, Stanford University, Stanford, CA
afm@stanford.edu

The catalyzed growth of nanowires (NWs) can provide us with a useful platform for studying nanoscale phase transformations that are readily observed using in situ transmission electron microscopy. For example, following vapor-liquid-solid (VLS) growth of Ge and Si NWs, the re-solidified catalyst, typically Au, remains at the end of the NW, with an abrupt, planar interface between the two materials. Fundamental behaviors of these nanoscale eutectic systems, such as melting and crystallization, as well as metastable phase formation, can be studied by heating and cooling the NWs in the TEM. The use of a MEMS based heating holder (Protochips AduroTM) allows for a large range of heating and cooling rates, including quench rates that are comparable to those used in more traditional rapid quench studies. Here we present details of the melting and crystallization behavior of the metastable hexagonal close-packed beta phase of the Au-Ge eutectic system.

We have previously shown that the metastable hcp phase, which occurs following NW growth [1], can also be formed by melting and rapid quenching of the Au nanocatalyst at the tip of a Ge nanowire [2]. Fig. 1 shows the melting behavior of the quenched-in metastable phase. Melting occurs over a timeframe of seconds; it begins at the edges of the Ge NW-catalyst interface (Fig. 1a and b). In Fig. 1b melting is also visible at the top of the catalyst indicating that initial melting continues along the surface. An abrupt formation of additional stacking faults (Fig. 1c), characterizes a transition to a large volume of melt regions that form parallel to the {0001} planes of the remaining crystal (Fig. 1d). As the melt volume grows, the crystal pulls away from the surface, adopting a spherical shape, and “floats” in the liquid, then moves to the interface (Fig 1e), before abruptly dissolving into the liquid volume (Fig. 1f). This last process is accompanied by a notable darkening of the liquid as a result of mass contrast induced by the dissolved Au. We note that this same sequence of events has been observed a number of times in the hcp quenched structure, e.g. Fig 2. These results suggest that orientation of the crystal, and diffusion along the {0001} planes of the metastable hcp phase influence details of the melting process. We will also present results of cooling studies, which indicate a correlation between the formation and orientation of the metastable phase, and show that the amount of Ge that remains in the hcp structure can be controlled by the cooling rate and minimized by subsequent annealing.

[1] A.F. Marshall, et al, Nano Lett. 10 (2010), 3302. [2] A.F. Marshall, et al, Microscopy and Microanalysis 2013, Phoenix, AZ.


Acknowledgement: Financial support is provided by National Science Foundation grant DMR-1206511. This work was performed at the Stanford Nanocharacterization Laboratory.

Fig. 1: Fig. 1: Selected timeframes from a melting video of a quenched nanocatalyst with the metastable phase. Melting begins at the edges of the interface in the first frame and proceeds through a series of morphological changes until final melting at about 11 seconds.

Fig. 2: Fig. 2: Still images from another quenched in metastable structure (a) shows the same preferred melting along the <0001> direction (b), and a detached crystal within the melt (c) prior to final melting.

Type of presentation: Poster

IT-7-P-1721 Maestro: a Matlab-based centralized computer control system for an electron microscopy laboratory

Bergen M.1, Dalili N.1, Malac M.1,2, Hoyle D.3, Chen J.1, Taniguchi Y.4, Yotsuji T.4, Yaguchi T.4, Hayashida M.5, Howe J.3, Kupsta M.1
1National Institute for Nanotechnology, Edmonton, Canada, 2Univ. of Alberta, Edmonton, Canada, 3Hitachi High Tech Canada, Toronto, Canada, 4Hitachi High Tech, Naka, Japan, 5AIST, Tsukuba, Japan
marek.malac@gmail.com

A modern electron microscopy (EM) laboratory needs to integrate a number of auxiliary devices with a (transmission) electron microscope (TEM). Integration refers to controlled operation of all the hardware and the TEM, image and spectra recording at precise moments of the experiment and accurate logging of the status of the microscope and all connected devices. Here we report extensive development of a Matlab-based central computer system for an EM laboratory, referred to as Maestro [1].

Matlab has been successfully used to control TEMs [2]. The Maestro computer control system offers extensive functionality beyond the microscope control. At present, the Hitachi HF-3300 TEM / scanning TEM and H-9500 environmental TEM (ETEM) can be fully controlled. Additional devices, such as video recording software, gas handling system for ETEM, custom-built controllers for sample heating, electron biprisms, Gatan Image Filter, Gatan DigiScan, electron tomography holders are included in the Maestro system. The status of all active devices is recorded within each data set (image, diffraction pattern or spectra) and can be later reloaded to reproduce the exact instrument status. Control and logging of multiple data sources is possible.

Fig. 1 shows a generic layout of an EM laboratory controlled by Maestro. The communication between the central computer with Maestro and Matlab is typically over LAN, but non-LAN methods, such as RS232 are possible. Maestro allows operation using a control Matlab script that accurately executes an experiment with control of multiple parameters and efficiency far exceeding that of a manually operated microscope. In an ETEM, the possibility to control and log multiple experiment parameters leads to about five fold decrease of experimental time and the elimination of user errors in the experiment execution. Fig. 2 shows an example of low loss EELS trace of gas composition obtained in an H-9500 ETEM with 100 ms time resolution. Both the gas composition and data acquisition were controlled by Maestro. Maestro can be also operated through a graphical user interface (GUI) shown in Fig. 3. Often experiments are developed as a script and the GUI is implemented for frequently repeated experiments. The settings of all devices used in an experiment are saved either as tags in traditional Digital Micrograph (DM) files or as a Matlab mat file. The tags with device status can be accessed in DM, as shown in Fig. 4, as well as in Matlab. Maestro can execute existing DM scripts within a Matlab instrument control script. Maestro can be used both as point-and-click GUI tool and an accurate script based advanced instrument control.

[1] M. Bergen et. al. Micr.  Microanal. 19 S2 (2013), p. 1394

[2] TOM toolbox: www.biochem.mpg.de/278655/tom_e


The work was made possible by extensive support of Mr. I. Cotton and Hitachi High Tech, Canada and by funding from NRC/NINT and Alberta Innovates Tech. Futures.

Fig. 1: Layout of a generic laboratory controlled by Maestro

Fig. 2: EELS trace of gas composition in a Matlab-controlled environmental TEM

Fig. 3: Graphical user interface for one click connection to multiple devices

Fig. 4: Tags with device status during data acquisition viewed in Digital Micrograph

Type of presentation: Poster

IT-7-P-1777 In situ investigations of the thermodynamic behavior of dislocation loops in nanopillars and their impact on nanomechanical properties

Kiener D.1, Jeong J.2, Lee S.2, Zhang Z.3, Oh S. H.2
1Department Materials Physics, Montanuniversität Leoben, Austria, 2Department of Materials Science and Engineering, Pohang University of Science and Technology, Korea, 3Erich Schmid Institute of Materials Science, Austrian Academy of Sciences, Austria
daniel.kiener@unileoben.ac.at

Studying the nanomechanical behavior of miniaturized objects was enabled by the availability of focused ion beam (FIB) microscopes to create nanoscale structures, and boosted by unique deformation mechanisms encountered in nanoscale dimensions. Quantitative testing in situ in the TEM [1] was seminal in aiding understanding of underlying processes. However, a remaining issue concerns near surface crystal defects created by the FIB [2] and their influence on the properties of nanoscale samples [3].
Here we combine in situ heating and in situ nanomechanical TEM testing to study the thermodynamic behavior of FIB induced crystal defects in the confined volume of nanopillars, and their influence on mechanical properties on fcc, bcc, and hcp metals.
We show that during annealing, initially the FIB induced prismatic dislocation loops undergo an Oswald ripening process (Fig. 1). From this time resolved process we were able to determine the activation energy of lattice diffusion in Al, similar to recent pipe diffusion measurements along dislocations [4]. Upon further annealing to about 0.6 Tm, the remaining loops exit the sample due to image forces, leaving behind a pristine pillar as confirmed by HRTEM (Fig 2a, b). Loading these pristine samples in compression in situ in the TEM, we observed that dislocation plasticity initiates by surface nucleation of dislocations at very high stresses, much higher than required to plastically deform specimens that still contain FIB induced loops (Fig. 2c, d). This demonstrates that we can undo the FIB induced damage and restore pristine crystals and probe their intrinsic mechanical behavior. Finally, annealing the samples to even higher temperatures closer to the melting point, we could study the sublimation, melting, and evaporation processes of such confined metallic volumes. For the case of Mg, we show that the Ga from the FIB processing stimulates the sublimation process of the Mg crystal with a flat interface. This sublimation causes enrichment of Ga at the interface, leading to formation of a lower melting point Mg-Ga alloy. Upon melting the surface forms cusps, and evaporation of the molten Mg encapsulated in a MgO shell continues (Fig. 3). Notably, the rate of material loss during sublimation and evaporation did not change.
These observations underline the importance of direct in situ observation in the TEM when attempting to investigate nanoscale thermal or mechanical processes.

References:
[1] Dehm G, Howe JM, Zweck J. In-Situ Electron Microscopy. Weinheim: Wiley-VCH 2012
[2] Kiener D et al. Mater. Sci. Eng. A 2007;459:262
[3] Shim S, et al. Acta Mater. 2009;57:503
[4] Legros M, et al. Science 2008;319:1646


DK acknowledges support from the Austrian Science Fund (FWF), projects I 1020-N20 and P 25325.

Fig. 1: Images showing FIB prepared Al pillar (a) and during loop growth (b - d).

Fig. 2: HRTEM image of a Cu pillar after FIB fabrication (a) and subsequent annealing (b). Stress-strain curve of annealed (c) and FIB prepared (d) pillar.

Fig. 3: Sublimation of Mg nanopillar: (a-d) Straightening and continuous sublimation of the Mg pillar. (e) Evaporation of Mg. Note the lack of diffraction contrast and the cusps formed.

Type of presentation: Poster

IT-7-P-1812 Dislocation mediated creep/relaxation in nanocrystalline palladium thin films revealed by on-chip high resolution TEM in-situ testing

Amin-Ahmadi B.1, Colla M. S.2, Idrissi H.1,2, Malet L.3, Godet S.3, Raskin J. P.2, Pardoen T.2, Schryvers D.1
1Electron Microscopy for Materials Science (EMAT), University of Antwerp, Belgium, 2Université catholique de Louvain, Institute of Mechanics, Materials and Civil Engineering, Louvain-la-Neuve, Belgium, 3Université Libre de Bruxelles, Matters and Materials Department, Belgium
behnam.amin-ahmadi@uantwerpen.be

The high rate sensitivity of nanostructured metallic materials demonstrated in recent literature is related to the predominance of thermally activated deformation mechanisms favoured by a high density of internal interfaces. In the present study, we report for the first time in-situ high resolution transmission electron microscopy (HRTEM) creep/relaxation tests on electron beam evaporated nanocrystalline (nc) palladium (Pd) thin films using an original on-chip nanotensile method resembling the technique used in [1]. Unexpectedly, large creep/relaxation rates have been observed at room temperature. Figure 1 shows the microstructure of the as-deposited Pd films characterized in both cross-sectional and plan-view thin foils prepared by focused ion beam (FIB). In this figure, columnar nanograins can be observed with 2 or 3 grains confined over the thickness of the films with an in-plane grain diameter of ~30 nm. Automated Crystallographic Orientation Mapping in TEM (ACOM-TEM) shows a clear [110] fibre texture parallel to the growth direction. The microstructure involves Σ3 60° {111} coherent twin boundaries (CTBs) in ~ 25% of the grains.
The in-situ HRTEM characterisation of the evolution of the microstructure shows that, despite the small grain size, the creep/relaxation mechanism is mainly mediated by the stress driven thermally activated nucleation and propagation of dislocations. Interestingly, the formation and the destruction of sessile Lomer-Cottrell dislocations have been observed in-situ. Furthermore, clear loss of the coherency of CTBs with time was observed as indicated by the progressive increase of the thickness of these boundaries as shown in Figure 2. Such feature is attributed to the interaction of CTBs with lattice dislocations. The impact of these elementary plasticity mechanisms on the creep/relaxation behaviour of the Pd films is discussed and compared to recent experimental and simulation works in the literature. This constitutes a key issue in the development of a variety of micro- and nanotechnologies, such as Pd membranes used in hydrogen applications.

References
1. H. Idrissi, B. Wang, M.S. Colla, J.P. Raskin, D. Schryvers and T. Pardoen, Adv. Mater. 23 (2011), P.2119.


Fig. 1: Figure 1. a) ACOM-TEM orientation mapping of as-deposited Pd film. Corresponding inverse pole figure along different directions are shown. b) Grain size distribution of (a). (c) Bright field micrograph obtained on cross-sectional as-deposited Pd film (d) HRTEM image obtained in as-deposited films showing a Σ3 60° {111} coherent TBs.

Fig. 2: Figure 2. HRTEM images showing Σ3 {111} TBs at (a) t=0 and (b) t=3 days, respectively. Note the increase of the TBs thickness from (a) to (b) in the filtered images at the upper right insets. Corresponding Fast Fourier Transform (FFT) showing the twin character is shown in the lower right inset of each image.

Type of presentation: Poster

IT-7-P-1868 In situ mechanical testing in the transmission electron microscope and finite element method simulations on nanoscaled amorphous silica spheres – Densification, hardening and improved intrinsic properties on nanoscale

Mačković M.1, Niekiel F.1, Spiecker E.1
1Center for Nanoanalysis and Electron Microscopy (CENEM), FAU Erlangen-Nürnberg, Cauerstr. 6, 91058 Erlangen, Germany
mirza.mackovic@ww.uni-erlangen.de

As functional members in electronic devices and fiber-based telecommunication techniques oxide glasses have become subject of intense research. In addition to functionality, oxide glasses often have to fulfil mechanical reliability. Because silica is relevant in electronic and optical applications, it is chosen as a suitable model system towards a general understanding of factors which control intrinsic strength, deformation and elastic properties of non-crystalline materials. Since glasses lack of long-range periodicity, usual strengthening strategies, which comprise introduction of defects or grain boundaries with the aim to inhibit dislocation motion [1], are not working. Hence, novel approaches are appreciated to improve their mechanical properties. In the past, electron beam (EB) irradiation was used to tailor the properties of materials [2], but facing the problem of increased specimen temperature. EB irradiation is also known to densify amorphous silica (a-SiO2) on macroscopic scale [3]. Recent in situ transmission electron microscopy (TEM) studies have shown that moderate EB irradiation is very useful to induce enormous ductility in nanoscale a-SiO2 [4].

In the present study combined in situ mechanical testing in TEM and finite element method (FEM) simulations were used to characterize the mechanical properties of nanoscaled a-SiO2 spheres. First, the dose-dependent densification of a-SiO2 upon EB-irradiation was monitored in situ in TEM (Fig. 1). At low beam current doses (LD) a-SiO2 spheres densify clearly less compared to higher dose (HD) irradiation. In order to investigate the effect of EB irradiation on the mechanical properties, the spheres were irradiated with either LD or HD and then compressed under beam-off (Fig. 3a,b) or beam-on (Fig. 3c) conditions. We observe a pronounced hardening effect (Fig. 3), whereby higher loads are required to compress a-SiO2 spheres, which are treated with HD irradiation prior to compression [5,6]. FEM simulations based on an elastic / ideally plastic model (set-up in Fig. 2) reveal an increase in Young’s modulus upon HD irradiation (not shown) [6], as well as different plastic strains for beam-off and beam-on compression (Fig. 3). This clearly proves that the intrinsic glass properties can be tailored by EB irradiation [5]. Our approach is highly promising and opens opportunities for fundamental studies on structure-property relations of nanoscaled glass.

1. K. Lu et al., Science (2009) 324:349; 2. A. Krasheninnikov et al., Nature Mater. (2007) 6:723.
3. W. Primak et al., J. Appl. Phys. (1968) 39:5651.
4. K. Zheng et al., Nature Comm. (2010) DOI: 10.1038/ncomms1021
5. M. Mačković et al., Microscopy Congress MC2013, Regensburg, Proc. (Part 1), pp. 470-471.
6. M. Mačković et al., submitted.


Financial support by DFG via SPP1594 „Topological Engineering of Ultra-Strong Glasses” and the Cluster of Excellence (EXC 315) is acknowledged. The authors thank M. Hanisch and R.N. Klupp-Taylor for providing the amorphous silica spheres.

Fig. 1: Quantitative in situ observation of electron beam induced densification of nanoscaled a-SiO2 spheres.

Fig. 2: FEM simulation showing an a-SiO2 sphere compressed at maximum displacement.

Fig. 3: Electron beam hardening of nanoscaled a-SiO2. TEM images of a-SiO2 spheres after compression (top); experimental load-displacement curves and corresponding FEM simulations (center); plastic strain fields at maximum loads from FEM simulations (bottom).

Type of presentation: Poster

IT-7-P-1941 Deformation modes of Au nanowires revealed by mechanical testing in TEM

Lee S.1, Im J.1, Bitzek E.2, Kiener D.3, Oh S.1
1POSTECH, Pohang, Republic of Korea, 2Friedrich-Alexander Universität Erlangen-Nürnberg, Erlangen, Germany, 3Montanuniversität Leoben, Leoben, Austria
shoh@postech.ac.kr

Shrinking the size of metallic structures not only leads to an increase of strength (i.e. the ‘smaller is stronger’ size effect), but also to a change in the deformation mechanism. In the case of uniaxial deformation of face-centered cubic (fcc) metal nanowires, the deformation mechanism can also change with the loading condition. According to recent molecular dynamics (MD) simulations, ultra-thin Au [110] nanowires (diameters of a few nm) deform predominantly by dislocation slip in compression, but in tension by deformation twinning. Here we report, by combination of in-situ transmission electron microscopy (TEM) and molecular dynamic simulation, the conditions under which particular deformation modes take place during the uniaxial loading of [110]-oriented Au nanowires [1].

In our deformation setup in TEM, a wedge-shaped top end of Au [110] nanowire was first compressed with a flat diamond punch (Fig. 1a), thus the initial deformation was localized near the contact region. Under such a strain gradient condition, the initial compressive deformation began with the emission of small prismatic loops from the top corner (white arrows in Fig. 1b). Initially, the loops appeared to replicate the perimeter of the contact line, but after a certain number of closed loops were punched out (typically less than ten), there was a clear transition in the nucleation mechanism; open loop dislocations started to bulge out and then released from the contact area (yellow arrows in Fig. 1b-c). As the contact area increased, ordinary dislocation slip along the inclined {111} slip planes dominated the compressive deformation (Fig. 1c and d).

The deformation mode of Au [110] nanowires changes from dislocation slip to deformation twinning as the loading condition is reversed from compression to tension. Moreover, once a Au nanowire has been twinned by the initial tensile loading, the subsequent compressive deformation was carried predominantly by detwinning instead of the expected dislocation slip (Fig. 2a and b). This twinning-detwinning behavior is capable of accommodating large plastic strains (> 30%) reversibly and repeatedly over many tension-compression cycles (Fig. 2c). Molecular dynamics simulations rationalize the observed behaviors in terms of the orientation-dependent resolved shear stress, i.e. Schmid factor, on the leading and trailing partial dislocations, their potential nucleation sites and energy barriers. The present in-situ TEM results demonstrate the primary role of the loading direction in determining the governing deformation mechanisms under uniaxial loading conditions.


This work was supported by the National Research Foundation of Korea (NRF) grant funded by the Korea government (NRF-2011-0029406).

Fig. 1: a. Deformation setup of Au nanowire in TEM. b. TEM DF image showing the initial compressive deformation by prismatic loops (white arrows) and half-loops (yellow arrows). c. TEM DF image showing ordinary dislocation slip (orange arrows). d. MD simulation showing the emission of half-loops (yellow arrow) and ordinary dislocation slip (orange arrows).

Fig. 2: a. A series of TEM DF images showing reversible twinning-detwinning deformation of Au nanowire during cyclic tension-compression. b. Schematics illustrating the corresponding loading condition for each TEM image in a. c. Plot of the axial strain versus the lateral displacement during the cyclic loading.

Type of presentation: Poster

IT-7-P-1969 Observation of Strain Aging Behavior in Strain Based Line Pipe Steels using in-situ Heating and Straining TEM stage

Hong S.1, Ahn T.1, Kim S.1, Ro Y.2, Lee C.2, Kim Y.1
1Seoul National University, Seoul, Republic of Korea, 2Technical Research Laboratory, POSCO, Pohang, Republic of Korea
hong0133@snu.ac.kr

One of the new concepts for American Petroleum Institute (API) X100 grade line pipe steels was the strain-based design (SBD) approach. As demands increased for the harsh environmental applications such as the artic and seismic area, SBD line pipe steels were considered as a key solution. Even though the strength could be diminished by the processing or design, uniform elongation is the top-most property to attain in the line pipe steel. Many researchers have focused on the alloy design, combined with microstructure analysis and mechanical properties, to fabricate line pipe steel delivering both the transport efficiency and the performance. Full size X100 steel plate and pipe with 32mm thickness were selected and investigated in this study. The pipe shaping was achieved through UOE (U-ing, O-ing, and Expansion) piping process. The mechanical properties such as yield stress (YS), tensile stress (TS), and uniform elongation (uEl) were measured from the tensile test. Furthermore, microstructures were observed by scanning electron microscope (SEM) and transmission electron microscope (TEM). The dislocation structures of the plate and pipe were analyzed by selecting several layers through the thickness. Because the plastic deformation history of the surface is different from that of the center during the UOE piping process, it is expected that the dislocation density and structures were formed differently through the thickness. UOE process is typically followed by the anti-corrosion coating process, which requires heating the pipe up to 200 ~ 250°C. During the heating process for the anti-corrosion coating, the pipe reveals the strain aging phenomena giving yield drop in the stress-strain curves. To investigate both the strain and the thermal effect on the strain aging behavior of SBD X100 steels, in-situ heating and straining TEM stage was designed and applied to test the alloy. Each step of process conditions, such as applying stress and heat/cooling, was simulated in the TEM while observing the microstructural change. The analysis of strain aging behaviors was conducted.


This work was supported by the Development program(No.10040025) of the Korea Evaluation Institute of Industrial Technology grant funded by the Korea government the Ministry of Trade, Industry and Energy.

Fig. 1: TEM images of dislocation structures of the line pipe: (a) surface, (b) 1/4t, and (c) 1/2t

Fig. 2: in-situ heating and straining TEM stage: whole view (left), detail view (right)

Type of presentation: Poster

IT-7-P-2081 In Situ 3D Studies of the Chlamydomonas Chloroplast Using Cryo-Focused Ion Beam Milling and Cryo-Electron Tomography

Schaffer M.1, Engel B. D.1, Cuellar L. K.1, Villa E.1, Plitzko J. M.1, Baumeister W.1
1Department of Molecular Structural Biology, Max Planck Institute of Biochemistry, Martinsried, Germany
schaffer@biochem.mpg.de

A comprehensive understanding of eukaryotic photosynthesis, the process that converts light energy into biochemical energy, requires a molecular-resolution three-dimensional model of the chloroplast’s intricate structure. Although the first transmission electron microscopy (TEM) studies of this important organelle date back to the early days of TEM in the 1950s, these observations, and the subsequent studies in the following decades, were limited by artifact-inducing sample preparation techniques. While valuable knowledge has been gained by both freeze-fracture and conventional heavy-metal stained plastic section preparations, the three-dimensional native architecture of the chloroplast can only be visualized by cryo-electron tomography (cryo-ET) of vitreous samples. In situ cryo-ET of specific subsystems within larger eukaryotic specimens requires selected areas of vitreous material to be thinned to electron transparency (less than 500 nm). Until recently, cryo-sectioning with an ultramicrotome was the only method capable of achieving this goal. However, cryo-ultramicrotomy is a laborious and technically demanding technique, and furthermore, mechanical sectioning introduces inevitable artefacts such as compression deformations.
In this work, we show that cryo-focused ion beam (cryo-FIB) milling (1,2,3) provides an alternative method of sample preparation. As a compression-free technique for thinning vitreous material to any specified thickness, it can produce ideal artifact-free specimens for cryo-ET. We combined cryo-FIB with cryo-ET in a complete integrated cryo-workflow to obtain in situ 3D tomograms of the chloroplast within the unicellular green alga Chlamydomonas reinhardtii, the canonical algal model organism for studying photosynthesis.

References:
[1] M Marko et al., Nat Methods 4(3) (2007) p.215.
[2] A Rigort et al., PNAS 109(12) (2012) p. 4449.
[3] E Villa et al., COSTBI 23(5) (2013) p.771.


Fig. 1: Cryo-FIB preparation of Chlamydomonas reinhardtii cells. (a) An SEM image of a vitrified specimen on a TEM grid. (b) A FIB SE image of a lamella edge-on and (c) a top-down SEM image of a lamella. (d) A 2D slice from a tomographic reconstruction of a chloroplast within the intact cell, revealing the thylakoids and the chloroplast double membrane.

Type of presentation: Poster

IT-7-P-2095 In-situ Lorentz microscopy of high Bs and low core-loss Fe85Si2B8P4Cu1 nanocrystalline alloys

Akase Z.1,2, Shindo D.1,2, Sharma P.3, Makino A.3
1Institute of Multidisciplinary Research for Advanced Materials, Tohoku University. , Sendai, Japan, 2Center for Emergent Matter Science, RIKEN, Wako, Japan, 3Institute for Materials Research, Tohoku University, Sendai, Japan
akase@tagen.tohoku.ac.jp

The Fe-Si-B-P-Cu nanocrystalline alloys which exhibit excellent magnetic softness and relative high saturation magnetic flux density have been newly developed. This material has a homogeneous nanocrystalline structure composed of alpha-Fe grains with a size of about less than 20 nm which are realized by crystallizing the heterogeneous amorphous alloys. In this study, we observed the movements of the magnetic domain walls in the heat-treated Fe85Si12B6P4Cu1 amorphous-ribbons by in-situ Lorentz microscopy using a transmission electron microscope equipped with a magnetizing system, in order to understand the dependence of the magnetic properties on the microstructures.

Figure 1 shows a schematic illustration of the magnetizing system installed on a JEM-3000F instrument. The magnetizing specimen holder and two deflection coils are connected to an electric power source via three independent amplifiers. The two deflectors control the incident angle of the electron beam to avoid shifting of the image on the screen. When the amplifiers are connected to the DC source, a static external magnetic field is applied to the specimen. In order to observe the motion of the magnetic domain wall, the amplifiers are connected to an AC source.

The smooth movement of magnetic domain walls was observed in the specimen which was heat-treated at 430 °C, while the specimen which was heat-treated at 470 °C showed less-smoothness of the domain wall motions. Both of two specimens have the nanocrystalline structure in which the size of alpha-Fe crystallite is about 5 nm, but the electron diffraction pattern indicates that the latter specimen contains precipitates of boride. Figures 2a to 2d show Lorentz micrographs of the specimen which was heat-treated at 470 °C in a static external magnetic field of 3.8 kA/m, 4.3 kA/m, 4.6 kA/m and 4.7 kA/m, respectively. The direction of the external magnetic field is indicated by arrow at the top right of the Fig. 2. The positions of the magnetic domain wall are indicated by a dotted line, and previous positions of the domain wall were also plotted on the images. It is noted that a pinning of the motion of domain wall was observed at the position indicated by a white circle in Fig. 2c. It was considered that the precipitates caused the less-smoothness of the domain wall motions.


This work was supported by "Tohoku Innovative Materials Technology Initiatives for Reconstruction (TIMT)" funded by MEXT and Reconstruction Agency, Japan.

Fig. 1: A schematic illustration of the magnetizing system used.

Type of presentation: Poster

IT-7-P-2143 Observing impregnation dynamics at the liquid-solid interface using scanning electron microscopy: Charged-controlled SAPO-34 zeolite particle dispersion on SiC substrates

Tran C. M.1, Fordsmand H.1, Appel C. C.1
1Haldor Topsøe A/S, Nymøllevej 55, DK-2800 Kgs. Lyngby, Denmark
chmt@topsoe.dk

Insight into colloidal and interface processes benefits from observations made by electron microscopy. A wide palette of materials manufacturing techniques rely on drying of colloidal systems. The drying step plays a key role in distributing the suspended solid particles on a substrate and the density, clustering and packing of solid particles will intimately depend on the evaporation of the solvent. In general, drying is believed to depend on several parameters such as the surface tension of the liquid, the zeta potential of the particles and the state of the substrate. Up to now, observations of the dynamic processes during drying are scarce. Here, such investigations are reported for the dispersion of approximately 1 µm wide SAPO-34 zeolite particles suspended in liquid onto a porous outer surface of the SiC substrate, as a model system for automotive exhaust abatement catalysts. Using scanning electron microscopy (SEM) in combination with a differentially pumped vacuum system and a Peltier cooling stage, time-lapsed image series are acquired in situ during humidity variation at constant specimen temperature, whereby the dynamic arrangement of particles on the substrates is directly observed. Specifically, the effect of surface-modifications for cationic, anionic and neutrally charged particles in the suspension is shown to markedly affect the distribution on the SiC. Moreover, complementary SEM observations under cryo conditions of freeze-fractures of the fully hydrated samples are pursued to provide a snapshot of the particle distribution inside the porous SiC, Fig. 1. Differences in the arrangements of the zeolite particles in the liquid, indicate that electrostatic interactions between the charged particles and the substrate in the porous structure. These results can directly be explained by the electrostatic interaction between the SAPO-34 zeolite particles and the SiC substrate and proposes a method for guiding particle dispersions in porous support systems.


Fig. 1: SEM SE images for cationic, neutral and anionic washcoats. The zeolites are in a different arrangement depending on the charge. The color maps show where the particles (yellow) are situated compared to the SiC (black). The blue color is representing the ice.

Type of presentation: Poster

IT-7-P-2418 EBSD Measurements of the Twinning Process in an Mg-4wt%Li-Alloy with an in-situ Tensile / Compression Module in the SEM DSM 982

Fahrenson C.1, Driehorst I.1, Lentz M.2, Camin B.2, Berger D.1
1Technical University Berlin, Center for Electron Microscopy (ZELMI), Straße des 17. Juni 135, 10623 Berlin, Germany, 2Technical University Berlin, Chair Metallic Materials, Ernst-Reuter-Platz 1, 10587 Berlin, Germany
fahrenson@tu-berlin.de

The combination of a tensile / compression module in a scanning electron microscope (SEM) enables the in-situ analysis of the microstructure modifications as a function of the applied load and strain direction. To analyze the microstructure in a SEM quantitatively, an electron backscattered diffraction (EBSD) measurement is the most promising tool. Unfortunately, the information depth of the EBSD signal is very small; therefore, it needs first to be clarified first if an EBSD-signal might be recorded from stressed specimen surfaces. This investigation is the purpose of this paper.
These measurements require a large specimen chamber with the possibility containing the module and a sample stage whose loading capacity is large enough for the module. Furthermore, for EBSD measurements it is essential that the EBSD-detector can be positioned as close to the sample as possible to optimize the data collection. Additionally, shadows of the module on the detector should be minimized. We used a tensile / compression module from Kammrath & Weiss in the "Narrow version" and a Zeiss SEM DSM 982 with Gemini optic. The used EBSD-detector is a Pegasus system from EDAX with a high speed CCD camera (Hikari). All technical difficulties with the integration of the module, the shadowing and the sample alignment are described in [1].
An Mg-4wt%Li-1wt%Al alloy was investigated whose misorientation relations should be determined. The deformation behavior of magnesium alloys is significantly influenced by the activation of mechanical twinning. Hence, twin nucleation and growth will be observed through characteristic reorientations during several load steps. Figure 1 and 2 show an EBSD-measurement of an Mg-4wt%Li alloy of an unloaded sample with the initial grain orientation. Figure 3 shows the surface of the same specimenposition after a compression of 14% in horizontal direction. The surface of the specimen became quite rough indicating a strongly strained surface. Nevertheless, it is still possible to record meaningful EBSD-maps (fig. 4). Only some areas close to grain boundaries are strongly deformed. In addition, the EBSD-map reveals that the grains have a preferred orientation after compression. The change of the orientation through twinning might be observed in subsequently recorded EBSD-maps during compression. In figure 4 several grains are still twinned while others are already completely sheared.
The presented results confirm that EBSD-measurements are still possible on strongly compressed specimens and that the complete twinning process during the increasing deformation might be observed and analysed in-situ.

[1] Microscopy Conference MC2013, Regensburg


The authors would like to thank Prof. Reimers for providing access to the tensile / compression module.

Fig. 1: SEM-Image from the initial state; the marked grain is in all images the same

Fig. 2: EBSD-map from the initial state

Fig. 3: SEM-Image after 14% compression

Fig. 4: EBSD-map after 14% compression

Type of presentation: Poster

IT-7-P-2483 High resolution cryo-CLEM: from cryo-light microscopy to cryo-TEM, through cryo-milling

de Marco A.1, Mayer T.1, Mahamid J.2, Arnold J.2, Plitzko J.2
1FEI Company, Munich, Germany, 2Max Plank Institute of Biochemistry, Munich, Germany
alex.demarco@fei.com

Correlative light and electron microscopy (CLEM) aims at combining the large field of view and chemical specificity of fluorescence microscopy with the high resolution ultra-structural details revealed by electron microscopy. CLEM can be extremely powerful in extending electron microscopy analysis to rare events that are impossible to target based on EM morphology alone. If CLEM is done on frozen hydrated samples there is also the opportunity to perform structural studies of complexes in situ.

Here is presented an innovative design for a cryo-light microscopy stage, developed to acquire data in cryo-light microscopy maximizing the resolution and minimizing the contaminations typically deposited on the sample during acquisition. The proposed design is extremely simple, where the stage is immobile and an inverted microscope is moved underneath. This allows the sample to be stored at cryogenic temperature, while the microscope and the objective are kept at room temperature in order to optimize the image quality.

Once the sample has been imaged in the light microscope, if suitable, it can directly go into the TEM for cryo electron tomography or single particle data acquisition. Considering the minimal amount of contamination accumulated during imaging it can easily be used for structural studies. In case the sample is too thick to be inspected in the TEM then a thinning procedure can be performed in a cryo dual-beam (Rigort A. et al JSB 2010; Rigort A. et al PNAS 2012). Relocation of a feature of interest identified in the light microscope and the dual beam is trivial thanks to the use of a cryo-shuttle which can be hosted in predefined orientation in both the light microscope and the dual-beam, as well as the use of a dedicated software framework.


Type of presentation: Poster

IT-7-P-2779 In-situ TEM studies of the electromigration process in a single InAs nanowire

Neklyudova M.1, Zandbergen H. W.1
1Kavli Institute of Nanoscience, Delft University of Technology, Delft, The Netherlands
m.neklyudova@tudelft.nl

Electromigration (EM) is a phenomenon in which the electrical current flow of high density through a solid can lead to intensive atomic motion due to the high speed electrons transfer part of the momentum to the atoms (or ions) by collision. This phenomenon can lead to morphological and structural instabilities not only in metallic interconnections but also in semiconductor nanowires [1,2]. Since semiconductor nanowires are the subject of active study in virtue of their usage as low-dimensional systems, as building blocks for future nanoscale circuits [3], the EM becomes the key issue that controls the lifetime and stability of a nanoscale device.
In this work the process of EM in a single InAs nanowire was investigated by in situ TEM technique using a FEI Titan microscope operating at 300 keV. The EM experiments were carried out in a bias-ramping mode which allowed to perform accelerated experiments for EM process visualization in-situ TEM. The voltage applied for all cycles of EM experiments was set to 1200 mV. The resistivity calculated for the nanowire diameter 221 nm was 2*10-2Ω·cm. The current density for EM activation was about 3.6*104A/cm2. It was found that the EM in InAs nanowire starts at a position close to the cathode with formation of the cubic-shaped nanoparticles in the place of failure. The EDX analysis of the nanowire after EM experiments showed that the particle formed near anode part is indium. In the presentation all structural and chemical evaluations of the InAs nanowires during the electromigration will be discussed.

[1] D. Kang, T. Rim, C.-K. Baek, M. Meyyappan. Appl. Phys. Let. 103, 233504 (2013).

[2] C.-X. Zou, J. Xu, X.-Z. Zhang, X.-F. Song and D.-P. Yu. Journal of Appl. Phys. 105, 126102 (2009).
[3] Law, M., Goldberger, J., Yang, P. Annu. Rev. Mater. Res. 2004, 34, 83–122.


I would like to acknowledge ERC project 26792.

Fig. 1: Snapshots from the real-time TEM movie showing the first cycle of EM in InAs nanowire. (a) Initial configuration of InAs nanowire before EM experiments. (b)-(e) Images of the nanowire part pointed by blue square on (a) and taken at B- E times on I-V curve respectively. (f) Typical I–V curve. The red arrows indicate the bias-ramping direction.

Fig. 2: Snapshots from the real-time TEM movie showing the 2nd EM cycle (a) Nanowire configuration after 1st EM cycle (b)-(i) TEM snapshots taken at B-I times on I-V curve respectively (j) TEM image of the nanowire part marked by green square on (i) (k) Magnified image of the nanowire part marked by red square on (j). (l) I–V curve for the 2nd EM cycle

Type of presentation: Poster

IT-7-P-2824 Understanding mechanisms of assisted sintering through dedicated in situ TEM experiments

van Benthem K.1
1University of California, Davis 1
benthem@ucdavis.edu

Sintering describes the densification of powder agglomerates through elimination of “empty space” between individual particles. [1] The application of electrical fields, currents and/or pressure in addition to heating can enable the accelerated consolidation of materials. While electric field assisted sintering, which includes spark plasma sintering and flash sintering, is already employed for the synthesis of a wide variety of microstructures with unique macroscopic properties, a fundamental understanding of the atomic-scale mechanisms that lead to enhanced densification is mostly absent from the literature. In this presentation, recent in situ transmission electron microscopy experiments will be reported that were designed to investigate specific densification mechanisms, including surface cleaning effects, i.e., dielectric breakdown of insulating surface oxides [2], mechanical properties of individual ceramic powder agglomerates [3], and electric field effects on the densification of yttrium-stabilized zirconia.

To quantitatively evaluate densification behavior we have developed an image processing tool to obtain three-dimensional densification curves from powder agglomerates. A variety of in situ TEM experiments was used to electrically contact individual nanoparticles (Figure 1a), apply mechanical pressure to particle agglomerates (Figure 1b), or expose particle agglomerates to electrical fields in non-contact mode. The results reveal that dielectric breakdown of insulating surface oxides on nanometric metal particles causes retardation of densification, while the morphology of ceramic powder agglomerates can limit densification through stabilization of pores. The application of electrical fields during in situ sintering experiments in the TEM reveals that the field strength in the absence of current has an appreciable influence on the densification behavior of Y-stabilized ZrO2. Moreover, the application of electrical fields promotes the formation of coincident site lattice grain boundaries and, hence, can accelerate grain growth in ceramic microstructures.

References

[1] Castro R, van Benthem K. Sintering: Mechanisms of Convention Nanodensification and Field Assisted Processes. Heidelberg: Springer, 2013.

[2] Bonifacio C, Holland TB, van Benthem K. Evidence of surface cleaning during electric field assisted sintering. Scripta materialia 2013;69:769.

[3] Rufner J, Holland TB, Castro R, van Benthem K. Mechanical properties of individual MgAl2O4 agglomerates and their effects on densification. Acta mater. 2014;69:187.


This work was supported by the University of California Laboratory Fee Program (12-LR-238313) and the Army Research Office (program manager: Dr. S. Mathaudu) under grant W911nf-12-1-0491-0.

Fig. 1: (a) Ni nanoparticles were contacted with an STM tip mounted on the TEM specimen holder to apply a local electrical bias [2]. (b) MgAl2O4 nanoparticle agglomerates were compressed by in situ nanoindentation [3]. Figures reproduced with permission.

Type of presentation: Poster

IT-7-P-2853 Correlating internal structure and mechanical properties of amorphous silica and gold micro-/nanoparticles using in situ mechanical testing in the scanning and transmission electron microscope

Herre P.1, Paul J.1, Romeis S.1, Niekiel F.2, Mačković M.2, Spiecker E.2, Peukert W.1
1Institute of Particle Technology (LFG), University of Erlangen-Nuremberg, Erlangen, Germany, 2Center for Nanoanalysis and Electron Microscopy (CENEM), University of Erlangen-Nuremberg, Erlangen, Germany
patrick.herre@fau.de

The ongoing miniaturization and reliability of functional materials and devices at micro- and nanoscale inevitably necessitates information on small scale mechanical properties. In the fields of e.g. optics and biomedical imaging, nanospheres of both, silica and gold, as well as hybrid core-shell structures made of these materials are of great relevance [1,2]. In order to elucidate deformation mechanisms and related mechanical properties on the nanoscale, compression of individual particles is particularly suitable due to the absence of strain gradient plasticity effects [3].

In the present work we use in situ mechanical testing in the scanning electron microscope (SEM) and transmission electron microscope (TEM) with the aim to characterize the mechanical properties of amorphous silica and gold micro-/nanoparticles. Thereby a custom built SEM supported manipulation device [4] and a TEM Picoindenter® (Hysitron, Inc.) are used.

Silica spheres are synthesized according to the Stöber-Fink-Bohn method [5] followed by thermal treatments at 400°C (S400), 800°C (S800) and 1000°C (S1000), respectively. Structural changes and mechanical properties are studied using in situ SEM indentation (Fig. 1b), combined with solid state nuclear resonance spectroscopy (NMR) and infrared spectroscopy. Reduced Young’s modulus and hardness of untreated silica particles (S70) (29.6±6.2 GPa and 6.0±0.6 GPa, respectively) increase after calcination at 1000°C (58.2±8.8 GPa and 13.4±0.9 GPa, respectively), as shown in Fig. 2b. The values achieved after calcination at 1000°C agree well with values known for bulk fused silica. In the course of heat treatment, infrared and NMR spectra (Fig. 3) reveal condensation of internal OH-groups and enhanced cross-linking of the silica network, resulting in particles with chemically and mechanically similar properties when compared to their bulk counterpart [6].

Gold nanoparticles (20 nm – 1 µm in size) are prepared by solid/liquid state dewetting of thin gold films. In situ SEM compression experiments show large strain bursts after elastic loading. Currently, cross-sections of as-prepared and deformed Au particles are investigated by ex situ TEM to get more insights into the internal particle structure. Nanomechanical testing in the TEM is focused on smaller Au particles with the aim to reveal deformation mechanisms in situ.

[1] Oldenburg et al., Chem. Phys. Lett. 288 (1998):243.
[2] Fuller et al., Biomaterials 29 (2008):1526.
[3] Gao et al., J. Mech. Phys. Solids 47 (1999):123.
[4] Romeis et al., Rev. Sci. Instrum. 83 (2012):095105.
[5] Stöber et al., J. Colloid. Interface Sci. 26 (1986):62.
[6] Romeis et al., Part. Part. Syst. Charact. (2014).


The German Science Foundation is gratefully acknowledged for financial support within the priority program “Particles in Contact” and the research training group 1896.

Fig. 1: a) Cumulative particle size distribution for untreated (S70) and heat treated (S400, S800, S1000) silica particles. Insets show representative SEM images for i) S70 and ii) S1000. b) Experimental setup for compression of single silica spheres between a diamond flat punch and a silicon substrate inside a SEM.

Fig. 2: a) Representative force-strain curves for silica particles from samples S70, S400, S800 and S1000. b) Reduced Young’s modulus, hardness and yield strength with respect to calcination temperature. At 1000°C, E* and HCEB approach the bulk values of fused silica.

Fig. 3: a) Infrared spectra of the heated silica particles. For higher temperatures densification and dehydroxylation occur. b) 29Si HPDEC and 29Si CP MAS NMR spectra of the samples. The mean number of siloxane (Si-O-Si) bonds per silicon atom (Qi, 1≤i≤4) increases from S70 to S1000, confirming enhanced cross-linking of the silica network.

Type of presentation: Poster

IT-7-P-2870 Revealing dislocation activities and deformation behavior in Nb2AlC using in situ nanoindentation in the transmission electron microscope

Schrenker N.1, Kabiri Y.1, Mueller J.1, Mackovic M.1, Spiecker E.1
1Center for Nanoanalysis and Electron Microscopy (CENEM),Erlangen, Germany
nadine.schrenker@ww.stud.uni-erlangen.de

MAX phases are layered crystals with ternary or quaternary chemical composition. Due to their excellent electrical and thermal conductivity, as well as high oxidation resistance, they are in focus of intense research activities. Dislocation activities in MAX phases at room temperature (RT) are believed to be limited to slip along basal planes. In cyclic stress-strain curves, fully reversible, rate independent and closed hysteresis loops are observed. These features are attributed to the formation and annihilation of incipient kink bands (IKBs) and dislocation walls (DWs) [1]. Fig. 1 illustrates the formation of a KB. Initiated by elastic buckling above a critical maximum shear stress dislocation pairs of opposite sign form and move in opposite direction. It is believed that a nucleated KB immediately extends to the free surface. Hence, the attraction force between DWs disappears and a kink boundary is formed. An IKB does not dissociate into mobile DWs and thus is reversible, when the load is removed. KBs were observed ex situ by transmission electron microscopy (TEM) in single crystal Ti3SiC2, after nanoindentation perpendicular to basal planes (Fig. 1e) [2]. However, to date, the precise nucleation mechanism of IKBs and DWs is not known.
By means of in situ indentation in the TEM we reveal dislocation activities in Nb2AlC [3]. Undeformed Nb2AlC specimens exhibit basal plane dislocations with 1/3<11-20> type Burgers vectors. In situ indentation in dislocation-free regions and parallel to the basal planes (Fig. 3) reveals that basal plane dislocations nucleate and move in the same slip system without cross-slip or entanglement (Fig. 4). This confirms that these dislocations are mobile at RT, as proposed by Farber et al. in Ti3SiC2 [4]. The strain bursts in the load-displacement curve (Fig. 2) are assumed to be caused by dislocation nucleation. Indentation perpendicular to the basal planes results in an elastic deformation response, followed by fracture. Currently the formation mechanisms of IKBs and DWs are investigated in situ. Furthermore, plastic anisotropy is investigated by comparing pillar compression in the µm- and nm-range using scanning electron microscopy (SEM) and TEM, respectively. Prior to compression the pillar orientation is determined by electron backscatter diffraction with SEM and electron diffraction with TEM. For compression edge-on to the basal planes it is assumed that KB formation is more likely to occur than in pillar compression parallel to the c-axis.
[1] Barsoum et al., Nat. Mater. 2 (2003): 107
[2] Molina-Aldareguia et al., Scr. Mater. 49 (2003): 155
[3] Kabiri, Master Thesis, University Erlangen-Nuremberg (2013).
[4] Farber et al., J. Amer. Ceram. Soc. 81 (1998): 1677
[5] Barsoum et al., Metall. Mater. Trans. A 30 (1999): 1727


Financial support by the DFG via research training group GRK 1896 is gratefully acknowledged. The authors further thank Prof. Dr. Peter Greil for providing the samples.

Fig. 1: Schematic of kink band formation: (a) Elastic buckling, (b) corresponding shear diagram, (c) formation of dislocation pairs and (d) kink band [5], (e) Cross-sectional TEM image of an indent in a Ti3SiC2 (0001) single-crystal thin film with a maximum load of 40 mN [2].

Fig. 2: Load-displacement and time-displacement curve of an in situ indentation parallel to the basal planes. Start and finish of the indentation corresponds to points 1 and 2 or 3, respectively [3].

Fig. 3: TEM image showing the sample tilted in two beam condition around the <0001> zone axis, before an in situ indentation parallel to the basal planes [3].

Fig. 4: TEM image after an in situ indentation parallel to the basal planes revealing nucleation and propagation of basal plane dislocations [3].

Type of presentation: Poster

IT-7-P-2887 In situ reduction of graphene oxide by Joule heating with TEM-STM system

Martín G.1, Claramunt S.1, Varea A.1, Yedra L.1,2, Rebled J. M.1,3, Sánchez-Hidalgo R.4, López-Díaz D.4, Velázquez M. M.4, Cirera A.1, Peiró F.1, Estradé S.1,2, Cornet A.1
1MIND/IN2UB, Departament d’Electrònica, Universitat de Barcelona, Marti i Franqués 1, 08028 Barcelona, Spain, 2CCiT, Scientific and Technological Centers, Universitat de Barcelona, C/Lluís Solé i Sabaris 1, 08028 Barcelona, Spain, 3Institut de Ciència de Materials de Barcelona-CSIC, Campus UAB, 08193 Bellaterra, Spain, 4Departamento de Química Física, Facultad de Ciencias Químicas. Universidad de Salamanca, E37008 Salamanca, Spain
gmartin@el.ub.es

Graphene has attracted a great deal of interest from scientists due to its intrinsic mechanical, thermal and electrical properties [1], [2]. Graphene, one-atom-thick layer of carbon, is a semiconductor with zero band gap [3] and high intrinsic mobility [4]. The excellent properties of graphene [5] have driven the search for methods for its large-scale production.

Graphene can be prepared by various methods [6] including micromechanical cleavage, epitaxial growth, chemical vapour deposition, exfoliation using graphite intercalation compounds and oxidation-reduction methods [7], [8]. These methods render high-quality graphene flakes although its low productivity makes them unsuitable for large-scale applications. The alternative strategy is the chemical oxidation of graphite or different carbon materials followed by chemical or thermal annealing.

Although the chemical oxidation of graphite is considered one of the most attractive methods to obtain graphene because it is cheaply, scalable and versatile, it presents the disadvantage that the O-containing groups produced by chemical oxidation, which make graphene oxide (GO) non-conducting [9], cannot be completely removed by the thermal annealing reduction. Thus, the level of reduction of GO is directly related to the conductivity, which can increase several orders of magnitude through the reduction process [10], [11].

In this work, GO, produced using a slight modification of the Hummers oxidation method from natural graphite flakes [12], has been in situ reduced by Joule heating in a TEM with a STM holder. The reduction of GO has been measured qualitatively from the comparison of conductivity of the sample before and after the reduction, all in the same experiment. Besides, with this technique it is possible to control the reduction from the measure of the conductivity of the sample and also characterize the sample during the experiment (both through TEM observation and through I-V characteristic). Indeed, the results show how GO has been reduced by observing a decrease of the resistance of more than four orders of magnitude.

[1] K.Novoselov et al, Science 306, 666-669(2004)
[2] A.K.Geim, Science. 324, 1530-1534(2009)
[3] Y.Zhang et al., Nature 459, 820-823(2009)
[4] K.I.Bolotina et al., Sol State Com 146, 351–355(2008)
[5] Y.Zhu et al., Adv. materials 22, 3903–3958(2010)
[6] D.Galpaya et al. Graphene 1, 30(2012)
[7] F.Bonaccorso et al. Materials today 15, 12(2012)
[8] Novoselov, K. S., et al. PNAS 102, 10451(2005)
[9] I.Jung et al. Nano Lett., 8 (12), 4283(2008)
[10] C.Gómez-Navarro et al. Nano Lett., 7 (11), 3499(2007)
[11] A.Bagri et al., nature chem. 2, 581(2010)
[12] B.Martín-García et al., ChemPhysChem 13, 3682(2012)


Fig. 1: TEM image of the tip contacting the GO during the experiment.

Type of presentation: Poster

IT-7-P-2942 Deformation Behavior of Silica Microparticles under Electron Beam Irradiation

Stauffer D.1, Bhowmick S.1, Major R.1, Asif S.1, Warren O.1
1Hysitron, Inc.
sanjit@hysitron.com

 

The studies of irradiation damage in silica are of significant interest because of its application in nuclear reactors, nuclear waste containers, optical fibers, and semiconductor devices. In this work, we investigate plastic flow and failure behavior of amorphous silica particles (1050±30 nm) under compressive stress inside a scanning electron microscopy (SEM). In situ quasistatic compression experiments were conducted using a PI 85 SEM PicoIndenter (Hysitron, Inc., Minneapolis, MN) with 2.5 mm flat punch diamond probe inside an SEM.

The deformation behavior of the particles before and after the experiments with beam on and beam off conditions can be seen in figures 1a-d. A large variation in the total plastic strain and tendency to fracture has been observed which varies with peak loads and beam condition. Here, plastic strain has been calculated as the ratio d/D, where D is the diameter of the particle and d is the amount of compression along the indentation axis. In quasistatic experiments with a 190s hold at 1 mN peak load, a particle deformed plastically to 55.5% strain when the beam was kept on during the test (fig. 1a). However, when the beam was turned off (fig. 1b), a similar diameter particle showed negligible strain (<0.05%). When the peak load is increased to 4 mN peak load with the beam on, a plastic strain of 57.8% strain was found with a crack that appeared on the surface as marked in fig 1c. In beam off condition, a similar sphere deformed plastically to 37% strain, occurring in conjunction with a large fracture which created a wedge-shaped missing segment as observed in figure 1d. The results in this study can be explained with the structural changes of the particles that has been reported in the literature. It has been observed that electron beam with sufficient intensity can change the pore structure of amorphous silica where small pores shrink and larger pores expand. The change in pore structure leads to softening of the particles which causes viscous fluid-type deformation. However, it should be emphasized that all the particles used in this study were exposed to electron beam before testing. So, it can be assumed that irradiation induced damage or defects in all the particles before loading were similar. This leads to a conclusion that the applied stress on the particles is playing a significant role in enhancing the structural changes and/or inducing more defects when electron beam is kept on. An important implication of this study is that electron irradiation under applied stress can induce significant instability and reduction in strength in silica resulting in lower lifetime in many devices where silica is an integral component.


Fig. 1: a-d: Images of deformation behavior of silica particles after quasistatic compression experiments with beam on and beam off conditions 1 mN and 4 mN. Fig e-f: Load-displacement plots at Pmax= 1 and 4 mN shows the effect of electron beam on plastic flow of the materials.

Type of presentation: Poster

IT-7-P-2970 Light Irradiation of ETEM Samples for In-Situ Studies of Photocatalysts

Miller B. K.1, Crozier P. A.1, Zhang L.1
1Arizona State University, Tempe AZ, USA
benmiller002@aol.com

Inorganic photocatalysts are currently being intensely studied for their potential use for the production of fuels from H2O and CO2. Designing new efficient photocatalysts requires an increased understanding of the link between catalyst microstructure and activity. Environmental TEM (ETEM) is a promising technique for elucidating this link. However, while gaseous environments and variable temperatures are common to ETEM work, illumination of the sample by visible, ultraviolet, and infrared light is much less common.
We have installed a variable wavelength light source to irradiate the sample area in an FEI Tecnai F20 ETEM [1]. This will allow detailed analysis of the interaction between light and photocatalysts under reaction conditions. The current design, as seen in Figure 1, consists of a broadband light source with filters, optical fibers with a vacuum feedthrough, and a manipulator to precisely position the fiber tip with respect to the TEM sample in the microscope. The Energetiq® light source we use is a xenon lamp which is powered by an infrared laser, rather than the standard arc discharge, providing a smaller and brighter source. As seen in Figure 2, the broadband light source is capable of illuminating the sample with high intensity over a broad range of wavelengths. Optical filters may be used to specify smaller pass-bands. The measured intensity distribution at the sample position is sharply peaked at the center, reaching a maximum of about 1400 mW/cm2. This is more intense than the typical solar irradiance on the Earth’s surface which is only about 100 mW/cm2. The area of the sample illuminated with at least 90% of the peak intensity is an ellipse about 200x400 μm in size. This is large enough to cover 8 grid squares in a 200 mesh TEM grid. As shown in Figure 3, the fiber comes into the TEM at 90° to the sample rod, and the tip is cut at an angle, which refracts the light up toward the TEM sample. The angle chosen for this design was 30°, in order to sufficiently refract the light exiting the fiber while avoiding total internal reflection at the tip. This configuration, with the fiber independent of the sample rod, gives flexibility in the choice of sample holders, allowing other in-situ capabilities simultaneous to the light illumination. We are using this new capability to study the structure of titania-based nanostructured photocatalysts, and have observed changes in the surface of the titania when exposed to water and UV irradiation [2].

References:

[1] Miller, B. K.. and Crozier, P. A. Microscopy and Microanalysis 19, 461-469. (2013).
[2] Zhang, L. and Crozier, P. A. Nano Letters 13, 679-684. (2013).


The support from US Department of Energy (DE-SC0004954) and the use of ETEM at John M. Cowley Center for High Resolution Microscopy at Arizona State University is gratefully acknowledged.

Fig. 1: Overview of the light illumination system, showing the laser driven broadband light source, optics, and fiber manipulator. The monochromatic laser is used to power a xenon plasma which produces the broadband light.

Fig. 2: Light characterization, both spatial and spectral. The spatial distribution shows the intensity variation at the sample position. A faint dashed circle indicates the size of a 3mm TEM sample. The spectral distribution of the light is compared to the solar spectrum incident on the Earth’s surface.

Fig. 3: Cutaway view of the fiber as it extends into the pole piece gap. The fiber is supported at a 90° angle to the sample rod. The fiber tip is positioned close enough to the sample to provide maximum intensity without interfering with sample tilt. The tip is cut at an angle to refract light up toward the sample.

Type of presentation: Poster

IT-7-P-3047 Reversible In-Situ TEM Electrochemical studies of Fluoride Ion Battery

Chakravadhanula K. V.1, Fawey M. H.2, Kübel C.1,2,3, Scherer T.2,3, Rongeat C.2, Munnangi A. R.1,2, Fichtner M.1,2, Hahn H.1,2
1Helmholtz Institute Ulm for Electrochemical Energy Storage (HIU), Albert-Einstein-Allee 11, 89081 Ulm, Germany, 2Institute of Nanotechnology (INT), Karlsruhe Institute of Technology (KIT), Hermann-von-Helmholtz-Platz 1, 76344 Eggenstein-Leopoldshafen, Germany, 3Karlsruhe Nano Micro Facility (KNMF), Karlsruhe Institute of Technology, Hermann-von-Helmholtz-Platz 1, 76344 Eggenstein-Leopoldshafen, Germany
cvskiran@kit.edu

New research directions in Li-ion batteries are focusing on improvements of battery performance. Alternative technologies are investigated based on different chemistries using, e.g., sodium, magnesium or chloride as charge transfer ions in secondary batteries. Batteries based on a fluoride ion shuttle (fluoride ion battery) are an interesting alternative to Li-ion batteries as they can theoretically provide substantially higher volumetric energy densities compared to Li-ion batteries. Recently, the principle of a secondary battery based on a fluoride ion shuttle has been demonstrated [1]. Here, the electrolyte is one of the key components to obtain good cycling properties (e.g., resulting from fast F- conduction in fluoride ion batteries)[2].

For performing in-situ electrochemical studies, the stability of the components towards the electron beam (with beam energy and beam current being critical parameters) is essential to clearly interpret the results for the battery system in terms of the electrochemical performance. In the case of the F- batteries, the components besides being stable under the electron beam do not require an inert transfer, thus being suited as a good model system for in-situ electrochemical studies inside the TEM.

Ball milling of a mixture of (1−y)LaF3 and yBaF2 was employed to prepare La0.9Ba0.1F2.9. Initially, the electrolyte (La0.9Ba0.1F2.9) was studied for its structure, composition, porosity and stability under the electron beam. The cathode material based on a mixture of Bi (active material), La0.9Ba0.1F2.9 (ionic conductivity) and C (electronic conductivity) was prepared. Both materials were pressed to form a pellet. A lamellae of 60X35µm was prepared and electricaly contacted on the Aduro Electrochemical device (E-AEK11 from Protochips Inc.) inside the focused ion beam system (FEI Strata 400S). An Aduro sample holder from Protochips Inc. along with a Keithley 2611 sourcemeter in the FEI Titan 80-300 TEM were used in this work. SAED and HRTEM studies indicated the formation of a BiF3 phase in the cathode (reflections corresponding to d-values of 5.85Å(100)BiF3 and 3.37Å(111)BiF3, which were absent in the as-prepared state). The electrolyte structure at the interface to the cathode also changed during charging, where reflections corresponding to La were observed, indicating local reactions in the electrolyte leading to the formation of a La/LaBaF3/BiF3 cell. During discharging, most of the BiF3 was again reduced indicating the reversible behavior of the battery system in the TEM.

[1] M. Anji Reddy, M. Fichtner, J. Mater. Chem 21 (2011), p17059.

[2] C. Rongeat, M. Anji Reddy, R. Witter, et.al., ACS Applied Materials and Interfaces, 6 (2014) p2103.


Robby Prang is acknowledged for discussions towards sample preparation.

Fig. 1: (A)Thin lamella on MEMS device through FIB preparation. HRTEM micrograph and SAED pattern of cathode (B),(E) at 0V and (D),(G) at 3V respectively. (C)I-V curve during charging. (F)Line profiles of diffraction patterns showing intensive peaks of Bi and BaF3 cleaving at 3V, as evidence of structural change in the electrode, through formation of BiF3.

Type of presentation: Poster

IT-7-P-3015 In situ applications of quantitative magnetic TEM imaging in magnetic nanostructures

Rodríguez L. A.1, Magén C.1, Snoeck E.2, Gatel C.2, Marín L.1, Serrano-Ramón L.1, Prieto J. L.5, Muñoz M.5, Ortolani L.6, Algarabel P. A.3, Morellón L.1, de Teresa J. M.3, Ibarra M. R.1
1LMA-INA, Universidad de Zaragoza, Zaragoza, Spain, 2CEMES-CNRS, Toulouse, France , 3ICMA, Universidad de Zaragoza-CSIC, Zaragoza, Spain, 4ISOM, Universidad Politécnica de Madrid, Madrid, Spain, 5IMM, CNM-CSIC, Madrid, Spain, 6IMM, CNR Bologna, Italy
cmagend@unizar.es

Magnetic imaging TEM techniques such as Lorentz Microscopy (LM) and Electron Holography (EH) are powerful tools to extract valuable quantitative information with nanometer-range resolution on the local magnetic states of nanomaterials. We usually study magnetic nanostructures at room temperature and at remanent state, but the use of special TEM specimen holders and/or set-ups opens the possibility to explore in situ the evolution of magnetization states upon the application of external stimuli such as temperature changes, magnetic fields or electric currents.

In this work, we present different applications of in situ LM and EH experiments under different scenarios: the in situ application of magnetic field with a calibrated objective lens is achieved by a smart control of the tilt angles of a double-tilt holder. A mathematical procedure has been developed to determine and/or quantify the in-plane component of the applied magnetic field. This capability is illustrated with two applications: the analysis of the domain conduit properties of magnetic nanowires by measuring the nucleation and propagation (depinning) fields [1], as shown in Figure 1(a); and the accurate determination of the magnetic hysteresis loops in nanoscaled magnetic tunnel junctions (MTJs) shown in Figure 2(b), where the different orientation of the magnetic induction component normal to the electron beam with respect to the induction in the sample’s plane can be quantified and corrected [2]. The potential of cryogenic conditions to study magnetism in nanostructures (down 100 K) is demonstrated by the investigation of the magnetic properties of La0.67Ca0.33MnO3 manganite thin films [2] and the strain effects on the suppressed ferromagnetism observed, an example of this characterization is displayed in Figure 2. Finally, the use of a dedicated two-contact TEM holder for the injection of spin-polarized currents in Py nanowires, illustrated in Figure 3, is applied to the investigation of current-induced domain wall manipulation phenomena.

[1] L. A. Rodríguez et al., Appl. Phys. Lett. 102, 022418 (2013).

[2] L. A. Rodríguez et al., Ultramicroscopy 134, 144-154, (2013).


This work was supported by the Spanish MINECO Projects MAT2009-08771, MAT2011-28532-C03-02 and MAT2011-28532-C03-03. The authors acknowledge the European Union under the Seventh Framework Programme under a contract for an Integrated Infrastructure Initiative Reference 312483-ESTEEM2.

Fig. 1: (a) Depinning DW processes in CoFe nanowires by application of parallel and transversal magnetic fields (yellow arrows indicate the magnetic field directions), (b) Hysteresis loops in a Fe/MgO/FeV MTJ.

Fig. 2: Amplitude, magnetic phase shift (MAG) and magnetic flux (B) recorded by Electron Holography at low temperature (100 K) of (a) full magnetized and (b) with a superficial non-ferromagnetic layer (NFL) epitaxial La0.67Ca0.33MnO3 thin films grown on SrTiO3 substrates.

Fig. 3: Defocused LM images recorded before and after injecting an electrical current pulse of an amplitude of 1 mA and a duration of 100 μs. Red arrows point a domain wall which is propagated after the pulse.

Type of presentation: Poster

IT-7-P-3115 In-situ heating using MEMS devices on FIB/SEM systems

Novák L.1, Vystavěl T.1, Faber P.2, Mele L.2, Šesták J.1
1FEI Company, Podnikatelská 6, 612 00 Brno, Czech Republic, 2FEI Company, Achtseweg Noord 5, 5600 KA Eindhoven, The Netherlands
libor.novak@fei.com

Introduction

Information on the kinetics of microstructural evolution is important in materials science fields like recrystallization, grain growth and phase changes. This requires reliable discrimination of differently oriented crystallites or different crystal phases, coupled with useful spatial resolution, temporal resolution and temperature change rate. Currently available SEMs have spatial resolution below 1 nm, temporal resolution below 10 ms (100 Hz frame rate), but existing heating holders only allow heating bulk samples up to 100°C per minute (~2°C/s). This prohibits experiments like quenching of metals and the long ramping time may cause the sample to change (oxidize, recrystallize) before the temperature range of interest is reached. In addition the backscatter (grain-, phase-) contrast is deteriorated because solid state detectors are blinded by the infrared radiation from the sample.

As a solution to these problems we present a MEMS heating holder [1], [2] in combination with in-situ sample preparation using a DualBeam FIB/SEM.

 

Sample preparation

A chunk of material is cut with the FIB and attached to the micromanipulator needle using beam-induced deposition (Figure 1a). After lift-out it can be further shaped using the FIB (1b). It is then placed on the MEMS heating holder, fixated with beam-induced deposition and cut loose from the needle (1c).

 

Ramping rates

The tiny thermal mass of the MEMS heater and sample allow temperature changes of 1000°C in just 50 ms (2·104°C/s) for a Cu sample of 20x50x50 um3 size (including settling to within 20°C, both for heating and cooling).

 

Imaging

The small heated area of the MEMS heater reduces the infrared radiation sufficiently that solid state detectors such as in-chamber BSE detectors and EBSD cameras can be used at elevated temperatures. Figure 2, for example, shows the melting of gold micro-particles at 1064°C imaged with the solid state BSE detector.

 

References:

[1] L. Mele et al. “A molybdenum MEMS microhotplate for high-temperature operation”, Sensors & Actuators: A. Physical, 2012 | 188 | 173-180

[2] B. Morana et al. “A silicon carbide MEMS microhotplate for nanomaterial characterization in TEM”, Micro Electro Mechanical Systems (MEMS), 2011 IEEE 24th International Conference on, 23-27 Jan. 2011, 380-383


This work was supported by Technology agency of the Czech Republic, project no. TE01020118 (Competence centre: Electron microscopy).

Fig. 1: Figure 1: Sample preparation from bulk sample: extraction using ion beam and manipulator (a); shaping of sample on manipulator needle (b); placement on MEMS heating holder (c). Horizontal field width is 50µm.

Fig. 2: Figure 2: Solid State Detector BSE imaging of gold particle solidification (a → b) and re-melting (b → c). Horizontal field width is 10µm.

Type of presentation: Poster

IT-7-P-3273 Atomic Level In-situ Characterization of NiO-TiO2 Photocatalysts under Light Irradiation in Water Vapor

Zhang L.1, Crozier P. A.1
1Arizona State University, Tempe, USA
liuxian.zhang@asu.edu

Photocatalysts have potential applications for solar fuel generation either through water splitting. It is now recognized that atomic level in situ observations are critical for understanding the structure-reactivity in photocatalysts in the presence of reactant and product species and during in-situ light illumination. NiO loaded semiconductor photocatalysts with Ni first reduced and then partially re-oxidized at the surface has been reported to have good photocatalytic properties by forming a metallic Ni ohmic contact between NiO and the semiconductors [1]. TiO2 is a promising photocatalyst which has attracted intense research interest for decades since photo-decomposition of water by TiO2 was discovered. The TiO2 photocatalysts are either anatase or rutile which has been well known. Herein we use anatase as a model material to develop in situ photocatalytic experimental methodology and explore structure changes of NiO/semiconductor photocatalysts. In-situ heat treatment in H2 or O2 is applied to prepare initially Ni/TiO2, NiO/TiO2 or NiO-Ni-TiO2 materials in an environmental transmission electron microscope (ETEM). Then, without exposure to air, analysis can be performed in the same modified ETEM under in situ conditions in the presence of light and reactants to explore oxidation/reduction or interface changes under photocatalytic reactions.
NiO-Ni-TiO2 was prepared using Ni(NiO3)2 as the precursor following impregnation, calcination, reduction and partial re-oxidation. Ex-situ experiments were performed to achieve preliminary observations under exposure of xenon lamp with mirror reflecting light in the range 360nm to 460nm light. TEM images for ex-situ experiments were recorded with a FEI aberration corrected Titan TEM. Figure 1A&1B show initial Ni-NiO core-shell structures on anatase particles. The inside rounded darker particles are Ni metals with outside shells of polycrystal NiO. After 6 hrs exposure to light in liquid water the oxide shells become porous and the Ni metal is absent leaving a void (Figure 1C&1D). Ni may either be oxidized to NiO or dissolved into the solution during photocatalytic reactions.
In-situ heat treatments using a hot stage sample holder with H2 or O2 allows Ni or NiO to be prepared as the starting material for in situ photocatalytic experiments. A FEI Tecnai F20 ETEM was modified to allow samples to be illuminated with light from a broadband laser driven light source (EQ-99, Energetiq Inc.) with the intensity up to 10 suns [2]. Changes taking place in these Ni metal and NiO structures under in situ light exposure in presence of water vapor will also be discussed.
References:
[1]. Domen, K.; Kudo, A.; Onishi, T.; J. Catalysis, 1986. 102,92-98
[2]. Miller, B.K.; Crozier, P.A. Microscopy and Microanalysis 2013, 19, 461-469


The support from US Department of Energy (DE-SC0004954) and the use of ETEM at John M. Cowley Center for HR Microscopy at Arizona State University is gratefully acknowledged.

Fig. 1: a) Initial 5%wt NiO on anatase particles, b) zoom in of initial NiO-Ni-TiO2 structure, c) after ex-situ 6hrs exposure to 360nm-460nm light in liquid water, d) zoom in of NiO particle on anatase after exposure

Type of presentation: Poster

IT-7-P-3335 Mems-based heaters for ultrahigh temperature in situ TEM studies

Erdamar A. K.1, Zandbergen H. W.1
1Kavli Institute of Nanoscience, Delft University of Technology, Lorentzweg 1, 2628 CJ Delft, The Netherlands
a.k.erdamar@tudelft.nl

MEMS-based heaters are presently used in situ TEM studies for different purposes and applications such as morphological transformations of gold nanoparticles [1], sculpting of graphene [2], gas nanoreactors [3], thermal stability of nanoparticles [4]. An obvious question is how high we can go with such heaters. The heaters used in references 1-4 were made out of Pt embedded in SiN. Other heater materials like Mo and W can be applied to allow a higher temperature, in particular for temperatures above 1000 0C. Table 1 gives some thermal properties of various materials. Note that these are bulk properties and that for a thin film or in combination with another thin films they can be quite different. The applicability of the various metals depends strongly on the layer package of the membrane and the heater in the MEMS fabrication. For instance we use Pt heaters are embedded in ~500 nm thick SiN membrane, with ~6 µm wide viewing windows of 10-20 nm thick SiN (Figure 2). Since the embedding requires two SiN fabrication steps, Pt has to be stable in the gases used in the second SiN deposition. However, W is not stable in this process and thus embedding of W requires a different fabrication route.
Our MEMS heater contains four electrical connections that allow for temperature determination and heating. The big advantage of MEMS-based heater holder that the heat produced is low and thus little drift. It will be the thickness of the membrane, the size e.g. 1000 µ wide, and its thermal conductivity that determines the heat transfer to the holder. In this respect it is useful to consider the total system from specimen to holder as a set of thermal resistors (Figure 3). The temperature of the sample on a thin window will depend on the heat transfer (and thus the thermal resistances of the components between the sample and the heater. At low temperatures (up to 500°C) the irradiation is relatively small and one can assume that the temperature of the sample is about equal to that of the heater even if the thermal resistance between the thin window and the sample is high. But at high temperature various components of the heater will irradiate which add to the uncertainty of the temperature of the sample.
We are exploring in particular the use of W as heater material, with SiN, SiC and Al2O3 as membrane material. Recently we made heaters that are at least stable at 1250°C over 24 hours and are trying to push this up to 1400°C by a optimization of materials and process steps. We will report on these optimizations in the presentation.

[1] Young, N.P. et.al. Ultramicroscopy 2010, 110, 506-516.
[2] Song, B. et.al. Nano Letter 2011, 11, 2247-2250.
[3] Malladi, S. et.al. Chemical Communication 2013, 49, 10859-10861.
[4] Yalcin, A.O. et.al. Nanotechnology 2014, 25, 055601.


This work is part of the research programme of The European Research Council (ERC) NEMinTEM 267922.

Fig. 1: Table 1: (a) Thermal properties of support materials, (b) and heater materials.

Fig. 2: (a) Optical images of the center of the MEMS-based heater with an embedded Platinum wire for local heating with four connections, (b) electron-transparent windows with a diameter of 6 µm and 20 nm thick SiN.

Fig. 3: Schematic representation of the various thermal resistors (R1-R6) that determine the heat transfer of the heater to the outer tube of the TEM holder. At high T the specimen itself, its contact with the support of the heater and the support-heater contact introduce thermal gradients and thus the real situation is more complicated than indicated.

Type of presentation: Poster

IT-7-P-3342 Applying of 6-carboxyfluorescein (6-FAM) to cytogenetics

Galkina S.1, Saifitdinova A.1, Bogomaz D.1, Radaev A.1, Gaginskaya E.1
1Saint-Petersburg State University, Saint-Petersburg, Russia
chromas.spbu@gmail.com

6-carboxyfluorescein (6-FAM) is one of the most commonly employed and simplest fluorescent reagents to use in oligonucleotide synthesis. 6-FAM is highly reactive, water-soluble single isomer of fluorescein, with absorbance/emission maxima in the visible region of the electromagnetic spectrum (492/517 nm respectively). 6-FAM plays a particularly important role in real-time PCR and SNP-analysis, being used in TaqMan probes, Scorpion primers and Molecular Beacons. Oligonucleotides labeled with 6-FAM at the 5’-end are widely adopted as PCR and DNA sequencing primers to generate fluorescently-labeled products for sequencing and genetic analysis. Per se 6-FAM-labeled oligonucleotides can be used as hybridization probes for fluorescent in situ hybridization (FISH), for example for a direct visualization of microorganisms in human and animal clinical samples (e.g. Behrens et al., 2004; Lin et al., 2011; Fontenete et al., 2013).

We used 6-FAM-labeled oligonucleotide probes specific for various chicken tandem repeats to detect RNA-transcripts and to localize them on giant transcriptionally active lampbrush chromosomes dissected from growing chicken oocytes. Lampbrush chromosomes have distinctive chromomere-loop patterns that enable high-resolution cytogenetic mapping of unique and repeat nucleotide sequences. We report that due to the high brightness and relatively long lifetime, the 6-FAM is found to be well suited for FISH proceeded accordingly with a DNA/(DNA+RNA) hybridization protocol.


The work is supported by SPbSU grant 1.37.153.2014 and grant for Leading Scientific Schools 3553.2014.4. The equipment used was provided by SPbSU Resource Research Centers “Chromas” and “Center for molecular and cell technologies”.

Fig. 1: FISH with oligonucleotide probe PO41 on chicken lampbrush microchromosomes. (a) Representative smallest microbivalents probed with PO41 labeled with 6-FAM (green signal); (b) microbivalents probed with PO41 labeled with Cy3 (red signal). Chromosomes are counterstained with DAPI. Left panels – phase contrast images. Scale bar 5 μm.

Type of presentation: Poster

IT-7-P-3408 In situ TEM deformation of a bulk metallic glass with a K2-IS detector

Gammer C.1, Rentenberger C.2, Karnthaler H. P.2, Czarnik C.3, Beitlschmidt D.4, Pauly S.4, Eckert J.4, Minor A. M.1
1National Center for Electron Microscopy, Lawrence Berkeley National Laboratory, Berkeley, USA, 2Physics of Nanostructured Materials, University of Vienna, Vienna, Austria, 3Gatan, Inc., Pleasanton, USA, 4IFW Dresden, Dresden, Germany
cgammer@lbl.gov

Bulk metallic glasses are an exciting new class of materials due to their unique mechanical properties, such as high strength and good wear resistance. However, potential applications are hindered by their low ductility caused by the formation of shear bands leading to catastrophic failure [1]. The origin of these shear bands remains unknown. In order to investigate the structural mechanisms of shear band formation, in-situ deformation was carried out inside a TEM using a Hysitron Picoindenter and a Gatan K2 high speed direct electron detector. Samples were made from a bulk CuZrAlAg rod produced by vacuum casting.

During the compression of nanopillars slipping events were observed but since the bands were not necessarily in projection they were difficult to analyze. To overcome this problem we used notched pillars to localize the deformation. Fig. 1 shows the results of a compression test carried out in dark-field mode. The video was recorded using a Gatan K2-IS direct detection camera at a frame rate of 400 f/s. The load displacement curve shows a long elastic regime followed by a sudden load drop. Two images were extracted, one right before and one right after the load drop (indicated with a and b, respectively). The images show the formation of a shear band as concluded from the shear offset that is indicated in (b). During further loading multiple small load drops can be observed followed by a larger loaddrop. The images taken before and after the large loaddrop (c and d) reveal that the load drop is caused by an abrupt slip event. The same shear band is reactivated (the trace of the shear band is indicated in (d)). Due to the fast frame rate of the camera it is possible to conclude that the time for the slip event was less than 2.5ms. The results indicate that after the initial formation of a shear band the pillar slips along this shear band in a stick-slip motion.

In addition to compression tests, tension tests of the metallic glass were carried out. Tension samples were made by transferring the sample to a Hysitron Push-to-Pull Device that allows using the Picoindenter as in-situ tensile apparatus [2]. Fig. 2 shows the result from an in-situ tensile test acquired in bright-field mode. The corresponding load displacement curve shows elastic deformation followed by an abrupt fracture with no indication of plasticity. In addition, digital image correlation from decorated samples was used to examine the origin of shear band formation. These results will be described in terms of the relationship between local shear transformation zones and eventual shear band formation and propagation.

[1] A.L. Greer. Science 267 (1995) 1947.
[2] H. Guo, et al. Nano Lett. 11 (2011) 3207.


The authors acknowledge support by the Austrian Science Fund (FWF):[I1309, P22440, J3397] and by the National Center for Electron Microscopy, Lawrence Berkeley Lab, supported by the U.S. Dept. of Energy under Contract # DE-AC02-05CH11231.

Fig. 1: Load displacement curve recorded during in-situ compression. After elastic deformation loaddrops can be observed. The initial one corresponds to the formation of a shear band (images from the dark-field video are shown in a+b). After some small loaddrops, a larger one is observed that results from an abrupt slip along the same shear band (cf. c+d).

Fig. 2: Stress strain curve recorded during an in-situ tensile test. The video acquired in TEM bright-field mode shows an elastic elongation of the tensile specimen. Two frames corresponding to the initial state and the elongated state are shown in (a) and (b). After elongation the sample fractures abruptly and shows no ductility (c).

Type of presentation: Poster

IT-7-P-3442 Direct evidence for orbital angular momentum transfer from electron vortex beam

Thirunavukkarasu G.1, Yuan J.1, McKenna K.1, Babiker M.1
1Department of Physics, University of York, Heslington, York, YO10 5DD, United Kingdom.
t.gnanavel@york.ac.uk

Optical vortices have become well known for a vast range of applications such as optical sensors, tweezers, nanoparticle trapping and manipulation etc., since they were first reported by Allen et al [1]. Compared to optical vortices, electron vortices are relatively new. They were first predicted theoretically by Bliokh et al [2] and experimentally realised by several groups in the subsequent years utilising either a phase plate method [3] or the holographic mask method in a transmission electron microscope [4-6]. These electron vortex beams have the characteristics of orbital angular momentum.

The rotation of gold nanoparticles subject to electron vortex beams has been reported by our group [6] as well as by Verbeeck et al [7]. In this presentation, we focus on experimental evidence that such rotation is direct proof of the mechanical transfer of orbital angular momentum from the beam to the particles. The experiment has been conducted in a JEOL 2200FS double-aberration corrected TEM operating at an acceleration voltage 200kV which utilises a specially designed condenser mask aperture with a fork dislocation to produce the required electron vortex beams. The motion of the nanoparticle subject to the vortex beam illumination is examined by video microscopy and frame-by-frame image analyse. The time series of particle rotation can be obtained and detailed analysis allows the rate and sense of the rotation to be determined.

The chirality of the beam is deduced by comparing the through focus images of the hologram mask in the condenser aperture with the simulation. From the phase structure of the simulated beams, the sense of the rotation of the particle flux can be deduced unambiguously (Fig. 1). As can be seen in figure 2a-d and figure 2e-h a clear trend of opposite rotation in both l = ±1 beams was observed. This shows that the rotation induced on the nanoparticle is consistent with the chirality of the electron vortex beam, indicating direct transfer of orbital angular momentum. However, a detailed examination of the induced rotation shows signs of stochastic processes, indicating that the rotation is dissipative due to friction between the nanoparticle and the substrate.

[1] L Allen et al, Physical Review A 45 11 (1992), p. 8185.

[2] K Bliokh et al, Physical Review Letters 99 (2007), 190404.

[3] M Uchida and A Tonomura, Nature 464 (2010), p. 737.

[4] J Verbeeck et al, Nature 467 (2010), p. 301.

[5] BJ McMorran et al, Science 331 (2011), p. 192.

[6] T Gnanavel, J Yuan and M Babiker, in Proc. European Microscopy Congress, ii, edited by DJ Stokes and J Hutchison (Royal Microscopical Society, Oxford, 2012).

[7] J Verbeeck et al, Advanced Materials 25 (2013), p.1114.


The authors gratefully acknowledge funding from the EPSRC (Grant No. EP/J022098). Thanks are also due to the York JEOL Nanocentre for the provision of microscopy facilities and JEOL, U.K. for financial support.

Fig. 1: Determination of chirality of the electron vortex beams.

Fig. 2: Series of images showing (a-d) anticlockwise rotation under = +1 and (e-h) clockwise rotation under = -1 beams and corresponding FFTs. The arrows are pointing to the same diffraction peak to highlight its angular orientation.

Type of presentation: Poster

IT-7-P-5752 Development of In-Situ Wet-Cell Electron Microscope Holder for Oxygen Nano-bubbles by Platinum

Zheng H. T.1, Liu S. Y.1, Tsai C. T.2, Haung T. W.1, Tseng F. G.1, Chen F. R.1
1Engineering and System Science Department/National Tsing Hua University, Hsinchu, Taiwan, 2Dept. of Material Science and Engineering, National Chung Hsing University, Tai-Chung, Taiwan
applebyapple1@gmail.com

Recently, wet-cell electron microscopy provides a new method for investigating crucial scientific issues within liquid which beyond the conventional electron microscopy. The progress in electron microscopy pushes the capability of viewing as close as the original phenomena occur and thus may open new scientific windows in multi-field due to the spatial resolution in sub-nanometer as well as tens of millisecond time-resolved power. Much more researches has been published and included various field such as electrochemistry [1], catalyst material [2, 3], and biophysics [4]. Heimei et al firstly visualize growth dynamics of Platinum nanocrystal with nanometer resolution in wet cell TEM [2]. The above discoveries provide the key to understanding toward whole mechanism for synthesizing more efficient catalyst materials. In this research, owing to the significance of catalyst, we put more focus on investigating catalytic process of Platinum. For this purpose and to approach real case, we built up the platform with function of liquid circulation to fit with currently TEM (JEOL, JEM-2010 LaB6 equipped with Gatan multi-scan CCD) observation respectively, that is in-situ wet-cell electron microscope holder. The wet-cell chip is made by micro electro mechanical systems (MEMS) process and the observing window is 50nm Si3N4 [5, 6].
For TEM observation, the platinum nanoparticles was carried by multi-wall carbon nanotube was dropped onto wet-cell chip with electron transparent Si3N4 membrane then sealing was completed by a set of o-rings. The solution contains 0.08wt% of H2O2 have been transported into the observation area by syringe pump; additionally, the flow rate was accurately controlled below 0.15 mL/hr that is the key to preventing membranes broken. The bubble formation is due to the well-known equation: 2H2O2→ 2H2O+O2. Platinum serves as catalyst to promote hydrogen peroxide decomposed into water and oxygen, which contribute to the source of bubble generation.
References:
[1] M. J. Williamson, R. M. Tromp, P. M. Vereecken, R. Hull and F. M. Ross, Nature Materials 2 (2003), p. 532.
[2] H. Zheng, R. K. Smith, Y. W. Jun, C. Kisielowski, U. Dahmen and A. P. Alivisatos, Science 324 (2009), p. 1309.
[3] J. M. Yuk, J. Park, P. Ercius, K. Kim, D. J. Hellebusch, M. F. Crommie, J. Y. Lee, A. Zettle and A. P. Alivisatos, Science 336 (2012), p. 61.
[4] M. J. Dukes, D. B. Peckys and N. D. Jonge, ACS Nano 4 (2010), p. 4110.
[5] T. W. Huang, S. Y. Liu, Y. J. Chuang, H. Y. Hsieh, C. Y. Tsai, Y. T. Huang, U. Mirsaidov, P. Matsudaira, F. G. Tseng, C. S. Chang, and F. R. Chen, Lab chip 12 (2012), p. 340
[6] T. W. Huang, S. Y. Liu, Y. J. Chuang, H. Y. Hsieh, C. Y. Tsai, W. J. Wu, C. T. Tsai, Utkur Mirsaidov, P. Matsudaira, F. G. Tseng and F. R. Chen, Soft Matter 9 (2013), p.8854


This work was supported by National Science Council (NSC102-2321-B-007-007 and NSC 102-2120-M-007-006-CC1).

Fig. 1: Fig. 1(a) The cross section view of TEM liquid holder tip: electron beam penetrate the liquid thickness which defined by metal spacer between chips and sets of o-rings are used for vacuum sealing. The nitride membrane is 50 nm for each piece and (b) assembling diagram of TEM liquid holder.

Fig. 2: Fig. 2, TEM images of nano-bubbles, phenomena of the bubbles merge as indicated in dash line region with time by injecting H2O2 solution (a) ~ (g).

Type of presentation: Poster

IT-7-P-5919 Real-time observation of in-situ cation exchange in CdSe-PbSe nanodumbbells during epitaxial solid-solid-vapor growth

Yalcin A. O.1, Goris B.2, Bals S.2, Van Tendeloo G.2, Casavola M.3, Vanmaekelbergh D.3, Tichelaar F. D.1, Zandbergen H. W.1, van Huis M. A.3
1Kavli Institute of Nanoscience, Delft University of Technology, Lorentzweg 1, 2628 CJ Delft, The Netherlands, 2Electron Microscopy for Materials Science (EMAT), University of Antwerp, Groenenborgerlaan 171, 2020 Antwerp, Belgium, 3Debye Institute for Nanomaterials Science, Utrecht University, Princetonplein 5, 3584 CC Utrecht, The Netherlands
a.o.yalcin@tudelft.nl

Both the synthesis and design of hetero-nanocrystals (HNCs) have undergone a rapid development, whereby PbSe and CdSe NCs are key materials acting as functional building blocks within a wide variety of heterogeneous nanostructures.1,2 Heat treatment of HNCs can induce new interface designs, exemplified by the transformation of PbSe/CdSe core/shell systems into PbSe-CdSe bi-hemispheres.2 Here, we report an in-situ heating-induced epitaxial PbSe NC domain growth at the solid-solid PbSe-CdSe nano-interface through cation exchange. We show that Pb replaces Cd at the PbSe/CdSe interface, resulting in growth of the PbSe phase at the expense of the CdSe phase.3 In analogy with vapor-liquid-solid4 and vapor-solid-solid5 growth mechanisms, the currently observed process could be called solid-solid-vapor (SSV) growth as the Cd evaporates, either as neutrally charged Cd atoms or in a molecular complex such as Cd-oleate. Figure 1 shows the elemental maps of CdSe-PbSe HNCs at each stage of the cation exchange during epitaxial SSV growth mechanism. As a result of the cation exchange from CdSe to PbSe, the crystal structure transformed epitaxially from hexagonal wurtzite (WZ) to cubic rock-salt (RS). Figure 2 shows this transformation at atomic resolution. When the HNC was heated from 160 ⁰C (Figure 2a) to 180 ⁰C (Figure 2b), the brighter intensity corresponding to PbSe advanced into the CdSe region. The PbSe RS (200) lattice spacings started to appear along the nanorod domain instead of the CdSe WZ (0002) lattice spacings, as confirmed by the Fourier Transformation (FT) patterns shown in the insets. It is clear that the cation exchange takes place at the PbSe/CdSe interface and propagates epitaxially (layer by layer) along the WZ<0001> direction.

[1] Son, D. H. et al. Science 2004, 306, 1009-1012.
[2] Grodzińska, D. et al. J. Mater. Chem. 2011, 21, 11556-11565.
[3] Yalcin, A. O. et al. Nano Lett. 2014, 14, 3661–3667.
[4] Gudiksen, M. S. et al. Nature 2002, 415, 617-620.
[5] Persson, A. I. et al. Nat. Mater. 2004, 3, 677-681.


This work is part of the research programme of the Foundation for Fundamental Research on Matter (FOM), which is part of the Netherlands Organization for Scientific Research (NWO).

Fig. 1: HAADF-STEM images and chemical maps of CdSe-PbSe HNCs at (a-d) 100 ⁰C (initial configuration), (e-h) 170 ⁰C, and (i-l) 200 ⁰C. In (e-h), a partially transformed nanorod is present. In (i-l), two PbSe-CdSe HNCs became full PbSe domains. The Se remains in place during the transformation. Reprinted with permission from Ref. [3].

Fig. 2: HAADF-STEM images of CdSe-PbSe HNC. With heating from 160⁰C (a) to 180⁰C (b), WZ CdSe nanorod started to transform to RS PbSe. The spot depicted with an arrow in the inset FT in Fig. 2a corresponds to WZ CdSe(0002) spacing. It disappeared in the inset FT of Fig. 2b, confirming the WZ to RS transformation. Reprinted with permission from Ref. [3].

Type of presentation: Poster

IT-7-P-6054 In-situ observation of gold nanorod self-assembly

Novotný F.1, Wandrol P.2, Proška J.1, Šlouf M.3
1FNSPE, Czech Technical University, Břehová 7, 115 19 Prague, Czech Republic, 2FEI, Podnikatelská 6, 612 00, Brno, Czech Republic, 3Institute of Macromolecular Chemistry, Academy of Sciences of the Czech Republic, Prague, Czech Republic
filip.novotny@fjfi.cvut.cz

Self-assembly organize gold nanorods (AuNRs) encapsulated by cetyltrimethylammonium bromide (CTAB) bilayer into an ordered material. The adsorbed molecules not only stabilize colloidal dispersion but also create “glue” among nanorods in the superlattice. Configurational entropy and depletion mediated interactions are commonly considered in the process of supracrystals growth. However, many questions arise about the dynamics of self-assembly in general and particularly about gold nanorods self-assembly in drying colloidal drop. Here we demonstrate the vizualization of the dynamic behaviour of the AuNRs (20 nm x 60 nm) in viscous CTAB/water environment using scanning transmission electron microscopy (STEM) in environmental conditions (STEM-in-ESEM). We observed several distinct stages of AuNRs self-assembly at the liquid-gas interface under space confinement, during controlled evaporation of solvent. The formation of free standing membrane of close-packed nanorods with vertical orientation around the inner edge of the carbon membrane hole, the formation of side-by-side AuNR chains and the sensitivity of the self-assembly process to the irradiation of the electron beam will be shown. Moreover, the viscous environment of the membrane enables to observe the dynamics of the self-assembly process on timescale of seconds. Particular events can be traced such as the nucleation and growth of the 2D crystals around the rim of the holey carbon membrane, the slowing down of the Brownian motion of loose tip-to-tip rod assemblies and convective flows in nano-environment revealed by their collective translation movement besides and the effects of the electron probe upon the prolonged exposure. A time lapse series of micrographs will used to demonstrate such capatibility.


This work was supported by the Czech Science Foundation project P205/13/20110S and P205/10/0348, Technology Agency of the Czech Republic project TE01020118, and also joint FEI & Czechoslovak Microscopy Society scholarship 2011.

Fig. 1: Scheme of in-situ observation of self-assembly of AuNRs. (a) AuNR solution is deposited on top of a holey carbon TEM grid. (b) Droplet undergoes rapid evaporation. (c) Collapsed drop concentrates the AuNR/CTAB volume to form a electron transparent hydrated viscous membrane. (below) Micrograph of AuNR array forming inside holey opening.

IT-8. Ultrafast microscopies

Type of presentation: Invited

IT-8-IN-1690 Imaging surface plasmon polaritons by fs-transmission electron microscopy

Carbone F.1
1LUMES, ICMP, SB, Ecole Polytechnique Fédérale de Lausanne
fabrizio.carbone@epfl.ch

In this seminar, we will review the recent advances in fs-transmission electron microscopy. The design and implementation of a fs-resolved transmission electron microscope will be briefly introduced and its overall performance in terms of time, energy and spatial resolution will be presented. Thanks to this technology, the direct imaging of light-induced surface plasmon polaritons in nanostructures is enabled. When electrons and photons are overlapped spatially and temporally on a nanostructure, the evanescent field photoinduced at the edges of the latter interacts with the electrons allowing them to absorb and emit photons from the pump laser beam. This results in sideband peaks spaced by an energy corresponding to the pump photon energy on both the energy gain and loss sides of the elastic electrons peak. By selecting one or more of these sidebands via energy filtered imaging, snapshots of the surface plasmon polaritons can be taken. Such a technique is called Photon Induced Near Field Electron Microscopy (PINEM). By controlling the properties of light excitation, its energy, polarization and intensity, the distribution of the field around a given nanostructure can be controlled, providing a unique tool for the characterization and manipulation of optoelectronic circuits. The life-time of the surface plasmon polariton waves on metallic materials is found to be ultrafast, comparable to the laser excitation pulse duration (100 fs), and reveals information about the surface morphology and its electronic properties.


This work was supported by an ERC starting grant.

Type of presentation: Invited

IT-8-IN-5754 Time-Resolved Imaging of Surface Plasmon Polaritons by Photoemission Microscopy: The Next Generation

Meyer zu Heringdorf F. J.1
1Faculty of Physics and Center for Nanointegration Duisburg-Essen (CENIDE), University of Duisburg-Essen, 47048 Duisburg, Germany
meyerzh@uni-due.de

Observing surface plasmon polaritons (SPPs) in a photoemission electron microscope (PEEM) is possible via a two photon photoemission (2PPE) process, if ultra-short laser pulses of a suitable wavelength are directed onto a surface with plasmonic structures. In the past, we used a grazing incidence angle of 65-74° of the laser light relative to the surface normal for PEEM-based SPP imaging. The resulting SPP contrast was in this case described as a Moiré-pattern [1,2]. Properties of the SPP, however, can only be inferred indirectly from the Moiré pattern in grazing incidence geometry. For instance, SPPs propagating in different directions across the surface produce Moiré-patterns with a different fringe-spacing. A “normal incidence” geometry – harder to achieve due to the geometrical restrictions of the available PEEMs – is overall better suited for SPP imaging. The cylindrical symmetry caused by the incidence of the laser pulses normal to the surface results in the same imaging conditions for all SPPs, independent of their propagation direction. Also, the spacing of the Moiré fringes resembles the SPP wavelength, and in this respect normal incidence 2PPE PEEM provides a direct conceptual visualization of the SPP phase fronts in time and space. In time-resolved experiments under normal incidence conditions the direct observation of isolated SPP wave packets is then possible. Normal incidence 2PPE PEEM offers the possibility to study SPP reflection, transient SPP interference, and SPP focusing in time and space.

[1] L.I. Chelaru, F.-J. Meyer zu Heringdorf, Surface Science 601 (2007) 4541
[2] N.M. Buckanie, P. Kirschbaum, S. Sindermann, F.-J. Meyer zu Heringdorf, Ultramicroscopy 130 (2013) 49


Financial support by the German Research Foundation (DFG) through programs "SFB616: Energy Dissipation at surfaces" and "SPP1391: Ultrafast Nanooptics" is gratefully acknowledged.

Type of presentation: Oral

IT-8-O-1703 Electron Pulse Properties and PINEM Aberrations in Ultrafast Transmission Electron Microscopy

Flannigan D. J.1, Plemmons D.1
1Department of Chemical Engineering and Materials Science, University of Minnesota, Minneapolis, MN, USA
flan0076@umn.edu

In ultrafast transmission electron microscopy (UTEM), extension of the static imaging and analytical capabilities of transmission electron microscopy to the ultrafast temporal domain relevant for many atomic and nanoscale processes allows for visualization of non-equilibrium structural phenomena. As in pump/probe spectroscopic techniques, the operating principle of UTEM requires – at some point during the experiment – temporal overlap of the femtosecond photon and electron pulses at the specimen. At time zero (i.e., precise overlap of the pulses), significant photon absorption by the freely-propagating electrons and population of virtual states occurs, and peaks occurring at integer multiples of the photon energy can be observed to the gain-side of the zero-loss peak in electron-energy spectra. Because this process, called photon-induced near-field electron microscopy (PINEM), is observed when the pulses are overlapped in space and time at the specimen, proposals for using this phenomenon to measure the total UTEM response function and the electron pulse shape and duration emerged during the initial experimental observations.

In this talk, we will discuss considerations for isolating the inherent artifacts of the highly non-linear near-field interactions from the actual pulse characteristics. Using theory developed to describe these interactions, we will discuss how temporal cross-sections of peaks in the electron-energy spectra corresponding to high-order transitions are expected to exhibit the true temporal behavior of the electron pulses. In general, the exceedingly small portion of the pump laser pulse capable of initiating such transitions results in the temporal widths converging to the electron packet duration. Additionally, population of quantized virtual states occurring for an electron beam focused on the edge of a nanostructure suggest that the resulting energy distribution may produce well-defined chromatic aberrations in images arising from the velocity-dependence of magnetic lens focusing (Fig. 1). As such, we will discuss the prospect for detecting such phenomena and its potential as a means of determining the UTEM instrument response without the need for a spectrometer. Appropriate interpretation of observed spectroscopic and image features should in principle enable systematic temporal and spatial deconvolution allowing for a more accurate depiction of intrinsic ultrafast dynamics, especially the critical initial excitation rising edge which, as advances continue, will drop below 50 fs.


This work is supported by 3M, and acknowledgment is made to the Donors of the American Chemical Society Petroleum Research Fund.

Fig. 1: (a) PINEM energy spectrum. Real-space annular point-spread functions due to (b) quantized chromatic aberration (i.e., PINEM aberration) and (c) spherical aberration. Convolution results in the point-spread function shown in (d). (e) Temporal variation of Fourier intensity for a discrete frequency arising from the convoluted aberrations.

Type of presentation: Oral

IT-8-O-1901 Diffract-before-destroy with electrons?

Egerton R. F.1, Li R. K.2, Zhu Y.3
1Physics Department, University of Alberta, Edmonton, Canada T6G 2E1, 2Department of Physics & Astronomy, UCLA, Los Angeles, California, USA, 3Center for Functional Nanomaterials, Brookhaven Nat. Lab., Upton, NY11973, USA
regerton@ualberta.ca

Free-electron lasers provide x-ray pulses with short enough duration (< 100 fs) to record diffraction patterns from biological molecules (allowing their structure to be determined) before the molecules are destroyed by radiation damage [1]. Since fast electrons are elastically scattered more strongly than x-rays [2], it is reasonable to ask whether damage by radiolysis can be similarly outrun using electrons.


Electrons carry greater momentum than photons of the same energy, leading to additional knock-on damage, but in organic materials the knock-on damage is less severe than that caused by radiolysis [3]. More importantly, electrons carry electrostatic charge that can cause charging of insulating specimens (deflecting the beam or disrupting the specimen) and limit the incident-current density because of electrostatic repulsion between the electrons.


The Brookhaven ultrafast electron diffraction (UED) apparatus [4] produces 100fs pulses containing as many as 106 electrons at 2.8 MeV kinetic energy and this high energy helps to reduce Coulomb repulsion effects. To record ten diffracted electrons per pulse from a 10nm particle (e.g. macromolecule), the beam would need to be focused down to about 500 nm, giving a current density of almost 109 A/cm2. At these high current densities, space-charge effects are more important than statistical repulsion. Even so, it appears impossible to focus the beam to below a few micrometers diameter by means of a solenoid (Fig 1), due in part to the increased energy spread (approaching 0.3%) caused by Coulomb repulsion.


Lateral coherency of the beam is of concern for diffractive imaging, which nevertheless offers advantages over direct imaging since it avoids imaging lenses that degrade resolution because of aberrations and beam crossovers [5].


[1] JCH Spence, U Weierstall and HN Chapman, Rep. Prog. Phys. 75 (2012) 102601.
[2] R Henderson, Quarterly Rev. Biophys. 28 (1995) 171.
[3] RF Egerton, Microsc. Research & Technique 75 (2012) 1550.
[4] XJ Wang et al., J. Korean Phys. Soc. 48 (2006) 390.
[5] BW Reed et al., Microscopy and Microanalysis 15 (2009) 272.


Ray Egerton wishes to thank the Natural Sciences and Engineering research council of Canada for financial support. Renkai Li acknowledges DOE Grants No. DE-FG02-92ER40693 and No. DEFG02-07ER46272, and ONR Grant No. N000140711174. Work at BNL was supported by US DOE-BES under Contract no. DE-AC02-98CH10886

Fig. 1: Beam diameter versus distance along optic axis, calculated using the 3Dmesh method for a 0.16pC bunch of 2.5MeV electrons focused by a solenoid with four different values of maximum field strength B0.

Fig. 2: Pulse profile, calculated using the 3Dmesh method and assuming an initial energy spread of 0.001%. Colors represent electron energy, red being the highest.

Type of presentation: Oral

IT-8-O-2020 Resolving Landau State Dynamics with Electron Vortex Beams

Schachinger T.1, Schattschneider P.1,2, Löffler S.1,2, Stöger-Pollach M.2, Steiger-Thirsfeld A.2, Bliokh K. Y.3, Nori F.4,5
1Institute of Solid State Physics, TU Vienna, Vienna, Austria, 2USTEM, TU Vienna, Vienna, Austria, 3iTHES Research Group, RIKEN, Wako-shi, Japan, 4Center for Emergent Matter Science, RIKEN, Wako-shi, Japan, 5Physics Department, University of Michigan, Michigan, USA
thomas.schachinger@tuwien.ac.at

Since the advent of electron vortex beams (EVB) [1] and techniques to routinely produce them in the TEM [2], some exciting applications have been proposed, such as particle manipulation and mapping magnetic moments on the atomic scale, owing to the fact that they carry quantized orbital angular momentum (OAM) of Lz=mħ as well as a quantized magnetic moment of MBm per electron. With m being the topological charge m=…,-1,0,+1,….

Recently, it has been argued quantum mechanically [3] and classically [4] that EVB in a homogeneous magnetic field resemble free electron Landau states (LS), exhibiting peculiar azimuthal dynamics that deviate from standard electron optics’ Larmor rotation (Ω), where M=0. Depending on the orientation of M (for M≠0) with respect to the magnetic lens field B, cyclotron (double-Larmor) rotation (2Ω) and no rotation (0Ω) have been predicted. In solid-state physics, these states are well known describing phenomena like the diamagnetism of metals [5]. Even though great advances in mapping LS in solid-state systems have been made, their azimuthal dynamics could not be resolved experimentally so far [6].

The application of EVB opens up the road for the observation of free electron LS in the quasi-homogenous objective lens field in the TEM. In breaking the rotational symmetry of the annular structure of an EVB with a Si-knife-edge, it is possible to resolve azimuthal variations by scanning the EVB along the propagation (z-) direction, see Fig. 1a. Due to the converging character of EVB in the TEM, see Fig. 1b, the free electron LS are only approximated well by the EVB in a specific z-shift region. Fig. 2 shows experimental images from that region. Indeed, m-dependent rotational speeds can be seen. Fig. 3 summarizes the experimental data of many measurements by giving the fitted slopes ‹ω›, which represent the electrons’ rotational speed. There, the comparison to the theoretical calculations indicates excellent agreement with the predicted dynamics of no rotation for m<0 states, Larmor rotation for m=0 and cyclotron rotation for m>0. It conclusively shows the OAM dependent Landau state behavior. Note that with the Larmor frequency being Ω~2π·19GHz, the discrimination between those three rotational states represents an energy resolution of ~100µeV. These findings provide new insight into the fundamental properties of LS and prepare the route towards detailed investigations of their otherwise hidden characteristics.

[1] K Y Bliokh et al, PRL 99 (2007), 190404, M Uchida et al, Nature 464 (2010), 08904
[2] J Verbeeck et al, Nature 467 (2010), 09366
[3] K Y Bliokh, et al, PRX 2 (2012), 014011
[4] T Schachinger, Master Thesis, TU Vienna (2013)
[5] L Landau, Z f Phys 64 (1930) , 629

[6] K Hashimoto et al, PRL 109 (2012), 116805


This work was supported by the Austrian Science Fund FWF (I543-N20) and RIKEN iTHES Project.

Fig. 1: (a) Sketch of the experimental setup. Half of the incoming beams (blue, green, red), immersed in the B-field of the lens, are blocked by a knife-edge, which is scanned in the z-direction, showing azimuthal dynamics in the observation plane. (b) Cross section of an |m|=1 beam showing the z-region corresponding to the lateral LS extension 2wB.

Fig. 2: Experimental images of the cut EVB with azimuthal angle measurements (in false colors), showing m-dependent azimuthal dynamics. Negative m-states show decreased angular variations, while the m=0 and the positive m-states show larger angular variations. The red line acts as a guide to the eye and represents a constant azimuthal angle.

Fig. 3: Scatter plot of the fitted slopes for the LS z-region validating the m-dependent rotational speeds. The solid lines stand for the theoretical slopes, e.g. no rotation (blue), Larmor rotation (green) and cyclotron rotation (red) and the dashed lines represent averaged slopes.

Type of presentation: Oral

IT-8-O-3019 High Throughput Imaging in a Multibeam SEM

Ren Y.1, Hagen C. W.1, Kruit P.1
1Delft University of Technology, Delft, the Netherlands
y.ren-1@tudelft.nl

Nowadays there is a growing demand to increase the throughput of scanning electron microscopes, especially in biological research where 3-D images of organ’s structures are desired but too time-consuming. By adopting our Multi-Beam Scanning Electron Microscope (MBSEM) and proper stage moving strategy, the time for constructing a 3D image of 1mm3 brain can potentially be dramatically reduced from 200 days to 1 day. There is also a place for high throughput imaging in the semi-conductor industry where inspecting patterned wafers often asks the high resolution of the SEM, but the standard SEM is too slow.
The MBSEM which we have developed is based on a regular FEI Nova-Nano 200 SEM, but equipped with a novel multi-electron beam source module containing a MEMS fabricated aperture array that delivers a 14x14 array of focused beams with a resolution and current per beam comparable to a state of the art single beam SEM 1,2 .
A Secondary Electron (SE) imaging system, a Transmission Electron (TE) imaging system and a Backscatter Electron (BSE) imaging system have been designed for this MBSEM. The most challenging issue for these 3 imaging systems is how to separate and collect different signals belonging to corresponding beams, especially considering that the pitch of the primary beams on the sample is only 0.5~5 µm, and that the systems should work well for different landing energies and working distances.
Here we will present analysis and the simulation results of the SE and BSE detection systems, and recent experimental results of TE imaging. For the latter we have made use of the in-vacuum high resolution light microscope developed for in-situ correlative microscopy3 as shown in figure 1. We demonstrate that all beams arrive at the specimen in a regular grid (figure 2) and that each beam gives a focused image (figure 2). We will discuss the detection problems that arise when so much data can be detected simultaneously.

References:

1. A. Mohammadi-Gheidari, P.Kruit, Nuclear Instruments and Methods in Physics Research A 645 (2011) 60
2. A. Mohammadi-Gheidari, C. W. Hagen and P. Kruit, J. Vac. Sci. Technol. B28, (2010) 1071
3. A.C. Zonnevylle, R.F.C. van Tol, N. Liv, A.C. Narvaez, A.P.J. Effting, P. Kruit and J.P. Hogenboom; Journal of Microscopy, (2013) doi: 10.1111/jmi.12071.


Fig. 1: Multi-Beam SEM with transmission detection

Fig. 2: Top: Multi-beam probes on the YAG : There are 14*14 beams in the system and the 4 quarters are used to identify beams. Bottom: TE images, based on the beam identification indicated above, TE images of different beams are shown, with the same field of view 4.3 µm, collected simultaneously.

Type of presentation: Poster

IT-8-P-1428 Analysis of image distortion on projection electron microscope image

Iida S.1, Hirano R.1, Amano T.1, Terasawa T.1, Watanabe H.1, Murakami T.2
1EUVL Infrastructure Development Center, Inc.(EIDEC), Tsukuba, Japan, 2EBARA CORPORATION, Fujisawa, Japan
susumu.iida@eidec.co.jp

We have been developing a high throughput projection electron microscope (PEM) for EUV (Extreme ultraviolet) patterned mask inspection system. The PEM provides a sample target with areal illumination at a throughput higher than that obtained from a conventional SEM as shown in Fig. 1. However, image distortion is one of the main issues to be fixed. In order to understand the mechanism behind this issue, simulated PEM images through the imaging electron optics (EO) were analyzed using an upgraded advanced Monte Carlo software CHARIOT. Fig. 2 shows a schematic illustration of a target sample using this approach. Near the pattern with 100 nm half-pitch lines and spaces (L/S), a metal contact with an applied 10 V was added to the substrate. This metal mimics a positively charged area. When a simulated L/S pattern image was obtained by an image sensor placed 30 nm above the target sample, the electrons forming this image could not pass through the imaging optics and remained unaffected by the local charge. On the other hand, when secondary electrons could pass through the imaging optics, the image from the detector placed at a focal plane 200 mm away from the target sample resulted in a distorted image as shown in Fig. 3. These results clearly show that image distortion can be reproduced not by the near-field image but by the focal plane image because the virtual source image is projected on the focal plane in PEM. In the case of focal plane, the L/S patterns bent away from the positively charged area in spite of the fact that secondary electrons should be attracted by the charged area. This phenomenon can be explained using Fig 4. If the SE was bent by the charged area, the virtual source, which SE generates, shifts away from the positively charged area. When the SE bends more, the source shifts farther. As a result, focused L/S patterns are bent away from the positively charged area.These simulation results of image formation including electron scattering and long-range effects of the charging help to understand the mechanism of image distortion and to overcome this issue. At higher energies of SE, the effects of bending become smaller. The energy of SE can be controlled by the extraction voltage. In a novel concept of PEM under development, we applied an extracting electrical field eight times stronger than in conventional PEMs in order to considerably reduce the charging distortion. This reduction effect of image distortion was confirmed by this simulator using the EO data of the novel PEM.


This work was supported by New Energy and Industrial Technology Development Organization (NEDO)

Fig. 1: Schematic illustration of PEM

Fig. 2: Schematic illustration of the test sample

Fig. 3: Simulated image from a detector placed behind the electron optical system

Fig. 4: Schematic illustration of distortion due to local charging

Type of presentation: Poster

IT-8-P-1696 New Operation Modes with the direct detecting pnCCD-camera in Transmission Electron Microscopy

Simson M.1, Hartmann R.2, Huth M.2, Ihle S.2, Müller K.3, Rosenauer A.3, Ryll H.2, Schmidt J.2, Soltau H.1, Strüder L.2
1PNDetector GmbH, Emil-Nolde-Str. 10, 81735 Munich, Germany, 2PNSensor GmbH, Römerstr. 28, 80803 Munich, Germany, 3Universität Bremen, Otto-Hahn-Allee 1, 28359 Bremen, Germany
martin.simson@pndetector.de

The pnCCD’s ability to directly detect and image single electrons in a transmission electron microscope (TEM) at energies from 300 to 80 keV has already been demonstrated [1,2] with the dedicated mechanical setup shown in Fig. 1. For these measurements a pnCCD with a physical pixel size of 48x48 μm² with 264x264 pixels was read out as fast as 1000 fps. At low dose conditions a sub-pixel resolution of 1320x1320 after processing of the raw data has been reached. Meanwhile the pnCCD-camera was used to detect TEM electrons with energies down to 20 keV.
A common limitation of CCD-based cameras is the charge handling capacity. Primary electrons from the TEM generate many signal electrons when they scatter in the bulk of the pnCCD. These electrons are stored in potential wells (the pixels) until the CCD is read out. If the number of signal electrons in one pixel exceeds its charge handling capacity, surplus electrons will spill over to neighboring pixels. This effect is usually called “blooming”.
In order to have an optimal behavior of the pnCCD-camera under many different experimental conditions, several new operation modes were developed. While the normal operation settings offer the best spectroscopic properties of the camera, especially at very low energies, the charge handling capacity is limited to about five 80 keV electrons per pixel and frame, i.e. 1 ms. In the high charge handling capacity (HCHC) mode this can be increased by a factor of 3 to 4 without noteworthy losses of other performance parameters. This means, for a primary electron energy of 80 keV and a readout rate of 1000 fps, up to 16,000 TEM electrons per pixel can be processed during one second.
To process even higher electron rates the anti-blooming (AB) mode is available. In this operation mode the number of electrons that can be stored in one pixel is reduced compared to the HCHC mode but excess charge does not overflow into neighboring pixels. Instead it is drained from the detector and does not contribute to the readout signal. Therefore it is possible to image spots of very high intensity without degrading spatial information. At the same time single electrons can be detected in other sections of the detector.
The XPLUS mode is a hybrid of HCHC and AB mode. The charge handling capacity is increased compared to the AB mode, while overflowing signal electrons are drained off.
In the presentation the advantages of the different operation modes will be explained and visualized by different measurements under varying TEM conditions and primary electron energies from 300 to 20 keV.

[1] H. Ryll et al., Microscopy and Microanalysis 19 (2013), p.1160-1161.
[2] K. Müller et al., Appl. Phys. Lett. 101 (2012), p. 2121101-2121104.


Fig. 1: The directly detecting pnCCD- camera for TEM applications.

Fig. 2: Images of 300 keV electrons taken in different operation modes. In normal operation mode (a) the high electron rate causes blooming. The HCHC mode (b) allows to store more charge in each pixel. In the AB mode (c) surplus charges are removed, suppressing all blooming effects. The XPLUS mode (d) is a hybrid of the HCHC and AB modes.

Type of presentation: Poster

IT-8-P-1895 Design of the novel flange-on high lateral and energy resolution ultrashort electron pulse compression system for ultrafast microscopy

Grzelakowski K. P.1
1OPTICON Nanotechnology
k.grzelakowski@opticon-nanotechnology.com

An instrumental realization of the idea [1,2] for the ultrashort electron pulse source based on the newly developed imaging energy filter called α-SDA (Spherical Deflector Analyzer) [3] is reported. Its compact design enables the realization of the flange–on instrument concept. It consists of six independent subsystems: photocathode/immersion lens, primary electron column, pass energy tuning element, α-SDA as a central part, focusing/ compression column and detector/target with XY-stage, Fig1. The ultrashort photoelectron bunch created by an attosecond laser pulse propagates through the primary column towards the mirror plane of the α-SDA, where according to simulations, the focusing and temporal reversion occurs [2]. As a consequence, the time-divergent primary electrons at the mirror plane are transformed to a time-convergent pulse at the same plane after 2π deflection. It has been also previously shown, that the aberration free imaging properties of the α-SDA assure a very high lateral (<<1nm) and energy (ΔE/E<10-3) resolution.  In the symmetric case with the first time compression exactly at π, the shortest electron pulse behind the α-SDA analyzer is mirror symmetric to the original electron pulse at the photocathode [2]. As a consequence, an extremely dense: ultrashort (<<1fs) and perfectly focused (<<1nm) high energy (104-105 eV) electron bunch strikes the target.

1 K.P. Grzelakowski, US Patent Nr. 7,126,117
2 K. P. Grzelakowski, R. M. Tromp, Ultramicroscopy, 130 (2013) 36
3 K.P. Grzelakowski, Ultramicroscopy 116 (2012) 95


The author acknowledges the financial support by the NCBR (National Centre for Research and Technology) in Warsaw. I would like also to express my gratitude to Dr.Rudolf Tromp for stimulative discussions.

Fig. 1: General outlook of the electron pulse compression system

Type of presentation: Poster

IT-8-P-2148 Sub-picosencond electron beam and femtosecond optical pump system in spin-polarized TEM

Kuwahara M.1, Nambo Y.1, Saitoh K.2, Ujihara T.1, Asano H.1, Takeda Y.3, Tanaka N.2
1Graduate School of Engineering, Nagoya University Nagoya, 464-8603, Japan, 2EcoTopia Science Institute, Nagoya University, Nagoya 464-8603, Japan, 3Nagoya Industrial Science Research Institute, Nagoya 464-0819, Japan
kuwahara@esi.nagoya-u.ac.jp

Time-resolved measurements at nanometer spatial resolutions are very important for investigating relaxation processes, catalyzed reactions and phase-transition phenomena. It is possible to carry out such time-resolved analysis using transmission electron microscopy (TEM) by using a pulsed electron beam as a probe beam. Such an approach has been applied in dynamic TEM (DTEM) and ultra-fast electron microscopy (UEM), which use metals and LaB6 for a photoemission source driven by a pulsed laser. These methods have led to the possibility of four-dimensional electron microscopy with high spatial and temporal resolutions. Spin-polarized transmission electron microscopy (SP-TEM) can satisfy two abilities of spin-resolved imaging and pulsed electron gun operation simultaneously, because the instrument consists of a laser-driven polarized electron source (PES) and a conventional TEM system [1].

Spin-polarized electron can be generated by photoemission from III–V semiconductors with a negative electron affinity (NEA) surface. Several beam parameters of a PES are vastly superior to those for conventional thermal electron beams. A high spin-polarization of 92% and a high quantum efficiency of 0.5% have been simultaneously realized using a GaAs–GaAsP strained-layer-superlattice photocathode. In addition, such a photocathode has the ability to generate a sub-picosecond multibunch beam [2]. In order to realize a pump-probe method using the spin-polarized pulse electron beam, we have developed a synchronizing system and demonstrated a phase-locked TEM image of wobbling probe beam by using a pulse electron beam [3]. TEM images were already acquired with a pulsed electron beam with a 1.4-ns pulse duration. Now we have constructed a new optical system, which can provide a sub-picosecond pulse laser and femtosecond pulse simultaneously, to realize an ultrafast temporal resolution. The figure 1 (b) and (c) show the photograph and the schematic diagram, respectively. The sub-picosecond pulse laser is used to drive the electron gun. Another femtosecond laser is transferred to pump a specimen to create an excited state. The sub-picosecond pulse is generated by narrowing a bandwidth of a seed laser which is emitted from a mode-lock Ti:Sapphire laser. Figure 2 shows the femtosecond and picosecond pulse. The sub-picosecond pulse laser is necessary to keep the high polarization. These results suggest the possibility of pump-probe measurements in SP-TEM using the pulsed electron beam as a probe, allowing nanometer-scale time-resolved spin mapping.

[1] M. Kuwahara et al., Appl. Phys. Lett. 101 (2012) 03310

[2] Y. Honda, et al., Jpn. J. Appl. Phys. 52, 086401-086407(2013).

[3] M. Kuwahara et al., Microscopy 62, 607-614 (2013).


The authors thank Drs. H. Shinada, M. Koguchi and M. Tomita of Hitachi Central Research Laboratory for fruitful discussions and encouragement. This research was supported by MEXT KAKENHI Grant Number 51996964, 24651123, 25706031.

Fig. 1:  (a) Photograph of the spin-polarized TEM, (b) Photograph of pulse laser system and (c) the schematics.

Fig. 2: Auto-correlation amplitudes of pulse lasers as a function of correlation time. (a) a correlation amplitude of a femtosecond pulse laser for excitation of a specimen. (b) a correlation amplitude of picosecond pulse laser for driving an electron source.

Type of presentation: Poster

IT-8-P-2668 Ultrafast measurement in spin- polarized pulse TEM

Nambo Y.1, Kuwahara M.1, Kusunoki S.1, Sameshima K.1, Saitoh K.1,2, Ujihara T.1, Asano H.1, Takeda Y.3,4, Tanaka N.1,2
1Graduate school of Engineering, Nagoya University, Nagoya, 464-8603, Japan, 2EcoTopia Science Institute, Nagoya University, Nagoya, 464-8603, Japan, 3Nagoya Industrial Science Research Institute, Nagoya, 464-0819, Japan, 4Aichi Synchrotron Radiation Center, Aichi Science and Technology Foundation, Seto, 489-0965, Japan
nambo.yoshito@f.mbox.nagoya-u.ac.jp

Investigations of ultra-high-speed phenomena in nanometer scale are important for analysis of dynamics in advanced nano-devices. Furthermore, measurement of a spin information is necessary for densification of magnetic recordings and developments of spintronics devices. It is expected that spin-polarized transmission electron microscopy (SP-TEM) can clarify the spin-dependent dynamics.
  The SP-TEM consists of a laser-driven electron source and a conventional TEM (Hitachi H9000-UHV). The electron source is configured with GaAs-GaAsP strained superlattice photocathode using a negative electron affinity surface, which has high polarization of 92% and high quantum efficiency of 0.5% [1]. Moreover, the photocathode also has an ability of generating a sub-picosecond pulsed electron beam [2]. The Pulsed electron beam is emitted by illuminating a pulse laser to the photocathode.
Figure.1 (a) and (b) show the TEM images obtained by using a continual electron beam and pulsed electron beam, respectively. The image of Fig.1 (a) is wobbled TEM image by an alignment deflector coil. The wobbling amplitude is dramatically decreased by using the pulse beam as shown in Fig.1 (b) [3]. Fig.2 is an illustration of an optical bench for a pump-probe method. A femtosecond pulse laser is separated two beam line by a polarized beam splitter. One of the beam line is delayed by passing a variable delay-line. The other is irradiated to a dispersive grating as a pulsewidth stretcher and extracts sub-picosecond pulse laser. Each of lasers are transferred to the SP-TEM by using optical fibers. Fig.3 shows the pulse-duration of a picosecond pulse laser measured by an auto-correlator. The picosecond pulse lasers are generated by selecting a part of the wavelength of femtosecond pulse laser.
  Consequently, a pump-probe method which picosecond and femtosecond pulse lasers are used for photocathode and specimen excitation respectively can be carried out in the SP-TEM. This instrument gives a possibility of investigations of spin dependent phenomena with a high temporal resolution such as a time-evolution of photo-induced magnetism on nanometer scale.

[1] X G Jin et al., Appl. Phys. Express 1. (2008) 045002
[2] Y. Honda et al., Jpn. J. Appl. Phys. 52 (2013) 086401.
[3] M. Kuwahara et al., Microscopy 62 (6) (2013) 607-614.


The authors thank Drs. H. Shinada, M. Koguchi and M. Tomita of Hitachi Central Research Laboratory for fruitful discussions and encouragement. This research was supported by MEXT KAKENHI Grant Number 51996964, 24651123, 25706031 and Kurata Research Grants from the Kurata Foundation.

Fig. 1: Wobbling TEM images acquired by illuminating (a) continual electron beam and (b) pulsed beam.

Fig. 2: A schematic of the optical bench for the pulse laser.

Fig. 3: Correlation time-structures of series of picosecond pulse lasers measured by an auto-correlator.

Type of presentation: Poster

IT-8-P-2678 Implementing in situ experiments in liquids in the(scanning) transmission electron microscope and dynamic TEM

Abellan P.1, Woehl T. J.2, Russell G. T.1, Schroeder W. A.3, Evans J. E.1, Browning N. D.1
1Pacific Northwest National Laboratory, Richland, USA, 2U.S. DOE Ames Laboratory, Ames, USA, 3University of Illinois at Chicago, Chicago, USA
patricia.abellanbaeza@pnnl.gov

Developing a fundamental understanding of phenomena that take place in liquids, such as nanoparticle growth, protein conformational dynamics or the transformation of active materials during battery operation requires characterization tools able to provide in situ information with nanometer spatial resolution. In principle, this can be achieved using fluid stages in the (scanning) transmission electron microscope ((S)TEM). One of the main experimental challenges in the field is obtaining reproducible data free of beam-induced effects to enable quantitative analysis. Methods of calibration of the amount of radiation damage resulting from beam-induced reactions with the sample continue to be needed [1,2]. For instance, in situ growth of particles in solution by the electron beam is typically observed in (S)TEM experiments and has been used to calibrate the effect of electron dose in a Ag precursor solution in an in situ fluid stage [3] (Figure 1 (a) shows areas where Ag was grown under different experimental conditions). Custom image analysis algorithms can be applied to analyze movies of nanocrystal nucleation and growth and extract important information on growth dynamics and parameters such as the induction threshold below which no nucleation occurs [3] (see an example of image analysis in Figure 1(b)). Besides electron dose, factors such as accelerating voltage, imaging mode (e.g. TEM, STEM, SEM), liquid thickness, and solution composition are expected to affect the results of in situ experiments. Reproducing an experiment in a different instrument operating with different electron optical settings, introduces a large set of variables whose effect must be calibrated. Here, we present our recent developments in the design and implementation of calibration experiments using in situ fluid stages, including an identification of beam-sample interactions for changing imaging and experimental conditions. Since fluid stages are designed to fit in any transmission electron microscopy, the different capabilities of each instrument can be applied to the study of liquid phase reactions. When using fluid stages in combination with the dynamic TEM (DTEM), a combined temporal and spatial resolution of ~10-6 and ~10-10 m, could be achieved (see schematic of the DTEM design at the Pacific Northwest National Laboratory (PNNL) in Fig. 2). The unique qualities of the DTEM that benefit the in-situ experiments with fluid environmental cells will be also discussed.

[1] T.J. Woehl et al., Ultramicroscopy 127 (2013) 2927

[2]J.M. Grogan, Nano Letters 14 (2014) 359

[3] T.J. Woehl et al., ACS Nano 6 (2012), p. 8599; Nano Letters 14 (2014), p. 373.

[4] J.E. Evans et al., Microscopy 62 (2013) p. 147-156


This work was supported by the CII; under the LDRD Program at PNNL. PNNL is a multiprogram national laboratory operated by Battelle for the U.S. DOE under Contract DE-AC05-76RL01830. A portion of the research was performed using the EMSL, a national scientific user facility sponsored by the DOE's Office of BER and located at PNNL.

Fig. 1: (a) Low magnification BF STEM image showing a set of nanocrystals growth experiments using different electron dose rates. (b) Number of particles grown from solution as a function of time measured from an in situ dataset using 300kV, 7.1pA beamcurrent, 3 μs pixel-dwell time and M=40000x in STEM, to give a dose per frame of 39.1 e-/nm2f.

Fig. 2: Schematic of the DTEM showing theupgrades planed at PNNL. Modified from [4]. Copyright 2013 Oxford UniversityPress.

Type of presentation: Poster

IT-8-P-3023 SPIM-Fluid: High-throughput platform based on Light-Sheet Microscopy

Gualda E. J.1, Pereira H.2,3, Pinto C.2,3, Simaõ D.2,3, Brito C.2,3
1Instituto Gulbenkian de Ciências, Portugal, 2Instituto de Biologia Experimental e Tecnológica, Portugal, 3Instituto de Tecnologia Química e Biológica, Universidade Nova de Lisboa, Portugal
emilio.gualda@gmail.com

Drug screens on complex cell models and organisms are a key factor to understand and treat human diseases. However, fast and effective conclusions have been hindered by the lack of robust and predictable models amenable to high-throughput (HT) analysis. Animal models can mimic s pathological features, however species-specific differences may occur and are prone to increase experimental costs. On the opposite side, adherent cell cultures have been used in drug screening and tumor modeling but they do not properly represent biological tissues. Recently, important advances have been made towards the development of 3D cellular models, using human immortalized cell lines, stem cells and other patient derived cells, which better recapitulate features of tissues. These advances bridge the gap between adherent cell culture and animal models, making 3D cellular aggregates an extremely powerful in vitro model for preclinical research.

A major hurdle, hampering the widespread utilization of complex in vitro models, is the lack of robust analytical tools. The development of innovative methodologies will allow more comprehensive readouts, generating more accurate and predictive human cell-based 3D models for drug and toxicity screenings. Imaging techniques like confocal microscopy are not optimal for thick samples, providing a short penetration and long imaging times. As an alternative approach, light sheet microscopy (LSM) has been proposed to overcome those limitations. Novel LSM configurations fusing its inherent capabilities with microfluidics will allow massive live 3D cell cultures studies in real-time and with a high spatio-temporal resolution, enabling sophisticated cell-based assays in 3D cell cultures (disease diagnosis and therapy; drug screening; cell differentiation; etc.). Using this approach we will be able to make HT quantitative analysis of the spatio-temporal organization of the different cell types in a spheroid, as well as the response to different environmental conditions with high resolution, high speed and minimal photo-damage.

We will present new designs and prototypes, and how the use of 3D-cell cultures and full system automation will contribute to measure a large set of biological parameters with statistical relevance to investigate drug response on the central nervous system (CNS), cancer therapy and cell differentiation. Also, it would facilitate the development of new typologies for 3D-cell cultures and optimize staining protocols. Those systems will be primarily devoted to 3D cell cultures studies, but the expansion to other biological systems, such as full brain imaging in zebrafish embryos with cellular resolution, will be also presented.


The authors acknowledge support from Fundação para a Ciência e Tecnologia, Portugal - grants SFRH/BD/80717/2011, SFRH/BD/78308/2011, EXPL/BBB-IMG/0363/2013 and PTDC/EBB-BIO/119243/2010; and from Innovative Medicines Initiative Joint Undertaking (EU), grant agreement n° 115188.

Fig. 1: Schematic of the Light Sheet Fluorescence Microscope at IGC (top). Detail of the SPIM-Fluid set-up (bottom).

Fig. 2: Viability of cells within differentiated neurospheres visualized with NucView 488 and MitoView 633 Apoptosis Kit (Biotium, Hayward, CA, USA) image during 15 hours. Tert-butyl hydroproxide (tBHP) (Sigma), an oxidative stress inducer, was used to trigger apoptosis at a concentration of 1mM in Hibernate medium (Invitrogen).

Type of presentation: Poster

IT-8-P-3340 In-Situ Lorentz Microscopy with Femtosecond Optical Illumination

Gatzmann J.1, Eggebrecht T.2, Feist A.1, Zbarsky V.2, Münzenberg M.2, Ropers C.1, Schäfer S.1
1IV. Physical Institute, Georg-August-University, 37077 Göttingen, Germany, 2I. Physical Institute, Georg-August-University, 37077 Göttingen, Germany
schaefer@ph4.physik.uni-goettingen.de

Ultrafast electron microscopy as a laser-pump/ electron-probe technique allows for the investigation of structural and electronic dynamics occurring at sub-picosecond timescales and nanometer length-scales. However, current implementations necessitate compromises in electron source brightness compared to conventional electron microscopy techniques. In-situ transmission electron microscopy with temporally-structured optical sample excitation, i.e. by employing femtosecond laser pulse trains, offers a complementary approach to access ultrafast processes, without the need for customized pulsed electron sources. To this end, we implement free-space-coupled femtosecond sample excitation in a Schottky field-emission electron microscope and investigate the optical response of magnetic domain structures with Lorentz microscopy. Specifically, we study laser-induced domain rearrangements in polycrystalline iron thin films on silicon nitride membranes which are pumped with single sub-50-fs laser pulses. By inverting the observed image contrast at large defocus, we reconstruct the local in-plane sample magnetization based on a transport-of-intensity approach. Prior to laser-excitation, the iron thin films display the well-known magnetic ripple domain structure (cf. Fig. 1A). Upon optical excitation, at laser fluences below a sharp threshold of about 5 mJ/cm2, single laser pulses induce local magnetic domain wall. At laser fluences above the threshold, a single laser pulse generates a network of magnetic vortex/anti-vortex (V/AV) structures, as depicted in Fig 1B-D. Subsequent laser pulses lead to nearly complete rearrangement of the V/AV network (left panels in Fig 1C and D). While the network is stable without optical excitation and shows no discernible dynamics on timescales of minutes to hours, V/AV annihilation can be triggered by illuminating the sample with laser pulses below threshold. After several low fluence optical pulses, the equilibrium ripple domain structure is recovered. The generation of a V/AV-network is remarkable as it presumably is the result of a partially melted, non-equilibrium spin system which is quickly quenched to a metastable state. Possible processes leading to a V/AV network are discussed on the basis of micromagnetic simulations and with respect to ultrafast all-optical pump-probe experiments. The nature and dynamics of the laser-driven magnetic reorganization will be further experimentally investigated with temporally-structured illumination utilizing femtosecond pulse pairs separated by variable time delays. In conclusion, we report the optically-induced vortex/anti-vortex generation mapped by in-situ Lorentz microscopy and discuss possible pathways for their generation.


We gratefully acknowledge financial support by the DGF through SFB 1073 "Atomic Scale Control of Energy Conversion" and by the DFG and the State of Lower Saxony under grant Inst186/867-1FUGG.

Fig. 1: (A) Electron micrograph with Lorentz contrast prior to optical excitation. (B-D) After single-fs-pulse laser excitation a network of vortices and anti-vortices appears (C, left panel). The reconstructed in-plane magnetization is displayed in (C, right panel) and (B). Subsequent laser pulses lead to a rearrangement of this network (D).

Type of presentation: Poster

IT-8-P-6038 Towards RF photo injector based dynamic transmission electron microscopy with REGAE

Manz S.1, Casandruc A.1, Keskin S.1, Zhang D. F.1, Bayesteh S.2, Hirscht J.1, Felber M.2, Gehrke T.4, Loch R. A.1, Marx A.1, Delsim-Hashemi H.2, Schlarb H.2, Hoffmann M.2, Hada M.1, Epp S. W.2, Floettmann K.1, Miller R. J.1,3,4
1Max Planck Institute for the Structure and Dynamics of Matter, Hamburg, Germany, 2DESY, Hamburg, Germany, 3University of Toronto, Toronto, Canada, 4University of Hamburg, Hamburg, Germany
stephanie.manz@mpsd.mpg.de

The relativistic electron gun for atomic exploration (REGAE) has been designed
to study structure and dynamics in a wide range of systems. Aiming for
a time resolution of far less than 100 fs, we plan to observe fast structural changes
in solid, solution and gas phase with single-shot femtosecond electron diffraction
in the energy range from 2 - 5 MeV.
As a prove of principle study, we investigated static electron diffraction of sample
thicknesses close to micrometer.
This poster will present latest feasibility studies of performing dynamic single shot
real space imaging with REGAE. The requirements for single shot imaging result in
bunch charges beyond pC. Both the electron’s high energy as well as space charge
in the electron bunches call for a special lens column pre- and post-sample.
The lenses need to be strong enough to diminish spherical and chromatic aberrations. In order to achieve nanometer resolution a focal length in the
millimeter to centimeter rage is necessary. For electromagnetic solenoid lenses
this means peak currents on the order of Tesla. Although the relativistic energy of
the electrons decreases space charge fields compared to a dc electron gun or
conventional electron microscopes, they come back in play when considering single
shot imaging. We study the effects of space charge on the resolution for our newly designed lens system. We find that the bunch
charge strongly affects both chromatic and spherical aberrations. Simulations show,
that space charge fields affect the resolution already from fC bunch charges on,
even though only the meanfield is considered yet. An optimized imaging system will be presented as well as
strategies to circumvent chromatic aberrations by temporal pulse shaping with an
additional RF cavity in order to achieve nanometer spatial resolution.


The project receives funding from the Centre for Ultrafast Imaging (CUI) at the University of Hamburg.

Type of presentation: Poster

IT-8-P-6048 Ultrafast transmission electron microscopy with nanoscopic electron sources

Feist A.1, Bormann R.1, Schauss J.1, Gatzmann J. G.1, Rubiano da Silva N.1, Strauch S.1, Schäfer S.1, Ropers C.1
1IV. Physical Institute, Göttingen, Germany
feist@ph4.physik.uni-goettingen.de

Ultrafast transmission electron microscopy (UTEM) is a laser pump/electron probe technique which enables the investigation of ultrafast processes on nanometer length scales [1]. Here, the dynamics of an inhomogeneous system after ultrashort laser excitation are probed by stroboscopic illumination with sub-picosecond electron pulses. However, current implementations to create short electron pulses, employing a flat photocathode, are intrinsically limited by their low emittance.
We present the implementation of a pulsed electron source, based on localized laser-triggered emission from a needle-shaped tungsten emitter [2], which we employ in a commercial Schottky field emitter TEM. Within this setup, we experimentally characterize the minimum spot size, overall brightness and intrinsic emittance of the electron beam. To further study the emission properties of the electron gun, numerical finite element calculations are carried out. In addition, photon induced near-field electron microscopy (PINEM) [3] of metallic nanostructures is utilized to investigate the temporal structure of the electron pulses, currently yielding pulse durations of 700 fs.
These electron bunches will allow us to study structural dynamics of heterogeneous systems at and near interfaces, defects and structural inhomogeneities with a sub-ps temporal and nanometer spatial resolution.
[1] A.H. Zewail, Science, 328, 187 (2010).
[2] C. Ropers, D. R. Solli, C. P. Schulz, C. Lienau, T. Elsaesser, Phys. Rev. Lett. 98, 043907 (2007).
[3] B. Barwick, D. J. Flannigan, A. H. Zewail, Nature, 462, 902 (2009).


IT-9. Electron and X-ray diffraction techniques

Type of presentation: Invited

IT-9-IN-1862 Multiple-scattering assisted electron crystallography

Koch C. T.1
1Institute for Experimental Physics, Ulm University, Ulm, Germany
christoph.koch@uni-ulm.de

The ab-initio determination of crystal structures typically requires highly complete single-crystal diffraction data, i.e. diffraction intensities should have been measured for almost all unique reflections. The reason for this is that, if many more reflections have been measured, than there exist atoms within the structure, the sparseness (peaked nature) of the real-space representation of the charge density (in the case of X-rays) or potential (in the case of electrons) can be utilized to solve the crystallographic phase problem (e.g. by direct methods, or charge flipping, or similar kinematic scattering based techniques). While electron diffraction has the great advantage over X-ray or neutron diffraction, that very small crystallites are already sufficient to produce such single crystal patterns, multiple scattering of electrons within the material generally prevents electron diffraction data to be used in as quantitatively a manner as X-ray or neutron data. This limits the application of electron diffraction tomography [1] to samples that are small along at least two dimensions (e.g. rods), and makes the investigation of other geometries (e.g. platelets) generally more difficult.

It is a well-established truth that, if electron diffraction data corresponding to a few different dynamical diffraction conditions is available, the relative phases of the structure factors that correspond to this data are uniquely determined. This fact is being exploited in structure-factor refinement by quantitative convergent-beam electron diffraction (QCBED) [2,3]. Applying the same real-space constraints as are used for solving the crystallographic phase problem from kinematical diffraction data, a lot less properly phased structure factors are necessary to find the corresponding arrangement of atoms than would be the case if the phases were not known.

In this talk I will show that by applying the recently developed large-angle rocking-beam electron diffraction (LARBED) technique [4], as implemented in the QED plugin [5] for DigitalMicrograph (Gatan), highly quantitative dynamical electron diffraction data sufficient to solve the structure can be acquired from nanocrystals even without tilting the specimen at all.

[1] U. Kolb, E. Mugnaioli, T. E. Gorelik, Cryst. Res. Technol. 46 (2011) 542 – 554

[2] C. Deininger, G. Necker, J. Mayer, Ultramicroscopy 54 (1994) 15-30

[3] J.-M. Zuo, M. Kim, M. O’Keefe, J.C.H. Spence, Nature 401 (1999) 49

[4] C.T. Koch, Ultramicroscopy 111 (2011) 828 – 840

[5] http://www.hremresearch.com

[6] C.T. Koch and J.C.H. Spence, Journal of Physics A: Mathematical and General 36 (2003) 803-816


Financial support by the Carl Zeiss Foundation as well as the German Research Foundation (DFG, Grant No. KO 2911/7-1) is acknowledged.

Fig. 1: Illustration of the recovery of structure factor phase triplets from a simulated 2D rocking curve (LACBED disc of radius 2°) for a single reflection of 3.5 nm thin GaAs by applying a recently developed scattering path expansion [6]. The structure factor phases can be further refined assuming sparseness of the potential in real-space.

Fig. 2: (001) LARBED pattern of SrTiO3. The range of beam tilts applied for acquiring this pattern spans the disc indicated by the red circle (diameter = 140 mrad). The beam tilt has been compensated by the diffraction shift coils to produce non-overlapping discs. Individual background-subtracted discs have been extracted and are shown as well.

Type of presentation: Invited

IT-9-IN-2477 Application of 3D EBSD: Growth of tin whiskers and hillocks

Michael J. R.1, Susan D. F.1, Rye M. J.1
1Materials Science Center, Sandia National Laboratories, Albuquerque, NM USA
jrmicha@sandia.gov

International agreements now require lead-free surface finishes for most electronic components. Pure tin finishes have been shown to form whiskers and hillocks. Tin whiskers can grow rapidly to long lengths that can cause faults in electronic devices and circuits. Long whiskers have been shown to be largely single crystals with specific crystallographic growth directions. Hillocks are smaller bump-like growths that can be made up of single grains or many grains and are less likely to cause electrical faults. Both whiskers and hillocks have been shown to grow from their bases.
The actual growth mechanisms for these features still remain elusive, although tin whiskers and hillocks are generally thought to grow as a response to stress in the plated tin films. In order to provide more information about the growth mechanism it is important to have a complete understanding of the crystallography of the growth substrate and the whiskers or hillocks that form. In this work we will apply the technique of 3D EBSD using the dual platform FIB/SEM to provide a more complete picture of whisker crystallography.

Shown in Figure 1 is a long thin whisker that contains low-angle grain boundaries. In this case, the whisker formed first followed by the formation of a hillock grain at the base, which caused subsequent whisker growth to stop. Note that the growth directions of the whisker segments are <001> and are aligned with the <001> direction in the hillock grain. The 3D crystallographic reconstruction shown in Figure 1 allows all of the grains at the base of the whisker to be visualized which is not possible with single 2D sections as are normally used for EBSD orientation maps.

Figure 2 is a 3D EBSD reconstruction of a hillock on the same Sn-film as the whisker in Figure 1. This reconstruction demonstrates that the hillock is polycrystalline with a grain size that is substantially larger than that of the electrodeposited Sn-film. When the area under the hillock is examined there is considerable grain growth associated with the film grains that continue into the hillock, indicating that grain growth and recrystallization may be contributing to hillock growth.

The capability to visualize whiskers and hillocks and the underlying grains in the plated films with 3D EBSD provides new insights into possible growth mechanisms and therefore may enable strategies to eliminate their occurrence on tin plated Cu surfaces


Sandia National Laboratories is a multi-program laboratory operated by Sandia Corporation, a wholly owned subsidiary of Lockheed Martin Company, for the U.S Department of Energy’s National Nuclear Security Administration under contract DE-AC04-94AL85000.

Fig. 1: 3D reconstruction of a straight tin whisker that has grown from a pure tin-plated surface on Cu.

Fig. 2: 3D reconstruction of a large tin hillock that grew on a pure tin-plated surface on Cu. Note that the hillock is polycrystalline as compared to the single crystal whisker in Figure 1.

Type of presentation: Oral

IT-9-O-1603 Split-Illumination Convergent Beam Electron Diffraction (SICBED) of Strained Crystals

Houdellier F.1, Arbouet A.1, Hÿtch M. J.1, Snoeck E.1
1CEMES-CNRS, Université de Toulouse, 29 Rue Jeanne Marvig, 31055 TOULOUSE FRANCE EU
florent@cemes.fr

Convergent beam electron diffraction (CBED) is a well-established Transmission Electron microscopy (TEM) method mainly used to characterize crystal structural properties like space group, charge density distribution or strain. CBED became very popular with the advent of stress engineering in microelectronics where the carriers mobility can be enhanced by tuning the strain of a channel in transistors. It is then of major interest to measure the strain state of the channel in order to understand the electronic properties of transistors. To tackle this problem, various strain measurement methods have been developed by X-Rays analysis, wafer curvature measurements or TEM techniques. Among different TEM methods (HREM, dark-field electron holography, (nano)diffraction, …) CBED is the most sensitive, because of the strong influence of the crystal parameters on the position of the High Order Laue Zone (HOLZ) lines. Furthermore, it has been shown by dynamical simulation method, that the HOLZ rocking curve is extremely sensitive to the displacement field changes along the electron beam path. This explains the occurrence of HOLZ line broadening when the strain is not constant along electrons trajectories in thin sample where surface relaxation occurs. CBED is therefore extremely efficient for local strain measurement, however, like all other strain measurement methods in TEM, the absolute strain measurements necessitates the use of an unstrained reference area. Depending on the sample geometry this reference can be located far from the area of interest. As example, in an epitaxial layer, due to stress relaxation, the substrate generally used as reference can be strained over hundreds of nanometers below the interface. In order to have access in a single CBED pattern to both the area where the strain measurement has to be performed and the unstrained reference (the latter being located microns apart from the former), we have developed a new convergent beam diffraction method, which combines split illumination and CBED optical configuration. This method called split-illumination CBED (Fig. 1) has been developed on the I2TEM microscope. I2TEM is an Hitachi HF3300C TEM fitted with a 300kV cold FEG, an electrostatic biprism located above the three condensors illumination system, two stages capability, a multibrism set-up, a 4k X 4k camera and a Cs-corrector from CEOS. The biprism located above the condenser system and the adjustment of the three condensors allow to separate the convergent beam in two parts which can be shifted apart on the surface of the sample and be recombined in the focal plane of the objective lens. In a single CBED disk, this allows half of the disk to from the strained region and half from a reference region located far from it (Fig. 2).


The authors acknowledge financial support from the European Union under the Seventh Framework Programme under a contract for an Integrated Infrastructure Initiative Reference 312483-ESTEEM2 and the National Research Agency under the program “Investissement d’Avenir” reference No. ANR-10-38-01-EQPX.

Fig. 1: A : SICBED optical configuration. The condenser biprism is used to split the beam convergent in two parallel half cones. B: Geometry of the FIB prepared SiGe/Si multilayers sample studied by SICBED

Fig. 2: SICBED patterns obtained when the condenser biprism voltage increase. Zone 2 (a=0V, b=3V, c=6V, d=9V), Zone 1 (1=0V, 2=3V, 3=6V, 4=9V)

Type of presentation: Oral

IT-9-O-2129 Mapping distortion and strain with EBSD in Cu single crystals

Kalácska S.1, Groma I.1, Ispánovity P. D.1
1Eötvös Loránd University, Budapest, Hungary
kalacska@metal.elte.hu

Cross-correlation based analysis of electron backscatter diffraction (EBSD) patterns is often carried out to map plastic strain variations in deformed polycrystalline samples [1]. In this work this method is applied to characterize the evolution of dislocation structures and corresponding distortion fields in Cu single crystals compressed to different levels. We aim at developing a statistical method that can be used to measure the total dislocation density in the specimen.

Firstly, the effects of sample surface preparation methods were investigated including Ar ion polishing and traditional electropolishing treatments. Then the distortion maps of the specimen are computed with the cross-correlation technique. This method is capable of detecting changes of the crystal orientation to higher accuracy than the commercial software provided for standard EBSD devices that analyse each EBSD pattern individually. The distribution of distortions shows broadening with increasing load and a slow decay. To give a more detailed evaluation of the microstructure the measurements are complemented with the analysis of broadened X-ray diffraction (XRD) peaks. The total dislocation density and its fluctuation within the sample are determined by the variance method [2,3]. The good qualitative agreement found between the two methods indicate that the cross-correlation method is capable of giving a statistical characterization of the dislocation structure.

In the last part of the talk EBSD measurements on thin foils are presented where the cellular dislocation structure can be directly observed by transmission electron microscopy. These results demonstrate the advantage of the EBSD method compared to XRD analysis, namely that the former is not only capable of determining the dislocation density but also yields the spatial distribution of dislocations.

References:

[1] T.B. Britton and A.J. Wilkinson, High resolution electron backscatter diffraction measurements of elastic strain variations in the presence of larger lattice rotations. Ultramicroscopy 114 (2012) 82-95.
[2] I. Groma, X-ray line broadening due to an inhomogeneous dislocation distribution. Phys.Rev.B 57 (1998) 7535-7542.
[3] F. Székely, I. Groma and J. Lendvai, Changes in the dislocation density fluctuations during plastic deformation. Scripta Mat. 45 (2001) 55-60.


Special thanks to Károly Havancsák, Zoltán Dankházi and Gábor Varga for consultation and valuable suggestions.

Type of presentation: Oral

IT-9-O-2147 Using electron vortex beams to distinguish enantiomorphic space groups

Juchtmans R.1, Verbeeck J.1
1University of Antwerp, Antwerp, Belgium
roeland.juchtmans@uantwerpen.be

Twenty two truly chiral space groups exist which are characterized by a screw axis. They can be divided in eleven enantiomorphic pairs of two space groups being each others mirror image. Telling apart crystals belonging to enantiomorphic space groups appears to be a difficult task. A few methods have been developed making use of dynamical scattering in which experimental observations have to be compared with numerical simulations [1-4]. We propose a new method to distinguish enanthiomorphic space groups without the need for simulations, based on the use of electron vortex beams in the kinematical approximation allowing a direct interpretation of the handedness of a crystal.
Ever since their first creation [5,6], electron vortex beams (EVB) have been subject of intensive  research [7]. EVB are solutions of the free space Schrödinger equations of the form Ψ(r,φ,z)=exp(imφ)Ψ(r,z). Being eigenfunctions of the orbital angular momentum operator, they carry a well defined orbital angular momentum (OAM) of mħ per electron and a transverse current around the vortex core. In order for the wave function to be continuous, the intensity of the beam has to be zero in the center of the beam resulting in the well known donut shape of the beam, a bright ring with a dark hole in the middle. As can be seen in fig.1, the wave fronts of such a beam have an helical form. Based on a simple model we have derived a relationship between the symmetry of the higher order Laue zones in the diffraction pattern and the OAM of the vortex when scattered kinematically on helically arranged atoms, as is shown schematically in fig.1. For crystals having one heavy atom near a 3-fold screw axis this provides a simple way of measuring the chirality of the space groups without the need for simulations. We verify our conclusions with multislice simulations of the diffraction patterns shown in fig.2 and fig.3 and we discuss the feasibility with experimental results.


[1] Goodman, P. & Johnson, A. W. S. (1977). Acta Cryst. A33, 997–1001.
[2] Goodman, P. & Secomb, T. W. (1977). Acta Cryst. A33, 126–133.
[3] Haruyuki I. et al. (2003), Acta. Cryst. B59, 802-810.
[4] Johnson, A. W. S. (2007), Acta Cryst. B63, 511-520.
[5] Uchida M. & Tonomura A. (2010), Nature 464, 737.
[6] Verbeeck J., Tian H. & Schattschneider P. (2010), Nature 467, 301.
[7] Verbeeck J. et al. (2014), C. R. Phys., http://dx.doi.org/10.1016/j.crhy.2013.09.014".


This research was supported by an FWO PhD fellowship grant (Aspirant Fonds Wetenschappelijk Onderzoek - Vlaanderen). The authors acknowledge support from the EU under the 7h Framework Program (FP7) under a contract for an Integrated Infrastructure Initiative, Ref. No. 312483-ESTEEM2, the European Research Council under the FP7, ERC grant N246791 – COUNTATOMS and ERC Starting Grant 278510 VORTEX.

Fig. 1: A vortex beam scattered on helically arranged atoms in a crystal. In our setup the vortex core coincides with the screw-axis and the size with the radius of the helix.

Fig. 2: Multislice simulation of the zeroth and first order Laue zone of the diffraction pattern of a focused 300keV vortex beam with convergence angle 8mRad and OAM=+1, scattered on a 3-fold screw-axis in right handed Mn2Sb2O7. The sample thickness is 20nm.

Fig. 3: Same as fig.2 for the left handed enantiomorph. The lack of 2-fold symmetry in the first order Laue zone, in contrast to fig.2, allows a direct interpretation of the handedness of the crystal.

Type of presentation: Oral

IT-9-O-2387 Retrieving nanoscale third-dimension information directly from TEM data using stacked-Bloch-wave simulations and artificial neural network tools

Pennington R. S.1, Van den Broek W.1, Koch C. T.1
1Institute for Experimental Physics, Albert-Einstein-Allee 11, Ulm University, 89081 Ulm, Germany
robert.pennington@uni-ulm.de

Transmission electron microscope (TEM) specimens are three-dimensional, but TEM images and diffraction patterns are two-dimensional. To retrieve the "third-dimensional" information, we have developed a direct-retrieval algorithm including dynamical diffraction that can use TEM data (such as a single convergent-beam electron diffraction [CBED] pattern) and retrieve variations of a range of nanoscale specimen parameters, including strain, crystal tilt, and chemical composition. The retrieval algorithm itself is detailed elsewhere [1], and uses the stacked-Bloch-wave algorithm [2-3] and artificial neural network optimization tools [4]. In this work, we show the effectiveness of our algorithm and discuss considerations for applying this algorithm to realistic experimental data.
A demonstration of this algorithm’s third-dimension (depth-dependent) retrieval ability is seen in Figures 1 & 2. Figures 1 and 2 show CBED patterns of a 100-nm-thick Si specimen at 80 kV at the [110] zone axis, simulated using the stacked-Bloch-wave [2-3] forward-simulation algorithm and 197 zero-order-Laue-zone reflections. Figure 1 has "asymmetric" diffraction features due to the third-dimension variation of crystal tilt. Figure 2 is a CBED pattern like that in Figure 1 but without third-dimension variation, and fails to reproduce the correct diffraction features. Figures 3 and 4 demonstrate our algorithm’s effective and accurate retrieval of third-dimension variation in crystal tilt (Δα, mean over all layers) from the specimen shown in Figure 1a, and decreasing mismatch between simulated and experimental CBED intensity (given by ΔE, mean over all reciprocal-space points). Figure 4 shows how well the unknown α is determined for a known E mismatch.
This algorithm can retrieve third-dimension material properties from a single CBED pattern; however, other techniques like dark-field image series or large-angle rocking-beam electron diffraction (LARBED) series can also be used. Each technique has its own advantages and challenges, especially for analysis of strain or compositional variations. Large lattice-parameter variations can also require a modification to the algorithm in [1].
In this work, we present practical considerations for using our third-dimension information-retrieval algorithm [1]. We demonstrate its effectiveness, discuss different acquisition techniques and consider how different parameters affect our algorithm.
[1]: R. S. Pennington, W. Van den Broek, C. T. Koch. (submitted)
[2]: R. S. Pennington, F. Wang, C. T. Koch. Ultramicroscopy, 2014. http://dx.doi.org/10.1016/j.ultramic.2014.03.003
[3]: D. J. Eaglesham, C. J. Kiely, D. Cherns, and M. Missous. Phil. Mag. A 60, 161 
(1989).
[4]: R. Rojas. Neural Networks: A Systematic Introduction (Springer Verlag, Berlin, 1993).


We acknowledge funding from the Carl Zeiss Foundation and Grant No. KO 2911/7-1 of the German Research Foundation (DFG).

Fig. 1: Simulated zero-loss-filtered convergent-beam electron diffraction (CBED) pattern, generated from a specimen with third-dimension crystal tilt variation. The specimen has ten 10 nm layers, tilted along the [001] direction {0.00, -0.04, -0.10, -0.20, -0.30, -0.30, -0.20, -0.14, -0.06, 0.00} degrees, respectively.

Fig. 2: A CBED pattern like Figure 1, but generated from a specimen with no layer-by-layer crystal tilt variation but with the same mean crystal tilt, which fails to reproduce the "asymmetric" diffraction features seen.

Fig. 3: Our algorithm [1] retrieves third-dimensional variation in crystal tilt (see text) using a (13x13) point reciprocal-space grid, each point 0.05 degrees apart, starting at the 000 point and moving in the [001] and [-110] directions. (This area does not correspond to the discs in Figure 1, but is from the same specimen.)

Fig. 4: The unknown third-dimension parameter mismatch (Δα, mean over all layers), plotted as a function of the known intensity mismatch (ΔE, mean over all points).

Type of presentation: Oral

IT-9-O-2763 Measuring strain with high precision and high spatial resolution using precession and convergent beam electron diffraction

Rouviere J.1, Martin Y.1, Beche A.2, Cooper D.3, Bernier N.3, Vigouroux M.3, Zuo J.4
1CEA, INAC/SP2M UJF-Grenoble Minatec campus, 17 rue des Martyrs, 38054 Grenoble Cedex 9, France, 2FEI Electron Optics, Achtseweg Noord 5, 5651 GG Eindhoven, The Netherlands, 3CEA, LETI, Minatec campus, 17 rue des Martyrs, 38054 Grenoble Cedex 9, France, 4Univ Illinois, Dept Mat Sci & Engn, 1304 W Green St, Urbana, IL 61801 USA
jean-luc.rouviere@cea.fr

Stimulated by the demand of the semiconductor industry, several new TEM based techniques have been recently proposed to measure strain with high sensitivity and high spatial resolution. In this presentation the interest of using diffraction techniques, either Convergent Beam Electron diffraction (CBED, fig. 1) or Nanobeam Precession Electron Diffraction (N_PED, Fig. 2 and 3) [1] will be shown. Off-axis CBED can give 3D maps of the complete 3D strain tensor ε or equivalently of the deformation gradient tensor F (Fig. 1b), but it is computationally and experimentally demanding. In constrast, N-PED is a straightforward and simple technique, although it is limited to the projected 2D strain tensor. Thanks to its robustness, great precision of about 2x10-4  and simplicity, N-PED should be the preferred tool for the microelectronics industry (Fig. 2).

Fig. 1 illustrates the principle of our strain measurement using off-axis CBED. The originality of our approach is to use both the deficient HOLZ lines of the transmitted beam and the excess HOLZ lines of the diffracted beams to measure the strain. Using Bloch wave calculated CBED patterns as tests, we could retrieve of 7 out of the 9 components of the deformation gradient tensor F (Fig. 1b); in particular the volume of the cells can be determined (Fig. 1c). By using two different directions which makes an angle of 22°, we show that it is possible to determines the whole tensor F. In addition, the method can also be extended to the analysis of split HOLZ lines that allow measuring the variations of the strain tensor along the electron beam.
For N-PED, best results were obtained on a FEI TITAN microscope using a 2kx2k CCD camera. Strain maps of 40x50 points can be acquired in about 20 minutes (Fig. 2). Precession can be used either with nearly parallel beam (NBED like condition, Fig. 3b) or with a convergent beam (on-axis CBED like condition, Fig. 3d). Slightly higher precision were obtained by using the CBED like condition. The main advantage of precession is to suppress the contrasts in the diffraction disks, which leads to improved strain precision.

A major advantage of diffraction based techniques is to be able to analyze samples of non-uniform thickness and non uniform composition along the electron beam. To demonstrate this, results on core shell nanowires (NWs) - Ge NWs embedded with SiN, or Si NWs with a surrounding polycrystalline gate - observed either parallel or perpendicular to the growth direction will be presented.


[1] J.L. Rouvière et al., Appl. Phys. Lett. 103 (2013) 241913.


This work was supported by several projects and contracts: the European catrene UTTERMOST project, the French ANR AMOS and the FEI-CEA common laboratory.

Fig. 1: (a) Simulated CBED pattern along a <651, 441,31> direction in Si. (b) Definition of the deformation gradient tensor F and its link to the strain tensor ε and rotation θ. (c)  Without using the excess HOLZ lines, fzz and (fxx+fyy) are correlated. The volume (fzz+fxx+fyy) can be determined only by using the excess lines.

Fig. 2: Maps of 2 transistors with SiGe source (S) and drain (D). As SiGe has a greater lattice parameter than Si the Si channel is compressed by the SiGe source and drain in the x-direction (εxx). (a) HAADF image computed from the series of N-PED patterns. (b-c-e-f) Strain and rotation maps. d) A typical N-PED pattern of the series.

Fig. 3: Diffraction patterns obtained for various semi-convergence angles α and beam diameters d. At the bottom right of each diffraction, an image of the associated electron probe passing through the [011] silicon crystal gives an estimation of d. In (a) and (b) α = 0.6 mrad (NBED like condition), in (c) and (d) α = 2.2 mrad (CBED like condition).

Type of presentation: Oral

IT-9-O-2773 Combining real space and reciprocal space tomography in the TEM

Eggeman A. S.1, Krakow R.2, Midgley P. A.3
1Department of Materials Science and Metallurgy, University of Cambridge, Cambridge, CB3 0FS, UK
ase25@cam.ac.uk

Precession electron diffraction (PED) [1] is a valuable technique for investigating crystal structures and when combined with a well-defined raster is able to produce high quality virtual dark-field (VDF) images and orientation maps [2] with ca. nm resolution. Many microstructures vary in all 3 dimensions and tomographic techniques are needed to investigate such complex structures. Combining scanning PED (SPED) with electron tomography offers a way to study local orientation across a volume of interest in 3D.

In this study a tilt-series (from -60o to +70o with 5o steps) of SPED images were recorded from a Ni-base superalloy sample, the scanned maps were recorded with 140x140px of 7.5nm step with a 5nm probe and a precession angle was 0.5o. The image processing is shown in Fig. 1: a) a VDF image (at 5o) that shows a large (ca. 200nm) precipitate, b) shows a VDF image of a smaller (ca. 50nm) inclusion, c) shows the components in the microstructure after segmentation. Geometric tomography (shape-from-silhouette) [3] was used to recombine the VDF tilt-series into a tomogram, Fig. 1(d). The tomogram allowed the contributions to each diffraction pattern in the tilt-series to be determined. As such, the individual diffraction components could be isolated and combined to produce 3D reciprocal lattice reconstructions.

The smaller particle was found to have the ordered η-phase structure (sp. gr. P63/mmc, a=5.314Å and c=8.351Å), the larger precipitate had the MC carbide structure (sp. gr. Fm3m, a=4.32Å) and the γ-matrix has the disordered fcc structure (sp. gr. Fm3m, a=3.59Å). The correspondence between the orientation of the diffraction pattern and the tilt step allowed the orientation of the different phases to be examined. The (001) plane of η-phase has a coherent registry with (111) of the γ-matrix. A test of the reciprocal lattice alignment confirmed that this registry existed across the ‘top’ facet of the inclusion. In the literature there has been no reported registry between the MC and γ phases. However, the front facet of the precipitate (shown in Fig. 2a) was found to be parallel to the (111) plane, the projected reciprocal lattice from this component at the appropriate orientation is shown in Fig. 2(b). The corresponding reciprocal lattice for the matrix is shown in Fig. 2(c) and returned the (531) plane as the matrix surface. Since the entire reciprocal lattice is projected, for clarity the ZOLZ reflections are highlighted and indexed where appropriate. Analysis of the interface showed a semi-coherent registry with the inclusion of a small interface strain (ca. 4%).
[1] R. Vincent & P. A. Midgley, Ultram. 53 (1994), 271-282
[2] P. Moeck et al. Cryst. Res. Tech., 46 (2011), 589-606
[3] Z. Saghi et al. J. Phys. Conf. Ser. 126 (2008) 012063


The authors acknowledge funding from the ERC though grant 291522-3DIMAGE, the 7th Framework Programme of the EC: ESTEEM2 and Rolls Royce plc.

Fig. 1: Figure 1a) and b) virtual dark-field images of second phase particles in a nickel-base superalloy microstructure, c) segmented image of the two particles and d) representative surface render of the reconstructed tomogram.

Fig. 2: Figure 2a) Tomogram surface of a carbide precipitate normal to its largest facet. b) and c) projections of the reconstructed reciprocal lattices from the carbide and matrix, respectively, at the same orientation. The ZOLZ reflections are highlighted in each showing that the interface is composed of tye carbide (111) and the matrix (531) planes

Type of presentation: Oral

IT-9-O-2884 Orientation measurements in TEM foils using transmission EBSD

de Kloe R.1, Nowell M. M.2, Suzuki S.3, Wright S. I.2
1EDAX, Ringbaan Noord 103, 5046 AA Tilburg, The Netherlands, 2EDAX, 392 E. 12300 S., Suite H, Draper, UT 840201, USA, 3TSL Solutions KK, #SIC2-401, 5-4-30, Nishihashimoto, Midori-Ku, Kanagawa, Sagamihara 252-0131, Japan.
rene.de.kloe@ametek.nl

For a number of imaging modes in the TEM, knowledge of the orientation is critical. For example dislocation analysis by weak beam dark field imaging requires orienting the grain of interest along one of a limited number of orientations (fig 1). Obtaining this orientation can be done by diffraction pattern analysis of multiple zone-axes. Only when multiple zone axes are identified can it be determined if the zone axes required for the imaging of certain defects are within tilting range. This can be a time consuming process with often limited success, especially on low-symmetry materials. There are a number of automated orientation mapping methods available in the TEM that can assist in this orientation determination [1,2], but the available analysis area in the TEM is limited and it is difficult to obtain a complete overview of a sample.
Combining orientation measurements in the SEM with subsequent TEM analysis can bridge this gap. Standard EBSD measurements can be obtained from most electron transparent crystalline samples that have been prepared for the TEM. Such samples can be mounted in the traditional 70 degree tilt orientation to collect larger area EBSD maps. Recently high resolution EBSD results have also been collected using TEM foils in transmission mode in the SEM (fig 2) [3,4]. But in addition to high resolution orientation mapping, this transmission analysis mode also allows identification of the electron transparent areas in the sample. And in combination with orientation simulations the transmission EBSD orientation results can be used to identify grains that are suitable for specific diffraction analysis on the same sample.

 

[1] Rauch E. F., Véron M., Portillo J., Bultreys D., Maniette Y., Nicolopoulos S., Automatic Crystal Orientation and Phase Mapping in TEM by Precession Diffraction. Microsc. and Anal. 93 (2008) S5-S8
[2] Dingley, D. J. (2006). "Orientation imaging microscopy for the transmission electron microscope." Microchimica Acta 155(1-2): 19-29
[3] Trimby P.W. Orientation mapping of nanostructured materials using transmission Kikuchi diffraction in the scanning electron microscope. Ultramicroscopy. 2012 Sep;120:16-24.
[4] Suzuki S. Features of Transmission EBSD and its Application. J.Japan Inst. Met. Mater’ Vol.77(2013), p268-275


Fig. 1: Weak beam dark field images of a dislocation structure in olivine imaged along different g-vectors

Fig. 2: Images of same area of 8Cr tempered martensite steel specimen. top: TEM bright field image (200kV), middle: t-EBSD IQ map (25kV), bottom: t-EBSD IPF crystal direction map // sample normal. The (sub)grain boundary structure is clearly represented in the t-EBSD images [4].

Type of presentation: Oral

IT-9-O-2923 Study of nanoscale local structures of ferroelectric barium titanate using convergent-beam electron diffraction

Tsuda K.1, Sano R.1, Yasuhara A.2, Tanaka M.1
1Institute of Multidisciplinary Research for Advanced Materials, Tohoku University, Sendai, Japan, 2JEOL Ltd., Tokyo, Japan
k_tsuda@tagen.tohoku.ac.jp

  Convergent-beam electron diffraction (CBED) is established as the most powerful technique to determine crystal point- and space-groups from nanometer-sized specimen areas.1) The CBED method was extended to quantitative crystal structure analysis by Tsuda and Tanaka,2, 3) which enables determinations of structural parameters such as atom positions, atomic displacement parameters (ADPs), as well as electrostatic potential and electron density distributions. In the present study, we applied the CBED method to examine nanometer-scale local structures of BaTiO3.

  It is well known that BaTiO3 undergoes successive phase transformations from the cubic paraelectric phase to three ferroelectric phases: tetragonal, orthorhombic and rhombohedral ones. Coexistence of the displacive and order-disorder characters in the phase transformations of BaTiO3 was pointed out from many experiments and theories. However, local structures related to the order-disorder character were discovered neither in crystal structure analyses using neutron and X-ray diffraction nor by TEM observations.

  Using the CBED method, rhombohedral nanostructures were observed in the orthorhombic and tetragonal phases of BaTiO3.4) It was found that the symmetry of the orthorhombic phase is formed as the average of two rhombohedral variants with different polarizations, and that of the tetragonal phase is formed as the average of four rhombohedral variants. These results indicate an order-disorder character in their phase transformations.4) Similar results were obtained in the ferroelectric orthorhombic phase of KNbO3,5) while it was found that the ferroelectric tetragonal phase of PbTiO3 does not have such rhombohedral nanostructures.6)

  We also proposed a combined use of STEM and CBED methods (STEM-CBED method7)) to observe the nanostructures of polarizations, which is schematically shown in Fig. 1. Using the STEM-CBED method, two-dimensional distributions of the rhombohedral nanostructures, or nanoscale fluctuations of the polarization clusters, were successfully visualized in the tetragonal phase of BaTiO3 as shown in Fig. 2.

References

1) M. Tanaka and K. Tsuda, J. Electron Microsc. 60(Suppl. 1), S245 (2011).

2) K. Tsuda and M. Tanaka, Acta Cryst. A 55, 939 (1999).

3) K. Tsuda et al., Acta Cryst. A 58, 514 (2002).

4) K. Tsuda, R. Sano and M. Tanaka, Phys. Rev. B 86, 214106 (2012).

5) K. Tsuda, R. Sano and M. Tanaka, Appl. Phys. Lett. 102, 051913 (2013).

6) K. Tsuda and M. Tanaka, Appl. Phys. Express 6, 101501 (2013).

7) K. Tsuda, A. Yasuhara and M. Tanaka, Appl. Phys. Lett. 102, 051913 (2013).


This study was supported by JSPS KAKENHI Grant Number 25287068.

Fig. 1: (a) Schematic diagram of the STEM-CBED method.7) (b) a STEM-CBED map of the tetragonal BaTiO3 and CBED patterns,7) which shows the intensity difference between the 020 and 0-20 reflections, (I020-I0-20)/I020. The CBED patterns obtained at positions A, B, and C are, respectively, shown in (c), (d), and (e).

Fig. 2:
Type of presentation: Oral

IT-9-O-2948 Direct determination of atomic structures from the observation of phase

Etheridge J.1,2, Nakashima P. N.2, Moodie A. F.1
1Monash Centre for Electron Microscopy, Monash University, VIC 3800, Australia, 2Department of Materials Engineering, Monash University, VIC 3800, Australia
joanne.etheridge@monash.edu

To determine a crystal structure, we need to determine the amplitude and phase of its structure factors from the intensity in its diffraction pattern. However, phase information is either missing or extremely difficult to extract from the diffracted intensities, the infamous “phase problem”. To compensate for this, conventional structure determination methods measure thousands of amplitudes and then deduce the missing phase information using computer-intensive statistical analysis. Although this is time-consuming and the solution is not unique, it has remained the only structure determination approach for a century because of the inability to measure phase.

Here we demonstrate the opposite approach. We show that a centrosymmetric structure can be determined purely from the observation of phase from 3-beam convergent beam electron diffraction (CBED) patterns [1], without the need to measure intensity or analyse it with computer simulations or statistical analysis.

The equations for three beam CBED patterns of centrosymmetric crystals can be inverted analytically, so that the crystal structure factors are described directly in terms of distances to specific features in the pattern [2,3]. This enables the direct measurement of the 3-phase invariant as well as the amplitudes of the structure factors, without recourse to pattern-matching routines [4,5]. Most notably, the sign of the 3-phase invariant can be determined directly by inspection, from the direction of deflection of the rocking curve near the 3-beam Bragg condition (Fig. 1) [4,5], and the individual phases can then be determined from the Bormann effect [6]. This then opens the possibility of solving a crystal structure starting from the observation of phases, rather than the measurement of amplitudes.

We illustrate the method with α-Al2O3, which has 30 atoms in its unit cell. We determine 9 of the structure factor phases, simply from observation of features in 3-beam CBED patterns [1,4,5]. Using these 9 phases only, we can determine the structure to better than 0.1Å precision with no a priori knowledge, except for its space group [1] (Fig. 2). In comparison, the determination of this structure using conventional X-ray diffraction required the measurement of over 2,000 structure facture magnitudes [7].

References

1. P.N.H. Nakashima, A.F. Moodie, J. Etheridge: Proc. National. Acad. Sci. 110 14144 2013.
2. A.F. Moodie, Chem. Scr. 14 21 1978.
3. A.F. Moodie, J. Etheridge, C.J. Humphreys Acta Cryst. A52 596 1996.
4. P.N.H. Nakashima, A.F. Moodie, J. Etheridge Acta Cryst A63 387 2007.
5. P.N.H. Nakashima, A.F. Moodie, J. Etheridge Ultramicroscopy 108 901 2008.
6. G. Borrmann Phys Z 42 157 1941.
7. E.N. Maslen, V.A. Streltsov, N.R. Streltsova, N. Ishizawa, Y. Satow, Acta Cryst B49 973 1993.


The data used in this work was obtained at the Monash Centre for Electron Microscopy. We are grateful to Prof. R. Withers for helpful discussions. This work was supported by the Australian Research Council (DP0346828 and FT110100427).

Fig. 1: An example of a 3 beam CBED pattern. The phase of a crystal structure factor can be determined by inspection of features in such patterns.

Fig. 2: The structure of α-Al2O3 was determined to <0.1Å resolution from the observation of just 9 structure factors phases from 3 beam CBED patterns. No intensities were measured, no computer simulations were required, no statistical analysis is used, no a priori information is needed, other than space group.

Type of presentation: Oral

IT-9-O-3034 3D Electron Diffraction Tomography without limits: structure analysis of a hyper-complex approximant to icosahedral quasicrystal

Oleynikov P.1, Ma Y. H.1, Fujita N.2, Garcia-Garcia J.3, Yoon K. B.4, Tsai A. P.2, Terasaki O.1,5
1Materials and Environmental Chemistry, Stockholm University, Stockholm, Sweden, 2IMRAM, Tohoku University, Sendai, Japan, 3Facultad CC. Químicas, Universidad Complutense de Madrid, Madrid, Spain, 4Department of Chemistry, Sogang University, Seoul, Republic of Korea, 5Graduate School of EEWS, KAIST, Daejeong, Republic of Korea
peter.oleynikov@mmk.su.se

Analyzing the crystal structure of approximants is of vital importance in deriving structural information of building units (or basic clusters) and their arrangements toward icosahedral quasicrystals (IQCs). The acquired knowledge is essential in performing a hyper-space modeling, which is the only feasible way of today to elucidate the structural details of IQC’s. For approximants conventional single-crystal X-ray diffraction can in principle be applied to analyze their atomic structure. However, it becomes quite challenging for the case of approximants to Al-based F-type IQCs, Al-Pd-TM (TM = transition metal) [1]. These approximants often have very large unit cells with lattice constants of over a few tens of Ångströms [2]. A recent study also suggests that, except for the solved case of [2], it is often very difficult to grow single crystals having coherent crystallinity within the width of the incident X-ray beam. It is therefore desirable if the crystal structure can be assessed using electron diffraction from a sub-micron sized crystal domain.
The aim of this study is to assess the possibility of taking the advantage of 3D Electron Diffraction Tomography (3D EDT) [3] in order to solve the crystal structure of the Al-Pd-TM IQC approximant (cubic, s.g. Pa-3, a = 40.54Å). Automated 3D EDT is a fast and efficient technique that has been recently developed by us [3]. It can be used for fast 3D reciprocal space scanning with a given fine step (0.01° – 0.1°) using conventional transmission electron microscopes.
The crystal structure of the individual sub-micron single crystal was determined from the EDT data collected in conventional selected area electron diffraction (SAED) mode using EDT-COLLECT software package [3] on JEOL JEM-2100 FEG CTEM equipped with a single high tilt holder (+/–50°) and Gatan UltraScan 1000 CCD (2048*2048). The acquired data set contains ~2000 unique electron diffraction patterns (exposure 0.5 sec/frame). Reciprocal space coverage was ~90°. The recorded frames were processed using the EDT-PROCESS software package [3] and assembled into a corresponding 3D volumetric representation of reciprocal space (Fig. 1). The crystal structure was successfully determined (Fig. 2) using the direct methods software Sir2011 [4] from the integrated intensities extracted by EDT-PROCESS program.
In this work we show that 3D EDT as a very powerful technique which offers a facile and systematic way to study complex crystal structures.

[1] A.P. Tsai et al, Mater. Trans. JIM 31 (1990), pp. 98-103.
[2] N. Fujita et al, Acta Cryst. A, 69 (2013), pp. 322-340.
[3] M. Gemmi and P. Oleynikov, Z. Kristallogr. 228 (2013), pp. 51-58.
[4] M.C. Burla et al, J. Appl. Cryst. 45 (2012), pp. 357-361.


We kindly acknowledge Swedish Research Council (VR, 1486801), JEOL Ltd., Japan and BK21Plus, Republic of Korea.

Fig. 1: Reconstructed 3D reciprocal space along 001 direction.

Fig. 2: The potential map of the solved structure using direct methods.

Type of presentation: Oral

IT-9-O-3230 Transmission Kikuchi Diffraction (TKD) in SEM

Palasse L.1
1Bruker Nano GmbH, Berlin, Germany
Laurie.Palasse@bruker-nano.de

It is well known that the study of ultrafine grained materials with grain/cell diameters smaller than ~100 nanometers is very difficult or impossible to characterise by Electron BackScatter Diffraction (EBSD) technique. The spatial resolution limitation of the EBSD technique is function of the electron probe diameter and energy as well as the backscattering coefficient of the analysed material. The incident angle between the beam and the specimen surface (~20º) is another critical parameter influencing the highly anisotropic character of the lateral spatial resolution of the EBSD technique.

As an alternative, the recently introduced Transmission Kikuchi Diffraction (TKD) technique is a SEM based method capable of delivering the same type of results as EBSD but with a spatial resolution improved by up to one order of magnitude [1, 2]. And it only requires a commercial EBSD system and a sample thin enough to be electron transparent, e.g. TEM thin lamellae.

The spatial resolution improvement of TKD compared to EBSD will be demonstrated using results obtained by both techniques. Examples on deformed sample as well as orientation contrast images acquired at unprecedented resolution will also be shown.

In addition, we aim to compare the grain size distribution results between the TKD and  the TEM based “Automated Crystal Orientation mapping” (ACOM) techniques in order to evaluate the feasibility of these advanced methods and discussed the parameters influencing the TKD analysis.

References:

[1] R.R. KELLER and R.H. GEISS, Journal of Microscopy, Vol. 245, Pt. 3, pp. 245–251, 2012.

[2] P. W. Trimby, Ultramicroscopy, 120, 16–24, 2012.


Type of presentation: Oral

IT-9-O-3245 Pushing the boundaries of symmetry determination with ‘digital’ electron diffraction

Beanland R.1, Woodward D. I.1, Evans K.1, Römer R.1, Smith K.1, Thomas P. A.1
11Department of Physics, University of Warwick, Coventry, CV4 7AL, UK
r.beanland@warwick.ac.uk

The symmetries in convergent beam electron diffraction (CBED) patterns and their relationship to crystal space groups were first explained almost 40 years ago, and there have been many investigations which have used this to solve crystal structures. The utility of CBED lies in the ability to obtain patterns from regions only a few nm in size, well below that attainable by other methods, sampling perfect crystal that is unaffected by defects or domain structure. Nevertheless, the technique is restricted by small Bragg angles, making it difficult or impossible to apply to materials with closely-spaced spots in a diffraction pattern. Use of computer control to collect patterns at different incidence angles is now relatively straightforward and overcomes this limitation. Capture of many hundreds or thousands of CBED patterns allows reconstruction of ‘digital’ large-angle CBED (D-LACBED) patterns from regions only a few nm in size. The vast increase in information allows previously intractable problems of symmetry determination – particularly for materials with lattice parameters >1nm – to be solved with relative ease. We give several examples, including AgNb7O18, Ca2Mn3O7, polarity measurements in thin PZT films, and polar nanodomains in Na0.5Bi0.5TiO3.
Figure 1 shows [001] diffraction patterns from AgNb7O18. X-ray diffraction showed the material to be orthorhombic with lattice parameters a = 1.4331, b = 2.6151 and c= 0.3836 nm, but was unable to distinguish between four possible space groups: I222, I212121, Imm2 and Immm. Selected reflections from the corresponding D-LACBED pattern, a combination of 2600 CBED patterns, are shown in Fig. 1b. The whole pattern has a vertical mirror but not a horizontal mirror. Opposing dark field patterns with ±g vectors are not equivalent when translated onto each other, demonstrating that the crystal structure is acentric and eliminating the space group Immm. The projection diffraction group of the pattern is therefore m1R, which fixes the point group as mm2. This is consistent with dielectric permittivity measurements which show that is AgNb7O18 is an ergodic relaxor ferroelectric.
Data from, the Ruddlesden-Popper phase Ca2Mn3O7, is shown in Fig. 2. Occasional stacking faults are visible in the HREM image (Fig. 2a) and these were avoided in the collection of D-LACBED patterns. Again, X-ray diffraction is able to limit the possible space groups to a small number of possibilities, in this case Cmcm or Cmc21. The spacing between spots in the SAED pattern (Fig. 2b) is such that no detail is visible in CBED patterns (Fig. 2c). The D-LACBED pattern projection diffraction group is m1R, indicating the lack of a centre of symmetry and confirming the space group to be Cmc21.


Fig. 1: Fig 1. (a) SAED pattern from [001] AgNb7O18. (b) D-LACBED patterns showing a vertical mirror, no horizontal mirror, and acentricity (projection diffraction group m1R).

Fig. 2: Fig 2. [001] Ca2Mn3O7. (a) HREM image of stacking faults; (b) SAED and (c) CBED patterns, (d) selected D-LACBED patterns (projection diffraction group m1R)

Type of presentation: Poster

IT-9-P-1450 Sr25Fe30O77 : A complex layered and modulated structure solved by electron diffraction

Lepoittevin C.1
1Institut Néel, CNRS et Université Joseph Fourier, Grenoble, France.
christophe.lepoittevin@neel.cnrs.fr

These past few years, many new structures have been solved using electron diffraction methods. Zone axis precession electron diffraction (PED) and tomography in reciprocal space are two methods enable to reduce importantly the multiple scattering of the electron beam, so that the reflection intensities can be used for structure determination by direct methods.

The ferrite Sr25Fe30O77 belongs to a family of phases whose structures consist of an intergrowth of m perovskite layers with complex rocksalt type layers [1-2]. Our compound of interest is the member m = 4 of this family and its structure has been solved by combining both electron diffraction methods cited above. This oxide crystallizes in an orthorhombic system with the sub-cell parameters a ≈ b ≈ 5.4 Å and c ≈ 42 Å. The structure exhibits modulation along axis with a modulation vector . Due to the commensurate nature of the modulation, the structure can be described in a supercell with the parameters a ≈ 27 Å, b ≈ 5.4 Å and c ≈ 42 Å. PED patterns were recorded in zone axis with a Spinning Star unit using a precession angle of 2°. The intensities were extracted with CRISP software [3] in “shape fitting” or “integer” modes. The data were then implemented in SIR2008 software[4] and many trials were made with or without application of geometrical Lorentz correction to obtain the structure. The tomography data collection, recorded by tilting manually every 0.5 degree from -30 to +30 degrees, was inserted in EDT (Electron Diffraction Tomography) software [5], which reconstructs the 3D reciprocal space and integrates automatically the reflection intensities. The resulting intensity file was then used on SIR2008 for structure resolution. The solved structure, by combining both methods, consists of four consecutive layers with Fe in octahedral environment alternating with one complex layer containing Fe in three different environments. The oxygen atoms in this last layer are responsable of the modulated nature of the structure.

References:

[1]Pérez, O., Mellenne, B., Retoux, R., Raveau, B. & Hervieu, M. (2006). Solid State Sciences. 8, 431-443, [2]Grebille, D., Lepoittevin, C., Malo, S., Pérez, O., Nguyen, N. & Hervieu, M. (2006). J.Solid State Chem. 179, 3849-3859, [3]Hovmöller S., www.calidris-em.com, [4]Il milione II suite http://wwwba.ic.cnr.it/content/sir2011-v10, [5]Oleynikov P.www.edt3d.com.


Fig. 1: [010] electron diffraction pattern of Sr25Fe30O77

Fig. 2: solved structure of Sr25Fe30O77

Type of presentation: Poster

IT-9-P-1516 Substrate threading dislocations imaged by weak beam dark field TEM on samples with GaN nano-LEDs

Lenrick F.1, Bi Z.2, Ohlsson J.3, Ek M.1, Hetherington C.1, Samuelson L.2, Wallenberg L. R.1
1Centre for Analysis and Synthesis/nCHREM, Lund University, Box 124, S-221 00 Lund, Sweden, 2Solid State Lighting Center, Lund University, Box 118, S-221 00 Lund, Sweden, 3QuNano AB, Ideon Science Park, Sheelevägen 17 , S-223 70 Lund, Sweden
filip.lenrick@polymat.lth.se

The semiconductor material GaN is used in blue and white light emitting diodes (LEDs). It’s also a promising material high power and RF electronics Traditional planar epitaxial fabrication of GaN is, however, not adequate due to the large lattice mismatch between GaN and the available substrates, such as sapphire, Si and SiC. At the strained interfaces threading dislocations (TDs) are formed, degrading efficiency, reliability and lifetime of the devices.
Nano-sized structures show the potential to be free of TDs due to their small dimensions, and morphologies such as nano-wires and nano-pyramids (grown along <0001>) have additional benefits. For instance, the quantum confined Stark effect can be reduced since these morphologies can offer non-polar and semi-polar planes, respectively.
Truncated GaN pyramids were grown by selective area metal-organic vapour phase epitaxy on a GaN substrate with high TD density. A 30 nm thick layer of amorphous Si3N4 (grown by low-pressure chemical vapor deposition) with openings about 100 nm in diameter, patterned by electron-beam lithography and etched by reactive ion etching, was used as the selective area mask. The mask blocks most of the TDs in the substrate from entering the pyramids, but the ones that cross through the mask are interesting to study due to their degrading impact on the device.
To clearly observe the threading dislocations, weak beam dark field (WBDF) transmission electron microscopy (TEM) was applied on focused ion beam (FIB) prepared cross sections. The images facilitate tracing of TDs through the material and how they enter the nano-structures. The FIB lamella, which was about 100 nm thick, showed a TD density of about 10 TDs/µm in projection. Six adjacent pyramids were analyzed where two was found to have TDs from the substrate coming through the mask. The WBDF technique is challenging on a high acceleration voltage microscope due to the low curvature of the Ewald sphere. WBDF condition such as 3g(9g) was found to be more suitable than the standard g(3g) since many diffraction spots are excited. By slightly defocusing the diffraction pattern and using Kikuchi lines as guide lines WBDF conditions became easier to set up.


Fig. 1: 3g(9g) weak beam dark field (WBDF) TEM image of FIB prepared cross sections of truncated GaN nano-pyramid grown through small openings in Si3N4 mask on a GaN (0001) surface. Threading dislocations (TDs) are visible as bright lines. Two TDs marked by red arrows are blocked by the mask, while one TD marked by a green arrow enters the pyramid.

Fig. 2: The low curvature of the 300kV Ewald sphere causes a challenge to set up the WBDF conditions. Kikuchi lines, visible at slight defocus, are usable as guidelines.

Fig. 3: Schematic illustration of one truncated GaN pyramid (grown through openings in amorphous Si3N4 on a GaN substrate) as seen in TEM projection. Threading dislocations (TDs) are marked as red lines. The Si3N4 mask acts as a filter, keeping the TDs in the substrate, but occasionally a TD pass through the opening.

Type of presentation: Poster

IT-9-P-1526 PRECESSED NANO-ELECTRON DIFFRACTION PATTERNS OF THE HUMAN TOOTH ENAMEL CRYSTALS

REYES-GASGA J.1,2, ADDAD A.1, BRÈS E. F.1
1Unité des Matériaux et Transformation (UMET). Université de Lille 1, Sciences et Technologies. Bâtiment C6. 59650 Villeneuve d’Ascq. Lille, France., 2Permanet Address: Instituto de Física, UNAM. Circuito de la Investigación s/n. Cd. Universitaria, 04510 Coyoacán, Mexico D. F., México
jreyes@fisica.unam.mx

In this work we present the precessed electron diffraction patterns of the nano-sized human-tooth-enamel crystallites. These diffraction patterns have allowed us to obtain crystallography information the enamel’s unit cell [1].
The intensity of selected area electron diffraction (SAED) and nano-electron diffraction (n-ED) patterns is difficult to interpret due to the multiple interactions which take place (dynamical diffraction, absorption, etc). However, when the electron beam is tilted and precessed at high frequency the dynamic effect is minimized [2]. The crystal is not moving but the Ewald’s sphere is precessing around the optical axis producing that the dynamical SAED patterns become close to kinematical conditions and they can be used to obtain information on crystal structures [3, 4].
We have obtained the precessed n-ED from human tooth enamel crystals along different zone axes. Human tooth enamel is composed in 95% of hydroxyapatite crystals (HAP, Ca10(PO4)6(OH)2). These crystals are elongated-plate-like of 30 to 60 nm wide and 100 to 200 nm long [5], approximately (figure 1).
The human tooth enamel samples were obtained from permanent non-carious human molar teeth, extracted for orthodontic or periodontal reasons. Samples were prepared in the FIB-FEI QUANTA 200 3D equipment using the two beams system. A Philips CM30 transmission electron microscope with LaB6 filament working at 300 KV was used for TEM observation, the n-ED and the precessed electron diffraction patterns obtaining using a double-tilt holder. The precession of electron diffraction patterns were obtained with the Nanomegas “Spinning Star” equipment. The patterns were recorded on a Gatan “ORIUS” CCD camera using the Digital Micrograph software. JEMS software (version 3.8431U2012) was used for electron diffraction simulation.
Therefore, we have obtained precessed nano-electron diffraction patterns from crystals in the range from 30 to 100 nm (figure 2).

References
1. See papers in Ultramicroscopy, vol.107, issue 6-7, July 2007.
2. R. Vincent, P.A. Midgley, Ultramicroscopy 53 (1994) 271-282.
3. J.P. Morniroli, A. Redjaımia, S. Nicolopoulos, Ultramicroscopy 107 (2007) 514-522.
4. H. Klein Acta Cryst. A67, (2011) 303-309.
5. J. Reyes-Gasga et al., Materials Sci. Eng. C 33 (2013) 4568-4574.


JRG thanks to DGAPA-UNAM (contract IN106713), CONACYT and PASPA-DGAPA-UNAM for sabbatical support.

Fig. 1: TEM bright field image of the human tooth enamel FIB sample used in this work. Note the nano-sized crystals.

Fig. 2: Nano-electron diffraction pattern along the [0001] zone axis (A) and the corresponding precessed electron diffraction pattern (B). C) Simulated [0001] electron diffraction pattern for a HAP sample with thickness sample of 15 nm.

Type of presentation: Poster

IT-9-P-1563 Nano-beam Diffraction of Pt/Al2O3 and Pd/Al2O3 Catalysts

Ward M. R.1, Boyes E. D.1,2, Gai P. L.1,3
1Department of Physics, University of York and the York Nanocentre, UK, 2Department of Electronics, University of York and the York Nanocentre, UK, 3Department of Chemistry, University of York and the York Nanocentre, UK
michael.ward@york.ac.uk

Pt-Al2O3 and Pd-Al2O3 catalysts are used in a wide range of applications including automobile emissions catalysts (1, 2). Although there have been many studies on this system, there are few studies which examine in detail the nanoparticle interface with the complicated nature of γ-Al2O3 and its associated polymorphs. Commercial Pt-Al2O3 and Pd-Al2O3 catalysts tend to be composed of small metal nanoparticles (< 10 nm) on high surface area Al2O3, which are often agglomerations of 10-30 nm crystallites. SAD analysis of such catalysts rarely provides quantitative crystallography of individual Al2O3 crystals but nanobeam diffraction (NDB) has been shown to be a useful tool for examining individual small crystals (3). The nano-sized probe, combined with nanosize crystals has been shown to produce well defined shape effects in the diffraction pattern which can provide further structural insights compared to images alone (4). Here, we have used nano-beam diffraction to investigate how this technique can provide useful insights into the structural relationships between the nanoparticles and the Al2O3 support.

A double aberration corrected JEOL 2200FS was used for this study. The catalysts were provided by Jonhson Matthey as powders. (S)TEM specimens were prepared by depositing an ethanol suspension of the powder onto a holey-C film Cu TEM grid.

Figure 1 shows a typical SAD pattern of an agglomeration of many Pd nanoparticle coated Al2O3 crystals. The bright rings are the {400} and {440} rings assuming γ-phase. The related θ and δ phases have similar bright spatial frequencies but are indexed differently. From the SAD pattern in Figure 1, no useful information can be extracted from it regarding the phase of individual Al2O3 crystals and the nanoparticles supported on them. Figure 2 shows NBD pattern of an individual Al2O3 crystal. The NBD pattern was taken with a probe of approximately 2 nm in diameter. The convergence angle of the probe in this case is sufficiently small to produce a diffraction pattern composed of well defined points rather than discs. The diffraction pattern shows well ordered spots of γ-Al2O3 in [110] orientation. Diffraction patterns such as this from the support and a nearby nanoparticle provide insights into the structural relationship between the two. Where it is applicable, for quantitative analysis the method is clearly superior to FFTs of images in terms of the effects of SNR, aberrations and drift etc and in not requiring an exact zone axis orientation.

References

1. A. Russell, W. S. Epling, Catal Rev 53, 337 (2011).

2. O. K. Ezekoye et al., J Catal 280, 125 (May 16, 2011).

3. M. R. Ward, T. Hyde, E. D. Boyes, P. L. Gai, Chemcatchem 4, 1622 (Oct, 2012).

4. F. Tao et al., Science 322, 932 (Nov 7, 2008).


The authors thank the EPSRC for support from critical mass grant EP/1018058/1.

Fig. 1: (a) TEM image and (b) SAD pattern of the Pd-Al2O3 catalyst. The SAD pattern rings are not useful for indentifying the phase of individual Al2O3 crystals.

Fig. 2: (a) TEM and (b) NBD pattern of a single crystal (indicated) in the Pd-Al2O3 catalyst with (b) being identified as γ-Al2O3

Type of presentation: Poster

IT-9-P-1573 Charge Density Determination for Transition Metals and Intermetallics by Convergent Beam Electron Diffraction and Density Functional Theory Validation

Sang X.1, Kulovits A. K.1, Wang G.1, Wiezorek J. M.1
1Department of Mechanical Engineering and Materials Science, University of Pittsburgh, Pittsburgh, PA, 15261, USA
wiezorek@pitt.edu

The electron charge density difference, ∆ρ(r), i.e., the difference between the crystal electron density and that of the equivalent independent atom model (IAM), represents quantum mechanical characteristics central for fundamental understanding of materials. Convergent beam electron diffraction (CBED) permits probing of nano-scale volumes of perfect crystal and can enable measurements of low-order structure factors, Fg, with sufficient accuracy to obtain ∆ρ(r) for transition metals and binary intermetallic phases [e.g. 1-3]. The high accuracy and precision of the CBED measurements warrants their use as additional metrics in validation of density functional theory (DFT) calculations for these d-electron system materials [2]. Here, sets of multiple Fg and the Debye Waller factors have been determined simultaneously by CBED for transition metals (e.g. Cr, Fe, Ni, Co, Cu, Ta) and chemically ordered intermetallic phases (e.g. NiAl, TiAl, FePd). Using the local density approximation (LDA), LDA + U, and different generalized gradient approximations (GGA) functionals implemented in WIEN2K low-order Fg and thus ∆ρ(r) have been calculated for comparison with the CBED measurements. While many of the different GGA calculations achieve good overall agreement with the experimentally determined low-order Fg for the elements, LDA and GGA functionals fail to predict accurately the low-order Fg for β-NiAl and γ1-FePd. For equiatomic γ-TiAl GGA based DFT achieved considerably improved agreement with experimentally determined ∆ρ(r), when compared with LDA calculations [2]. Fig. 1 shows the difference between the X-ray structure factors, Fg, determined by CBED for two different composition TiAl crystals (Ti-50at%Al and Ti-52at%Al) and the IAM based Fg. Select data from ∆ρ(r)-maps obtained from CBED measurements and GGA DFT calculations are compared in Fig. 2 for the equiatomic and slightly Al-rich TiAl phases. Effects from the small (2at.%) Al-excess in the intermetallic γ-TiAl have been detected by the CBED experiments and are discernible in the ∆ρ(r) most clearly for the (001)-sections (Fig. 2). The excess Al is incorporated substitutionally on Ti sites and appears to enhance delocalization of charge density between second nearest neighbor Ti atoms along <010>, while reducing it for nearest neighbor Ti atom bonds along <110> (Fig. 2).

References

[1] XH Sang, AK Kulovits, JMK Wiezorek, Acta Crystal. A66 (2010) p. 694

[2] XH Sang et al., J. Chem. Phys. 138 (2013) p.084504

[3] XH Sang, et al., Phil. Mag. 92 (2012) p.4408


The authors acknowledge support from the Office of Basic Energy Sciences, Division of Materials Science and Engineering (Grant No. DE-FG02-08ER46545).

Fig. 1: Fig. 1: Difference between the CBED determined X-ray structure factors and the IAM structure factors, ∆Fg, for equiatomic (TiAl) and Al-rich off-stoichiometric (Ti-52Al) γ-TiAl phase. The hkl are plotted along the abscissa. The [uvw] in the legend (inset) indicate the approximate incident beam direction in CBED experiments.

Fig. 2: Fig. 2: Example ∆ρ(r) sections in (001), all Ti plane for the equiatomic composition phase, of L1o-structure tP4 unit cell of TiAl. CBED derived for Ti-50Al (equiatomic) on the left, CBED derived for Ti-52Al (Al-rich) in the middle, and DFT calculated for Ti-50Al (equiatomic) on the right.

Type of presentation: Poster

IT-9-P-1631 Axial transmission electron diffraction in a scanning electron microscope

Volkenandt T.1, Müller E.1, Gerthsen D.1
1Laboratory for Electron Microscopy, Karlsruhe Institute of Technology (KIT), Karlsruhe, Germany
erich.mueller@kit.edu

Scanning transmission electron microscopy (STEM) at low electron energies is a well suited technique to achieve sensitive material contrast in the high-angle annular dark-field (HAADF) mode where contrast is attributed to incoherently scattered electrons. HAADF STEM can be exploited for sample thickness determination and composition analysis [1]. Transmission electron backscattered diffraction (t-EBSD) patterns were recently recorded from a thin specimen by a detector placed laterally to the tilted sample [2]. In our study the detector was placed on-axis below the sample and coherent electron scattering at energies up to 30 keV was analysed which yields axial Bragg-diffraction patterns with Kikuchi lines.
A FEI Strata 400S scanning electron microscope equipped with a segmented semiconductor STEM detector was used. A conventional imaging plate (IP) was inserted below the sample as a detector. The sample consists of a GaN layer with 140 nm thickness on a 120 nm AlN layer epitaxially grown on a Si(111) substrate. A TEM sample with a thickness of 120 nm was prepared by focused-ion-beam milling.
Figure 1 shows a 25 keV HAADF STEM cross-section image of the sample. Dislocations and columnar regions (indicated by dashed lines) with slightly different intensities can be seen in the GaN layer. A tilt series was recorded which shows changes and even contrast inversion within the GaN layer which is a strong indication for coherent scattering.
Figure 2 shows a transmitted on-axis IP-image taken at 25 keV at the position marked by a cross in Figure 1. Figure 2a depicts the illuminated area with the STEM detector segments marked by circles. Kikuchi lines are visible on the whole detector area which can be identified by comparison with simulated EBSD patterns. Diffraction patterns from different positions along the GaN layer show a shift of Kikuchi lines due to orientation changes in the columnar layer. Figure 2b depicts the inner region of the diffraction pattern. A GaN [1-100] zone-axis pattern is identified by measuring the scattering angles for the Bragg reflections. This pattern also yields information on the first-order Laue zone and shows (0002) two-beam excitation condition.
Axial diffraction patterns recorded with IP reveal Bragg reflections and Kikuchi lines within the scattering range covered by the STEM detector. They provide information on the crystal structure of the sample and show that coherent scattering must be considered even at large scattering angles at low electron energies. Moreover, the diffraction pattern shows the local orientation and excitation condition of the sample.

References
[1] T. Volkenandt, E. Müller, D. Hu, D. Schaadt, D. Gerthsen, Microsc. Microanal. 16, 604 (2010)
[2] N. Brodusch, H. Demers, R. Gauvin, J. Microscopy 250, 1 (2013)


This work was funded by the Deutsche Forschungsgemeinschaft (DFG).

Fig. 1: 25 keV HAADF STEM cross-section image with dislocations marked by arrows and columnar regions separated by dashed lines. The cross indicates the position where the diffraction pattern in Figure 2 was taken. The sample was covered with a Pt/C-layer for protection during FIB milling.

Fig. 2: a) Diffraction pattern taken at marked position in Figure 1 at 25 keV. The layout of the STEM detector is indicated by dashed-line circles. b) Inner region of a) showing a diffraction pattern of GaN [1-100].

Type of presentation: Poster

IT-9-P-1694 Strain mapping at the nanoscale using precession electron diffraction in a non-corrected Transmission Electron Microscope

Vigouroux M. P.1, Delaye V.1, Lafond D.1, Bernier N.1, Rouvière J. L.2, Chenevier B.3, Bertin F.1
1CEA, LETI, MINATEC Campus, 17 rue des martyrs, 38054 GRENOBLE Cedex 9, France, 2CEA, INAC, MINATEC Campus, 17 rue des martyrs, 38054 GRENOBLE Cedex 9, France, 3LMGP, CNRS, 3, parvis Louis Néel, 38016 GRENOBLE Cedex 1, France
mathieu.vigouroux@cea.fr

The electron precession [1] technique is a recent innovation in electron crystallography. The advantage of this technique is to minimize the dynamical effect to such an extent that diffraction images can be analyzed using a kinematical approach with minimal user intervention. As a first step we have performed Precession Electron Diffraction (PED) strain measurement on a simple calibration sample paving the way to the strain analysis on more complex devices from micro-electronic.
PED measurements were made using a JEOL-JEM2010FEF non corrected microscope operating at 200 kV. Precession beam scan alignment is performed employing NanoMEGAS’s “DigiSTAR” add-on device. Precession semi-angle was set to 1.44° to take full advantage of PED kinematical behavior. With a probe size as small as 4.2 nm FWHM is obtained on the sample with a convergence of 1 mrad.
The sample we have used is prepared from materials grown by RPCVD on a [001] Si Substrate. It is composed of four 10 nm SiGe layers with different contents in Ge separated by 30 nm of Si and covered with 150 nm Si capping layer. A 8 kV FIB operating voltage was used to provide 50 nm thin parallel-sided lamellae with reduced surface damage. This sample was specifically designed to benchmark strain studies [2] as it is easy to simulate the strain expected in TEM.
PED are recorded every 2.7 nm in a 185 nm x 240 nm region indicated in Fig. 2. using a 512 x 512 pixels Camera deported from the microscope optical axis. Classical projective geometry was used to correct most of distortions in misaligned cameras. Figure 1 (a) illustrates typical diffractions patterns acquired during experiments. Beam probe images were made with a CCD camera (Fig. 1) able to deal with high brilliance scenes and dedicated software was designed to compute PED patterns for strain analysis. The algorithm used takes advantage of the whole “kinematic” region in reciprocal space. The basis of vectors inherent in that periodic region is found using Delaunay triangulation and introduced in a reciprocal matrix G. From this matrix, the distortion matrix D can be retrieved, giving access to the strain ε matrix and rotation Ω matrix.
Figure 2 shows εxx strain mapping obtained by analysing the acquired diffractions patterns set with this method. The noise in 800 contiguous εxx values far from SiGe layers is rather small so that an rms of 3 10-4 is obtained. Strain profiles (Fig. 3) reveal the strong repeatability in measures. Both of them agree very well with finite element COMSOL simulation of the strain averaged along the beam direction and convoluted with the measured electron beam shape.

[1] Vincent, R., et al. « Ultramicroscopy, 53, 3,1994

[2] Rouviere, et al. Applied Physics Letters 103, 24,2013


This study was made possible through funding provided by ANR LABEX MINOS and ANR AMOS programs. Experiments have been done within the Nanocharacterisation Platform of the CEA/Grenoble, MINATEC Campus.

Fig. 1: (a) [110] PED diffraction patterns obtained in Silicon with the probe displays in (b). (a) 4.2 nm full width at half maximum spot size in silicon measured on CCD camera with 1.44° Precession semi-angle.

Fig. 2: SiGe Strain mapping with Precession (semi-angle set to 1.44°)

Fig. 3: εxx SiGe Strain profile along y=70 profile shown Fig. 3; mean εxx strain profile all over y; εxx SiGe Strain profile obtained by finite element COMSOL simulation.

Type of presentation: Poster

IT-9-P-2034 Imaging of grain boundaries in polycrystalline samples by HRTEM

Kiss Á. K.1, 2, Pécz B.1, Rauch E. F.3, Nicolopoulos S.4, Lábár J. L.1
1Institute for Technical Physics and Materials Science, Research Centre for Natural Sciences of the Hungarian Academy of Sciences (MTA TTK MFA), Budapest, Hungary, 2University of Pannonia, Doctoral School of Molecular-and Nanotechnologies, Veszprém, Hungary, 3SIMaP, Grenoble INP/CNRS, France, 4NanoMEGAS Sprl, Brussels, Belgium
kiss.akos.koppany@ttk.mta.hu

Simultaneous imaging of neighboring grains and the grain boundary between them is tedious if polycrystalline samples are to be examined with random orientation distribution of submicron sized grains. The tilting range of HRTEMs is limited to about 20° and there is a low probability to find simultaneously resolved planes and especially low index zones for both grains within this tilting range by chance. Operation of a computer assisted method is demonstrated here that aids such imaging. The method is a combination of the commercial precession electron diffraction (PED) system [1] deployed on a JEOL 3010 with a new computer program that predicts tilt values needed for simultaneous HRTEM imaging of the grains selected from the orientation map.

The best scenario is when we are able to orient low index zones parallel to the electron beam in both grains and the grain boundary is also parallel to the beam simultaneously. A solution with compromise is if only one of the grains is seen from a low index zone while only one plane-set is resolved for the other.

Miller indices of the grain boundary plane in the coordinate systems of both grains are determined from its projection and the local thickness (or from projections at two tilt values as an alternative). The method also comprises the calibration of the directions of the tilt axes in the image.

The evaluation process can be applied to both cubic and non-cubic crystal systems and even to phase boundaries since the calculation of orientations and sample tilts is based on the general metric matrix formalism.

Application of the method is demonstrated here on hcp ZnO thin film with grain size of ca. 20‑40 nm deposited on Si substrate. Figure 1 shows the orientation map of the interested area. Different colors represent different orientations (blue area at bottom is the Si substrate) therefore individual grains can be recognized. The chosen boundary is marked by the white arrow. Figure 2 shows BF image of the layer while Figure 3 presents the HRTEM image of the observed boundary. The area marked by the dashed rectangle indicates a region where the two grains do not overlap, so the boundary is almost in the beam direction here. Fast Fourier transforms of the two grains, shown as inserts, corroborate that both grains are seen from the predicted orientations.

[1] J.Portillo, E.F.Rauch, S.Nicolopoulos, M.Gemmi, D.Bultreys: Precession Electron Diffraction assisted Orientation Mapping in the Transmission Electron Microscope, Materials Science Forum Vol. 644 (2010) pp 1-7 doi: 10.4028/www.scientific.net/MSF.644.1


Z. Baji is acknowledged for the preparation of the ZnO layer by ALD.

Fig. 1: Orientation map; probe size: 10 nm, step size: 5 nm. The observed boundary is marked by the white arrow.

Fig. 2: Bright field image taken at the area of interest. The observed boundary is marked by the white arrow.

Fig. 3: High resolution image of the neighboring grains showing the first grain from [011] i.e. [-1 2 -1 3] zone. The (011) i.e. (0 1 -1 1) planes are only resolved for the second grain. The selected area shows the best insight into the structure of the boundary.

Type of presentation: Poster

IT-9-P-2120 Statistical evaluation of 3D planes in polycrystalline materials from 2D Electron Backscatter Diffraction (EBSD) maps

Jäger A.1, Klinger M.1, Tesař K.1, Malachov M.1
1Institute of Physics AS CR, Na Slovance 2, Prague, Czech Republic
jager@fzu.cz

Scanning electron microscope (SEM) fitted with electron backscatter diffraction (EBSD) detector reached widespread popularity for gaining crystallographic information from a surface of crystalline materials. The main limitation of EBSD technique during two-dimensional mapping is missing depth information. However, in comparison with time-consuming 3D EBSD that require focused ion beam (FIB), 2D EBSD technique needs simpler equipment and easier post-processing.

In this work, uncomplicated statistical approach is presented to find dominant planes such as grain boundaries and fracture planes in bulk polycrystalline materials. The model is based on analysis of intersections of demanded planes with plane of EBSD mapping. Intersection of the two planes generates traces which are further evaluated. For experimental verification metals with hexagonal close packed (hcp) structure were selected; namely magnesium and titanium since they are very attractive for many industrial applications. Data were acquired on SEM FEI Quanta 3D FEG fitted with Hikari EBSD camera. It is shown that the approach combining 2D EBSD mapping with calculations in Matlab software can evaluate the results very well even with moderate amount of experimental data. With this technique dominant planes such as abundant {10-12} <11-20> 86° twin boundaries in wrought magnesium alloy AZ31 (nominally Mg-3wt%Al-1wt%Zn) and preferred fracture plane in duplex phase titanium grade 2 (nominally <0.3wt%Fe, <0.25wt%O, <0.015wt%H) submitted to uniaxial tension at room temperature were successfully analyzed. An example of statistical evaluation in wrought magnesium alloy AZ31 with a number of {10-12} <11-20> 86° twin boundaries is shown in Fig. 1.


The authors would like to appreciate financial support offered by GACR GBP108/12/G043 and MEYS LM2011026.

Fig. 1: Fig. 1: Normals to boundary planes found with the help of the approach. The results correctly show abundant {10-12} <11-20> 86° twin boundaries in wrought magnesium alloy AZ31.

Type of presentation: Poster

IT-9-P-2149 Precession Electron Diffraction Tomography study of new materials derived from Aurivillius phases.

Mouillard-Stéciuk G.1, Boullay P.1, Barrier N.1, Pautrat A.1
1Laboratoire CRISMAT, UMR CNRS 6508, ENSICAEN, 6 Bd Maréchal Juin, F-14050 Caen Cedex 4, France
gwladys.mouillard@ensicaen.fr

Oxides of the Aurivillius family (Bi2O2)2+(Am-1BmO3m+1)2- (A = Ca, Sr, Ba, Pb, … and B= Ti, Nb, W, …) have attracted constant interest in the solid state chemistry community considering both their complex layered structure and their wide range of potential applications. A large number of Aurivillius phases exhibit ferroelectric properties at room temperature and present structural distortions leading to predictable structures and space groups [1]. While their dielectric properties have been intensively studied over past decades, Aurivillius phases have recently proved to also present good potential as semi-conductor photocatalyst [2,3].

In the search for new ferroelectrics derived from Aurivillius phases, we recently found [4] a series of layered materials in the pseudo-binary system Bi5Nb3O15-ABi2Nb2O9 (A=Ca, Sr, Pb, Ba). Preliminary observations made by Transmission Electron Microscopy (Fig. 1) indicate that these compounds exhibit a complex incommensurately modulated structure. Following the procedure described in [5], a (3+1)D structural model was obtained using ab-initio phasing by charge flipping (Superflip) based on the analysis of Precession Electron Diffraction Tomography (PEDT) data (Fig. 2). The (3+1)D structure was further validated by a refinement against powder X-ray diffraction (PXRD) in JANA2006 (Fig. 3).

The new materials possess a layered Aurivillius-type structure with periodic crystallographic shear planes (CSP) leading to the formation of “collapsed” structures with discontinuous (Bi2O2)2+ slabs and perovskite blocks (Fig. 3b) quite similar to what is known in the high-Tc superconductors and related compounds [6]. It appears that the structural difference between the compounds of this series is the length of the collapsed layers, related to the evolution of the modulation vector with the cationic radius A.

Nevertheless, instead of “conventional” Aurivillius phases, where the possibility of non-stoichiometry is mostly limited to a partial substitution of A cations for Bi in the (Bi2O2)2+ slabs, the newly found compounds exhibit a wide compositional stability domain.

Our results define the contour of what appears as a new family of layered perovskite oxides and emphasizing the role of PEDT in the search for new materials.

[1] P. Boullay, G. Trolliard, D. Mercurio, J.M. Perez-Mato and L. Elcoro, J. Solid State Chem. 164 (2002) 252.
[2] H.H. Kim, D.W. Hwang and J.S. Lee, J. Am. Chem. Soc. 126 (2004) 8912.
[3] X. Chen, S. Shen, L. Guo and S.S. Mao, Chem. Rev. 110 (2010) 6503.
[4] G. Mouillard, Master 2 Recherche (2013) Université de Caen.
[5] P. Boullay, N. Barrier and L. Palatinus, Inorg. Chem. 52 (2013) 6127.
[6] M. Hervieu, M.T. Caldes, S. Cabrera, C. Michel, D. Pelloquin, B. Raveau, J. Solid State Chem. 119 (1995) 169.


Fig. 1: a) [0100] electron diffraction zone axis patterns of one representative member of the new layered compounds. b) Enlarged area of a) revealing the existence of a modulation of the form q = αa*+γc*. c) Corresponding HREM image.

Fig. 2: Results for BaBi7Nb5O24, [0100] projection of a 14ax14c: a) Electron density map as obtained from the charge-flipping structure solution procedure. b) Cationic structural model obtained after interpretation of theelectron density map and the addition of discontinuous functions (crenels) (Bismuth red and Niobium green).

Fig. 3: a) Final observed, calculated, and difference plots obtained for the PXRD Rietveld refinement of BaBi7Nb5O24. The black tick marks indicate the main reflections and the green set the satellite reflections. b) [0100] projection of a 14ax14c supercell as obtained from the PXRD refinement.

Type of presentation: Poster

IT-9-P-2158 Multiple scattering in amorphous structures

Mu X.1, Koch C. T.2, Sigle W.1, Neelamraju S.3, van Aken P. A.1
1Max Planck Institute for Intelligent Systems, Stuttgart, Germany, 2Institute for Experimental Physics, Ulm University, Ulm, Germany, 3Max Planck Institute for Solid State Research, Stuttgart, Germany
muxiaoke@gmail.com

Electron diffraction is a convenient technique to study the structure of materials with the advantage of high spatial resolution compared to X-ray diffraction. This fact has recently also increased interest in measuring the pair-distribution function (PDF) of amorphous materials by electron diffraction.[1] However, electrons are likely to scatter multiple times on their path through the sample, due to their strong interaction with matter. Thus, understanding the effect of multiple scattering (MS) on extracting PDFs from electron diffraction is crucial for the quantitative interpretation.

It is generally accepted that for materials possessing a 3-dimensionally isotropic structure subsequent scattering events along the electron path are independent from one another. It implies that MS can be accounted for by a simple convolution.[2] The single-scattering signal should thus be extractable from a diffraction pattern containing the contribution from MS electrons by deconvolution.[3] In our study of amorphous MgF2,[4] we found that the PDF extracted from the deconvolved diffraction pattern does not differ significantly from the PDF extracted from the original experimental data in peak shape and positions, even though there has been a significant amount of MS.

In order to investigate this similarity between the original and the deconvolved data, we used the QSTEM package [5] for simulating a dynamical diffraction pattern of an amorphous structure [6] and extracted the PDF from it. The first multislice simulation (figure 1c) was done to simulate a diffraction pattern from a small model (figure 1a) obtained by molecular dynamics simulation, mimicking single scattering because of the very thin specimen. Another simulation (figure 1d) was done to simulate the diffraction pattern from a supercell being constructed by vertically stacking the original model 20 times (figure 1b), mimicking a 20 times thicker specimen. Figure 2 shows that, except for a reduction in peak height at low frequencies, the diffraction pattern containing MS agrees rather well with the kinematical one. The PDFs (figure 2d) extracted from the MS data and the kinematic data also show no difference in peak shape or position. We finally conclude that, apart from a reduction in peak height, MS has no significant effect on the PDF. Therefore, deconvolution is not necessary in case that correct retrieval of coordination numbers is not important.

[1] D. J. H. Cockayne, Annu Rev Mater Res 2007, 37, 159-187.

[2] G. R. Anstis et al., Ultramicroscopy 1988, 26, 65-69.

[3] J. E. Ankele et al., Z Naturforsch A 2005, 60, 459-468.

[4] X. Mu, Ph.D thesis, TU Darmstadt 2013, 91-94.

[5] C. T. Koch, Ph.D. thesis, Arizona State University 2002.

[6] C. T. Koch et al., Ultramicroscopy 2006, 106, 383-388.


Acknowledgements: The research leading to these results has received funding from the European Union Seventh Framework Programme [FP/2007-2013] under grant agreement no312483 (ESTEEM2).

Fig. 1: Figure 1. (a) A MgF2 cell containing 6150 atoms. (b) Supercell constructed by stacking 20 randomly orientated single cells (shown in a) to mimic the thick material for the dynamical diffraction simulation. (c) Simulated diffraction pattern from the single cell of the model shown in a. (d) Simulated diffraction pattern from the supercell shown in b.

Fig. 2: Figure 2. (a) Profiles of simulated diffraction patterns; (b) structure factors extracted from a; the black dotted line is a 4th-order polynomial function fitted to the red curve; (c) same as (b) but the polynomial has been subtracted from the red curve (d) PDFs obtained by Fourier sine transform of the structure factors in b.

Type of presentation: Poster

IT-9-P-2328 ACOM-TEM analysis of mineral particles ultrastructural organization in bone tissue

Verezhak M.1, Rauch E. R.2, Gourrier A.1 3
1Laboratory of Interdisciplinary Physics, Université Grenoble Alpes / CNRS, Saint Martin d’Hères, France, 2SIMaP laboratory, Université Grenoble Alpes / CNRS, Saint Martin d’Hères, France , 3European Synchrotron Radiation Facility, Grenoble, France
mariana.verezhak@ujf-grenoble.fr

Bone tissue has a complex hierarchical architecture that is self-assembled in order to perform diverse mechanical, biological and chemical functions. At the nanoscale it can be viewed as a composite material made up of two principal components: collagen fibrils of ~ 100 nm in diameter and platelet-shaped calcium phosphate mineral crystals of the 5 x 50 x 100 nm dimensions. The size, shape, organization, orientation and internal structure of mineral crystals has been a matter of disputes since bone sections were first studied by electron microscopy in the 1950’s [1].
Transmission electron microscopy (TEM) shed new light on this problem by allowing the direct visualization of bone structure. However, a lot of difficulties were faced related to image interpretation and to the choice of samples preparation technique. In collaboration with a medical team, we are now able to produce bone sections as thin as 70 nm. We are also currently exploring new bone sample preparation methods than ultramicrotomy as, e.g. tripod polishing and ion milling.
The novel use of the Automated Crystal Orientation Mapping with a TEM method (ACOM-TEM, also known as ASTARTM tool from NanoMEGAS) [2] to study the mineral particles ultrastructural organization in bone tissue with the spatial resolution of 20 nm is reported. The ACOM-TEM method operated in scanning mode and relied on the comparison between the high quality electron diffraction patterns collected at every scan position and the simulated patterns calculated for a given crystal in all possible orientations. This method, therefore, allows crystallographic indexing, high-resolution nanocrystal orientation (~ 1°) and crystal phase mapping.
The mineral particles in bone orientation 2-D mapping was, for the first time, analyzed and the presence of disorder, discontinuity and crystallinity degree variations is discussed. Current results are part of larger project aiming to understand the nanostructural characteristics of bone tissue and to identify key structural markers of pathological human bone [3], providing possible development of new diagnostic and pharmaceutical tools.

References:
1. Robinson R. A., Watson M. L. (1952). Collagen-crystal relationships in bone as seen in the electron microscope. Anatom Rec 114: 383–409.
2. Portillo J., Rauch E.F., Nicolopoulos S., Gemmi M., Bultreys D. (2010). Precession Electron Diffraction assisted Orientation Mapping in the Transmission Electron Microscope. Mater Sci Forum Vol. 644 pp 1-7.
3. Gourrier A., Li C., Siegel S., Paris O., Roschger P., Klaushofer K. and Fratzl P. (2010). Scanning small-angle X-ray scattering analysis of the size and organization of the mineral nanoparticles in fluorotic bone using a stack of cards model. J Appl Crystallogr 43, 1385-1392.


This project is supported by the NanoSciences Fondation (Grenoble, France), through the PhD Excellence Grant Programme for M. Verezhak.

Type of presentation: Poster

IT-9-P-2385 Investigation of dislocation structures by cross-correlation based EBSD mapping and TEM imaging

Kalácska S.1, Groma I.1, Ispánovity P. D.1
1Eötvös Loránd University
kalacska@metal.elte.hu

During unaxial compression of copper single crystals an inhomogeneous dislocation structure develops. With the use of cross-correlation based analysis of electron backscatter diffraction (EBSD) patterns it is possible to map plastic strain variations in deformed polycrystalline samples [1]. In this work this method is applied to visualize the dislocation structures and corresponding distortion fields in Cu single crystals compressed to different levels. The maps created by this method show inhomogeneous cell structure. Furthermore transmission electronmicroscopy is widely used to create micrographs that directly show dislocation arrangement within the sample.

Sample surface preparation plays a key role in creating ideal conditions for both TEM and EBSD measurements. Firstly, we applied various preparation techniques and investigated the efficiency of those methods. We used focused ion beam to create TEM foils of approximately 100 nm thickness. From samples with high dislocation content it's difficult to carve out such lamellas because during the thinning process the foil can spontaneously bend due to the inner stress field. We also made TEM samples with traitional electopolishing and ion polishing processes and compared the resultant TEM micrographs.

Then the distortion maps of the specimen are computed with the cross-correlation technique. This method is capable of detecting changes of the crystal orientation to higher accuracy than the commercial software provided for standard EBSD devices that analyse each EBSD pattern individually. The good qualitative agreement found between the two methods indicate that the cross-correlation method is capable of giving distribution characterization of the cellular dislocation structure. The results measured on the same surface area by cross-correlation based EBSD and TEM methods were compared and evaluated.

Reference:

[1] T.B. Britton and A.J. Wilkinson, High resolution electron backscatter diffraction measurements of elastic strain variations in the presence of larger lattice rotations. Ultramicroscopy 114 (2012) 82-95.


Special thanks to Károly Havancsák, Zoltán Dankházi and Gábor Varga for consultation and valuable suggestions. The help of Alajos Ö. Kovács and János Lábár is also appreciated.

Type of presentation: Poster

IT-9-P-2459 EBSD sample preparation: high energy Ar ion milling

Kalácska S.1, Baris A.1, Varga G.1, Radi Z.2, Lendvai A.2, Dankházi Z.1, Havancsák K.1
1Eötvös Loránd University, 2Technoorg Linda Ltd.
kalacska@metal.elte.hu

EBSD is a versatile tool providing grain size determination, orientation mapping, phase identification and 3D mapping. Since the EBSD information comes from a few tens of nanometers of the specimen surface regions the most critical issue of the EBSD measurement is the surface quality. The surface should be perfectly clean, free of amorphous or deformed surface layer and moreover it should be flat because of the shadowing effect. Lack of these factors can result either no or faded diffraction pattern.
As it is known, the usual mechanical grinding and polishing create an amorphous layer of (1-100) nm thickness on the surface. The commonly suggested colloidal silica polishment continues for hours and can embed residual polishing material in the surface grains. Electropolishing of the surface can also be tried, but this is a difficult and complex procedure, nevertheless in some cases it cannot lead to the desired result.
In the last decades a new surface milling method is spreading. This is based on energetic ion beam milling; the underlying physical process is the sputtering. One direction of this method is the focused ion beam technique (FIB) with ion energies up to 30 keV. The other direction uses near parallel inert gas (usually Ar) ion beams with energy up to 10 keV.
In this poster we present a newly developed Ar ion sample milling apparatus and show how advantageously it can be utilized to produce high quality sample surface. Surface quality development on series of metal samples was investigated using Technoorg Linda's SC-1000 SEMPrep Ar ion milling apparatus. The surface quality of samples was characterized by the image quality (IQ) parameter of the electron backscatter diffraction (EBSD) measurement. Ar ion polishing recipes have provided to prepare a surface appropriate for high quality EBSD mapping. The initial surfaces of samples were roughly grinded and polished. High quality surface smoothness could be achieved during the subsequent Ar ion polishing treatment. The optimal angles of Ar ion incidence and the polishing times were determined for several materials using a FEI Quanta 3D FEG SEM.


Type of presentation: Poster

IT-9-P-2531 Determination of dislocation density by electron backscattering diffraction and X-ray line profile analysis in ferrous lath martensite

Berecz T.1, Csóré A.1, Jenei P.2, Gubicza J.2, Szabó P. J.1
1Department of Material Science and Engineering, Budapest University of Technology and Economics, Budapest, Hungary, 2Department of Materials Physics, Eötvös Lóránd University, Budapest, Hungary
berecz@eik.bme.hu

Ferrous martensite can appear in several forms, such as lath, lenticular and plate, depending mainly on the composition. Among these martensite structures the lath martensite has high industrial significance because of its high strength and toughness. Lath martensite can appear usually in the technologically more important low (and medium) carbon, low cost and low alloyed steels.

The lath martensite morphology exhibits a characteristic multilevel microstructure. A parent austenite grain consists of several packets (the group of laths with the same habit plane). Each packet is divided into parallel blocks and a block is further subdivided into laths. The size of individual martensite laths is very small, therefore they cannot be seen by optical microscopes. There are high angle boundaries between the blocks and packets, while low angle (about 5-10°) boundaries between the single laths.

The strength and toughness of the lath martensitic steels strongly depend on the microstructure through packet and block sizes, as well as the size, shape and arrangement of the laths. The reason of their high strength and toughness is mainly the high angle boundaries between the blocks and packets which hinder the dislocation movements.

Dislocation density in the lath martensitic structure can be determined by both automated electron backscattering diffraction (EBSD) and X-ray line profile analysis (XLPA) method. Dislocations can cause local lattice distortion, which leads to misorientation between individual points in the lattice. Using automated EBSD, the local orientations are determined at individual points in a regular grid on a planar surface of a polycrystalline specimen.

From the difference between the neighboring orientations on planar surfaces the dislocation density can be calculated. XLPA is sensitive to microstrains around the individual dislocations, even if the dislocations are arranged into such configurations which do not yield any misorientation between the different volumes of the crystal. Thus, the dislocation density calculated by XLPA may be different from that measured by automated EBSD.

In our study dislocation densities are determined in individual laths and blocks by EBSD and these results are compared with the dislocation density measured by XRD in ferrous lath martensite.


This work was supplied by the Hungarian Scientific Research Fund (OTKA PD 101028).

Type of presentation: Poster

IT-9-P-2605 Quantitative local structure analysis of nanocrystalline FeAl by electron diffraction

Rentenberger C.1, Gammer C.2, Karnthaler H. P.1
1University of Vienna, Physics of Nanostructured Materials, Vienna, Austria, 2National Center for Electron Microscopy, LBNL, Berkeley, California, USA
christian.rentenberger@univie.ac.at

Profile analysis by X-ray diffraction has been proven to be able to obtain microstructural parameters averaged over a large sample volume (>10µm3). In nanocrystalline materials it is frequently the case that local information is required. This can be achieved by local quantitative analysis based on selected area electron diffraction (SAED). Using the method of PASAD [1] that provides profile analysis of SAED patterns we show that structural parameters can be deduced of volumes on a submicrometer scale (<0.01µm3).

Nanocrystalline B2 ordered FeAl with a mean grain size of about 35nm was made by high pressure torsion (HPT) followed by a heat treatment [2]. The achieved nanocrystalline material was exposed to a further HPT deformation (3 turns, 8 GPa). SEM studies indicate that the deformation occurs inhomogeneously in the form of shear bands. TEM studies were carried out using 200kV.

Fig. 1 shows a bright field image of a nanocrystalline FeAl sample after further deformation by HPT. The complex contrast variations are caused by orientation variations of individual grains and by lattice defects. The darker band in the middle of the image corresponds to a shear band (SB). The density of the dislocations is so high that it is not possible to determine it. Therefore, we use an alternative method. Fig. 2 shows an SAED pattern taken from the encircled area (cf. Fig. 1). The pattern consists of concentric rings. Using PASAD-tools [1] an intensity profile as a function of the diffraction vector g is obtained by integration along the rings (cf. inset Fig. 2). The broadening of the peaks (half-width at half maximum, HWHM) corrected for instrumental broadening was studied by fitting combined Voigt peak-functions. Since broadening by grain size and strain has different effects on the peak profiles both of them can be determined using the method of modified Williamson-Hall plots [3]. This is shown in Fig. 3(a) taking the contrast factors C of dislocations (slip system <111>{110}) into account. The slope of the curve is proportional to the square root of the dislocation density. The values of the slope were calculated from 35 SAED patterns arranged in a 5x7 array within the area indicated in Fig. 1. Fig. 3(b) shows a contour plot of the slope values as a function of the SAED positions. The values indicate that even in a nanocrystalline material the dislocation density within a shear band can be up to a factor 4 higher than in the neighbouring area.

[1] C. Gammer, C. Mangler, C. Rentenberger, H. P. Karnthaler. Scri. Mater 63 (2010) 312.

[2] C. Mangler, C. Gammer, H. P. Karnthaler, C. Rentenberger. Acta Mater 58 (2010) 5631.

[3] T. Ungar, A. Borbely. Appl. Phys. Lett. 69 (1996) 3173.


The authors acknowledge support by the Austrian Science Fund (FWF):[I1309, P22440, J3397] and C.G. by the National Center for Electron Microscopy, Lawrence Berkeley Lab, supported by the U.S. Dept. of Energy under Contract # DE-AC02-05CH11231.

Fig. 1: TEM bright-field image of nanocrystalline FeAl deformed by HPT. Structural parameters were measured by profile analysis of SAED patterns of 35 circular areas placed within the marked rectangle.

Fig. 2: TEM selected area electron diffraction pattern of the encircled area indicated in Fig. 1. The inset shows the corresponding intensity profile.

Fig. 3: (a) Modified Williamson Hall plot obtained from the intensity profile shown in Fig. 2. (b) Contour plot drawn from the slope values of the modified Williamson-Hall plots obtained from a 5x7 array of SAED patterns (of the area marked in Fig.1). The values are proportional to the square root of the dislocation density.

Type of presentation: Poster

IT-9-P-2650 Fluctuation electron microscopy of an amorphous-crystalline composite material

Ebner C.1, Gammer C.2, Karnthaler H. P.1, Rentenberger C.1
1University of Vienna, Physics of Nanostructured Materials, Vienna, Austria, 2National Center for Electron Microscopy, LBNL, Berkeley, California, USA
christian.ebner@univie.ac.at

Fluctuation electron microscopy (FEM) is a TEM technique that allows the characterization of the atomic structure in an amorphous material. It measures the spatial fluctuations in the scattering of electrons arising on a medium-range scale (1-3nm). Here, the FEM technique based on the acquisition of tilted dark-field images was applied to specimens containing nanometer sized crystalline regions embedded in an amorphous matrix.  
Intermetallic Co3Ti with the nominal composition of Co-23at.%Ti was made by mixing Co and Ti of high purity in an induction furnace under Ar atmosphere. The high oxidation tendency of Ti leads to the formation of some small titanium-oxide particles. After annealing at 950°C for ~100h under a static Ar overpressure to achieve the L12 long range ordered phase, the crystalline Co3Ti alloy was rendered amorphous by severe plastic deformation using high-pressure torsion (HPT with 20 turns at 8GPa).
Fig. 1a shows a TEM bright-field image of the Co3Ti sample subjected to HPT deformation. Dark dots (5-10nm in size) within a homogeneous speckle contrast characteristic for an amorphous sample indicate the presence of small crystalline particles. Some of these crystalline particles light up in the tilted TEM dark-field image (cf. Fig. 1b) when a certain scattering vector k is used. Fig. 2 shows the corresponding TEM diffraction pattern of a large area. The dominance of the diffuse rings is characteristic for the amorphous phase. The particles can be analysed by EELS but in this study we want to emphasize the capability of FEM. Therefore, FEM that is sensitive to spatial differences in diffraction was applied. Fig. 3 and 4 show the FEM results calculated from tilted TEM dark-field images taken from the entire reciprocal space by varying the direction and length of k. The images were analysed statistically by calculating the mean and the normalized variance V(k) of the image intensity I(k,r): V(k)=(<I(k,r)2>/<I(k,r)>2)-1, where <> means averaging over sample position r [1]. By averaging <I(k)> and V(k) of images taken with a given k, plots of <I(k)> and V(k) as a function of k are obtained (cf. Fig. 3).  In order to analyse the crystalline particles, V(k) values of two-phase areas V(k)tp are compared with the value of the amorphous area V(k)a. The plot V(k)tp - V(k)a shows peaks at positions corresponding to titanium-oxide lattice planes (cf. Fig. 4). The good agreement of the results by FEM and EELS reveals that FEM is able to identify crystalline particles and it underlines also the applicability of FEM for the characterisation of structural medium-range order in the amorphous phase.

[1] M.M.J. Treacy et al., J. Phys.: Condens. Matter 19 (2007) 455201.


The authors acknowledge support by the Austrian Science Fund (FWF):[I1309, P22440, J3397] and the National Center for Electron Microscopy, Lawrence Berkeley Lab, which is supported by the U.S. Department of Energy under Contract # DE-AC02-05CH11231.

Fig. 1: TEM bright-field (a) and tilted TEM dark-field image (b) of amorphous Co3Ti containing small crystalline particles.

Fig. 2: TEM diffraction pattern and the corresponding intensity profile of amorphous-crystalline Co3Ti.

Fig. 3: Plot of the mean intensity and the normalized variance Va of an amorphous area as a function of k. The position of the first peak in Va indicates the presence of Co3Ti-like structure on a medium-range scale.

Fig. 4: Plot of the normalized variance Vtp – Va. Depending on the orientation of the oxide particles in the taken area (area1-3) different peak positions corresponding to different lattice planes are observed.

Type of presentation: Poster

IT-9-P-2743 Illumination Wavefront Determination by Image and Diffraction Focal Series

McLeod R. A.1,2, Rouviere J.2, Zuo J.1,3
1Fondation Nanosciences, Grenoble, France, 2CEA, INAC/SP2M UJF-Grenoble Minatec campus, 17 rue des Martyrs, 38054 Grenoble Cedex 9, France, 3University of Illinois at Urbana-Champaign, Champaign, USA
robbmcleod@gmail.com

In a transmission electron microscope (TEM), the geometric optics of the illumination system typically are unknown to the user, outside of some basic principles such as condenser lens underfocus or overfocus. When the intensity of a nanoprobe is measured, the phase shift across the probe is lost. The phase shift contains fine oscillations that affect how the probe propagates through the specimen. We present here a method to measure the optical parameters of the illumination system. With an optical model of the illumination mode, an estimate of the probe phase can be found for any lens conditions. The resulting complex entry-wavefunction can then be used for simulation and optimization of the instrument for nanobeam diffraction or coherent diffractive imaging (CDI).

The illumination system is modeled as a single compound lens using the paraxial approximation, with a demagnification of the source and limiting condenser aperture above the lens as shown in Fig. 1. The method calculates the three degrees of freedom: (1) the electron probe diameter b, (2) the convergence angle of the illumination α (or equivalently numerical aperture) and (3) the focal length of the illumination system f (shown in Fig. 2). The dependent parameters, (4) the condenser aperture optical diameter a, and (5) the defocus from specimen to the cross-over zf , are calculated in-addition. The demagnification can be estimated (1⁄M~60) for the given spot size from the nominal aperture diameter. By Fourier optics, the wavefront at the aperture can be numerically forward propagated by za to estimate the complex wavefunction at the specimen.

Our method relies on acquisition of focal series of the nanobeam probe in vacuum via the Python scripting interface. The objective lens excitation is fixed at the eucentric focus. The operating condenser lens, C3 in the case of a FEI Titan, is varied through a large range, forming a series of nanobeam probes at the specimen plane, as shown in Fig. 3. The range is from an image of the condenser aperture conjugate on the specimen plane (C3 = -0.25 in Fig. 3), to the illumination focused on the specimen plane (C3 = 0.02 in Fig. 3). The TEM is then placed in diffraction mode and a series of vacuum diffractograms over both diffraction lens (DL) excitations, and the same range of C3 excitations, is collected (not shown). The diffraction series allows the convergence angle α to be measured, and the combination of both series allows the focal length to be stated in nanometers rather than nominal units. The magnifications in image-mode and camera length in diffraction-mode at each C3 and DL were measured from a second series of images and diffractograms from a monocrystalline Silicon specimen.


RAM acknowledges the financial support of Fondation Nanosciences and CEA.

Fig. 1: The simplified TEM illumination model forms a probe on the specimen of diameter b and convergence angle α. The limiting aperture is demagnified by a factor 1/M. By varying the power of the lens for a series of focal lengths f, the unknowns of the model may be calculated and a conversion from nominal lens excitation to focal length made.

Fig. 2: Focal length of illumination f. The result for the two condenser apertures used, 10 μm and 30 μm, vary only slightly.

Fig. 3: A representative example set for the probe series in imaging mode for the condenser aperture in focus, at C3 = -0.25, to the focal point at C3 = 0.02, and just into overfocus, using the 10 m aperture.

Type of presentation: Poster

IT-9-P-2769 Automated crystal orientation mapping in TEM for the statistical analysis of microstructure evolution in nano-grained polycrystalline thin-films

Aebersold A. B.1, Hébert C.1, Alexander D. T.1
1Interdisciplinary Centre for Electron Microscopy (CIME), École Polytechnique Fédérale de Lausanne (EPFL), Lausanne, CH-1015, Switzerland
arthur.aebersold@epfl.ch

Non-epitaxial polycrystalline films are important in many technological applications. Characterization of their microstructure is crucial for understanding their growth mechanisms and improving their properties. Their microstructure formation typically begins by dense nucleation of randomly oriented grains. During film thickening these grains impinge on each other, resulting in the overgrowth of unfavorably oriented grains and the subsequent formation of a textured film with a columnar or V-shaped grain structure. In this work our aim is to develop a methodology for quantitatively determining the grain size and orientation distributions throughout the thickness of such films, in order to help create accurate simulation models and to correlate film microstructures to their macroscopic properties.
The methodology is based on the principle of orientation mapping (OM); given the nm scale of the film microstructures, in this contribution we apply OM by nano-beam diffraction in TEM using a 2–3 nm electron probe and the NanoMEGAS ASTAR [1]. The high spatial resolution of this technique comes at the cost of needing suitably electron-transparent samples. A quantitative analysis of microstructure evolution requires the sample to have large thin areas from different heights within the film. These requirements can be met by a special double-wedge sample geometry previously proposed by Spiecker et al. [2], which provides continuous plan-view sections throughout the film thickness (see Fig. 1). The heights in the film of the plan-view sections are determined by cross-correlating the position of the thin area to the thickness of the film after dimpling, which was measured from visible light interferences after the dimpling step. Furthermore, the plan-view sections are perpendicular to the direction of the grain elongation thereby minimizing the regions of grain overlap within the projection of the specimen. This improves the reliability of the ASTAR measurements.
Here we report the application of this methodology to nano-grained polycrystalline low-pressure chemical vapor deposited ZnO films used as transparent conductive oxide layers in thin-film solar cells [3]. The applied methodology allows us to extract quantitative in-plane data on the evolution of grain size, orientation, and boundary misorientation as a function of height in the film (see Fig. 2), which can be compared to existing theory and simulations and help to provide new insights into the growth mechanisms that create these films.

References:
[1] EF Rauch et al., Microscopy and Analysis 22 (6) (2008) p. S5.
[2] E Spiecker et al., Acta Mater. 55 (2007) p. 3521.
[3] S Faÿ et al., Thin Solid Films 518 (2010) p. 2961.


The authors acknowledge funding from the SNSF, Grant Number 137833. L. Fanni, Dr S. Nicolay and Dr A. Hessler-Wyser of the IMT PV-lab, EPFL are thanked for the samples and discussions.

Fig. 1: Illustration of double wedge sample geometry. a) ZnO film on glass, b) dimpling of ZnO film (first wedge), c) wedge polish, d) bright-field TEM imaging at different heights along the electron transparent edge (red solid line)

Fig. 2: Inverse pole figure maps overlaid with reliability index, extracted grain size and orientation distributions from two different heights in a ZnO thin film. The data demonstrate the ability of the methodology for obtaining quantitative data on nanocrystalline grain distributions along the height axis.

Type of presentation: Poster

IT-9-P-2894 Analysis of the ordering state of pyroxenes using precession electron diffraction.

Jacob D.1, Wouossaju S.1, Palatinus L.2
1UMET, UMR 8207 CNRS-Université Lille 1, Villeneuve d’Ascq, France, 2Institute of Physics of the Academy of Sciences of the Czech Republic, 182 21 Prague, Czech Republic
damien.jacob@univ-lille1.fr

The precession electron diffraction (PED) technique [1] has been originally developed for structure determination at a submicrometer scale in a transmission electron microscope (TEM). Since, many structures have been solved using PED, recently combined with the tomographic acquisition of 3D electron diffraction data [2]. Using PED, integrated intensities of the diffracted beams as a function of the rocking beam orientation are collected. The resulting intensities keep dynamical in nature, due to residual multiple scattering, but are more closely related to the strength of the scattering events and ranking of reflections as a function of their intensities is generally correlated to the structure factor values, which is crucial for structure solution.

Recently, it has been shown that PED could also be used for structure refinement [3]. In this case, experimental intensities have to be compared with dynamical simulations of diffracted intensities, taking into account the multiple scattering occurring when the electron beam is passing through the crystal. Applied to structures with mixed occupancies, the analysis can be used to refine atomic occupancies of specific sites of the structure, giving access to the ordering parameter. In the field of mineralogy, the PED refinement has thus been used to analyze the ordering state of orthopyroxene (OPX) samples. Results have enabled the distinction between an equilibrated sample (natural OPX (Mg0.60Fe1.40)Si206) and a non equilibrated one (heat-treated (1000°C, 48h) and quenched sample from the same origin), giving ordering parameter values in good agreement with those obtained at the grain scale using XRD [4].

To go further and use PED data to decipher the thermal history of the sample with sufficient precision, the sensitivity of the PED refinement method still appeal for a detailed quantitative evaluation. In this work we discuss the influence of experimental parameters such as the irradiation dose and/or heating of the sample under the electron beam. Analyses are performed on the previously studied equilibrated OPX sample. Our results show a noticeable evolution of the ordering parameter with the electron beam irradiation duration (Fig. 1), which assesses for the high sensitivity of the technique. Possible evolution of the ordering state associated with the in-situ heating of the sample will also be explored, opening the road to the study of intra-crystalline diffusion kinetics at a very local scale in a TEM using PED. [1] R. Vincent and P.A. Midgley, Ultramicroscopy 53 (1994) p. 271. [2] U. Kolb et al., Crystal Research and Technology 46 (2011), p. 542. [3] L. Palatinus et al. Acta Crystallographica A (2013), 69(2), P. 171. [4] D. Jacob et al., American Mineralogist 98 (2013) p.152.


We gratefully acknowledge C. Domeneghetti (Univ. Pavia) and F. Camara (Univ. Torino) for supplying the OPX samples together with their XRD structural analysis

Fig. 1: Plot of XFe(M2) vs. XFe(M1) in a natural OPX sample as obtained from PED dynamical refinement as a function of the duration of the electron beam illumination (200kV, LaB6 Tecnai 20 microscope). Dashed line corresponds to the constant composition line. Green triangle corresponds to XRD results obtained at the grain scale

Type of presentation: Poster

IT-9-P-2929 Application of Large-Angle Convergent-Beam Electron Diffraction to APBs Recognition

Jezierska E.1
1Warsaw University of Technology, Faculty of Materials Science and Engineering, Woloska 141, 02-507 Warsaw, Poland
jezierska@op.pl

The Large-Angle Convergent-Beam Electron Diffraction (LACBED) technique was proposed by Tanaka in 1980 to improve the quality of the CBED patterns obtained with a large angle convergent incident beam (Kossel patterns) [1]. In this method a specimen is raised (or lowered) from its usual eucentric position in the object plane. The LACBED technique which uses a defocus incident beam has a unique property: the image of the illuminated area of the specimen is superimposed on the diffraction pattern composed of Bragg lines [2]. Therefore, the pattern is a mapping between the direct and the reciprocal spaces and “shadow image” of a defect is visible on the pattern.
TEM investigations were performed on JEM 3010 Jeol equipped with Gatan CCD camera. Conventional TEM studies and LACBED were used to elucidate the structure of ordered intermetallics. The antiphase domains structure in various ordered intermetallics with perfect L12 superstructure has been examined. For advanced studies of the nature of antiphase boundaries (APBs) LACBED method was employed.
Ordering of atoms occurs in a large number of alloys. Whereas in the disordered form the lattice sites are occupied at random, they will be occupied by atoms of a given chemical species in the ordered form. Ordering is accompanied by domains formation. The arrangement of domains is characteristic for ordered alloy and applied technology. Due to phase contrast the visibility of APBs is significant on TEM images. Using centered superlattice dark-field image the mapping of ordering can be achieved (Fig. 1). Perfect symmetry was confirmed from LACBED images and ordering is manifested in superlattice Bragg lines (Fig. 2). Any defects breaking the translational symmetry and perfect order can be visible on LACBED lines. For antiphase domain boundaries (APBs) in ordered compound the splitting of superlattice Bragg lines on LACBED images can be observed (Fig. 3). The superlattice excess line is split into two lines with equal intensity on bright-field LACBED pattern as well as on dark-field LACBED pattern if the domains are enough large to see the effect (Figs. 3-4). This splitting can be considered as typical and used to identify APBs. For very fine domains only subtle effect of affected Bragg lines can be noticed [3].
References:
[1] M. Tanaka, R. Saito, K. Ueno, Y. Harada, LACBED, Journal of Electron Microscopy, 29 (1980) 408-412.
[2] J.P. Morniroli, Large-Angle Convergent-Beam Electron Diffraction (LACBED). Applications to crystal defects, Sfμ , Paris (2002).
[3] E. Jezierska, J.P. Morniroli, Antiphase boundaries in Ni3Al ordered intermetallic – application of CBED method, Material Chemistry and Physics 81 (2003) 443-447


The financial support from the Polish Ministry of Science and Higher Education, Faculty of Materials Science & Engineering Warsaw University of Technology is gratefully acknowledged.

Fig. 1: TEM image of antiphase domains boundaries in (Al,Mn)3Ti ordered intermetallic phase with L12 superstructure (centered superlattice dark-field with 011 operating spot)

Fig. 2: LACBED image from perfect (Al,Mn)3Ti superstructure with [233] zone axis

Fig. 3: LACBED on antiphase domains boundary. The superlattice excess line is split into two lines with equal intensity on bright-field LACBED pattern

Fig. 4: DF LACBED of superlattice (01-1) Bragg line with splitting due to APB

Type of presentation: Poster

IT-9-P-2930 Quantitative CBED in a Nano-structured Material

Nakashima P. H.1, Bourgeois L.1,2, Etheridge J.1,2
1Department of Materials Engineering, Monash University, 3800 Victoria Australia, 2Monash Centre for Electron Microscopy, Monash University, 3800 Victoria, Australia
philip.nakashima@monash.edu

When rapidly quenched to room temperature from just below its melting point, aluminium can form octahedral voids of a few tens of nanometres in size, truncated with {001} facets. For convergent beam electron diffraction (CBED), this presents an interesting scenario that can be thought of as a “CBED sandwich”. For a focussed electron probe incident on a {001} void facet, the resultant CBED pattern is the product of diffraction from two totally coherent slabs of crystal, oriented along <001>, each slab sandwiching the free space in the void. Such a CBED pattern is not only sensitive to the thicknesses of the two slabs but also their separation across the void because the electron waves modified by the first slab of crystal then Fresnel propagate as they traverse the void.

In addition to highly constrained measurements of the thickness of the specimen, the dimension of the void in the beam direction and its position with respect to the entrance and exit faces of the specimen, quantitative CBED is used here to measure bonding-sensitive structure factors on either side of the void.

To investigate the sensitivity of the bonding electron density to the nanoscale geometry and size of these structures, we compare these results with recent work [1] where the same structure factors were measured with sufficient accuracy and precision as to be able to unequivocally determine the bonding electron density in aluminium. Our work takes advantage of the multislice formalism for electron scattering [2], which is conducive to the geometry of the “CBED sandwich” that a void in a metallic foil presents.

References:

[1] P.N.H. Nakashima, A.E. Smith, J. Etheridge, B.C. Muddle, Science 331 (2011), 1583.

[2] J.M. Cowley, A.F. Moodie, Acta Cryst. 10 (1957), 609.


The data collected for this work was obtained using the JEOL 2011 TEM in the Monash Centre for Electron Microscopy, funded by the ARC (RIEFP 99). PN thanks the ARC for grant FT110100427.

Type of presentation: Poster

IT-9-P-3029 Strain Analysis by Nano-Beam Electron Diffraction (SANBED) in semiconductor nanostructures

Mahr C.1, Müller K.1, Erben D.1, Schowalter M.1, Zweck J.2, Volz K.3, Rosenauer A.1
1Universität Bremen, Otto-Hahn-Allee 1, 28359 Bremen (Germany), 2Universität Regensburg, Universitätsstraße 31, 93040 Regensburg (Germany), 3Philipps Universität Marburg, Hans-Meerwein-Straße, 35032 Marburg (Germany)
mahr@ifp.uni-bremen.de

The measurement of lattice strain is an important aspect in the characterisation of semiconductor nanostructures. As strain has large influence e.g. on the mobility of charge carriers, methods for accurate strain measurement with high precision are mandatory. In the present work [1] we measure strain from the positions of diffraction discs in convergent-beam electron diffraction (CBED) patterns using dedicated algorithms. Large one- and two-dimensional series of CBED-patterns (~3000) in semiconductor nanostructures have been recorded at an FEI Titan facility in STEM mode. Contrary to parallel illumination in conventional Nano-beam electron diffraction (NBED), we show that focusing the beam with a semi-convergence of 2.6 mrad increases the spatial resolution drastically by a factor of 5 to be 0.5 – 0.7nm. We determined the precision of this method to be 0.07%.

The rich inner structure of CBED-discs causes a big challenge to recognize their positions accurately. In particular, three different algorithms have been developed: As shown in figure 1, the first algorithm detects edges around each disc and iteratively deselects erroneous edges by circle-fitting. In this way the fit converges to the disc boundary. A disadvantage of this parameter-free method is a long computation time. An improvement of speed by a factor of 15 is achieved with the Radial Gradient Maximisation method. This method positions two sets of rings around the initially estimated disc position, one set of rings with smaller radii than estimated and one with larger radii, as illustrated in figure 2. Disc position and radius are determined by maximising the difference between the integrated intensity on inner and outer rings. The third method is a cross-correlation with different masks. As it is nearly a factor of 100 faster than the edge detection algorithm it is capable for in-situ strain measurement during CBED pattern acquisition. The left part of figure 3 shows two different masks. Mask A assumes fully illuminated CBED-discs, whereas with mask B the inner structure of the diffraction discs has less influence on the result. The right part of figure 3 shows the change in the correlation function if the disc is shifted. The disc position can be determined from the shift between mask and experiment. We compare the precision of the three different algorithms among each other and with respect to former approaches [e.g. 2, 3]. Finally we show by application, that specimen cooling, zero-loss energy filtering and aberrations of the projection system only weakly affect the measured strain.

[1] K. Müller and A. Rosenauer et al., Microsc. Microanal. 18 (2012), p. 995.

[2] F. Uesugi et al., Ultramicroscopy 111 (2011), p.995.

[3] A. Béché et al., Appl. Phys. Lett. 95 (2009), p. 123114.


This work was supported by the Deutsche Forschungsgemeinschaft (DFG) under contracts numer RO2057/8-1, SCHO1196/3, V0805/4 and V0805/5.

Fig. 1: Disc position and radius recognition with selective edge detection. (a) Process from experimental raw data to fitted disc position. (b) Edge point deselection: Beginning with the first circle-fit to all edge points, the point with the largest distance to the fitted circle is deleted until only edge points on the disc boundary remain.

Fig. 2: Disc position and radius recognition with Radial Gradient Maximisation. Two sets of rings are positioned inside and outside an initial radius estimation and the intensity is integrated. Radius and position are fitted via maximisation of the difference between both sums.

Fig. 3: Disc-position recognition via cross-correlation with masks. Left: Masks for correlation. Right: Changes in correlation-function (down) when the experimental disc shifts. Disc-position recognition from shift between experiment and mask.

Type of presentation: Poster

IT-9-P-3080 Aberration-compensated large-angle rocking convergent-beam electron diffraction (LARCBED)

Koch C. T.1, Ozsoy Keskinbora C.2, Mu X.2, van Aken P. A.2, Ishizuka K.3
1Institute for Experimental Physics, Ulm University, Albert-Einstein-Allee 11, 89081 Ulm, Germany, 2Stuttgart Center for Electron Microscopy, MPI for Intelligent Systems, Heisenbergstrasse 3, 70569 Stuttgart, Germany, 3HREM Research Inc., 14-48 Matsukazedai, 355-0055 Higashimatsuyama, Japan
christoph.koch@uni-ulm.de

Convergent beam electron diffraction (CBED) is a very efficient technique for acquiring two-dimensional rocking curve information in a single exposure. This is possible, because, for crystal structures having very small unit cells, the space between the Bragg spots is large enough to provide space for many non-overlapping diffraction patterns. If the sample is thick enough (typically > 100 nm), the dynamical diffraction conditions between these diffraction patterns differ enough to produce strong variations in the diffraction intensities. For materials with larger unit cells, such rocking curves must be acquired sequentially, because the distance between reflections is much smaller. Also thin crystals (e.g. nanocrystals) require a much larger tilt range than thick crystals, in order for intensity variations to be significant [1]. For thin crystals with small unit cells one may therefore acquire many CBED patterns, each with a different beam tilt, and combine them to large angle CBED (LACBED) patterns [2]. However, at large beam tilts, care has to be taken to compensate for movement of the probe on the sample due to aberrations of the objective pre-field lens. Also, imperfections in the separation between beam tilt and shift will become significant at large beam tilts. Aberrations of the imaging system add to the complexity of precisely localizing the probe on the specimen. However, this problem has been solved by the commercially available QED plug-in for DigitalMicrograph (Gatan Inc.) [1,3] which allows for automated calibration and compensation of all aberrations up to 7th order.

Here we present results of acquiring large-angle rocking-beam electron diffraction (LARBED) patterns using the QED plug-in with a convergent probe. Fig. 1 shows two example CBED patterns from the data stack (Fig. 1a and c), as well as the sum of all CBED patterns in the stack (Fig. 1b). Note that the beam tilt range (radius of ‘beam tilt disc’) was 60 mrad in each direction. At such large beam tilts the central spot of the pattern would be outside the detector area if no de-scan (compensation of beam tilt by diffraction shift) would be have been applied.

Fig. 2 shows bright-field and dark-field LACBED discs with a diameter of 120 mrad (6.9°) that have been extracted from the data stack shown in Fig. 1 by simply placing each CBED disc at the position of the pattern where it would have been recorded without applying any diffraction shift.

[1] C.T. Koch, Ultramicroscopy 111 (2011) 828 – 840

[2] R. Beanland, et al., Acta Cryst. A69 (2013) 427–434

[3] http://www.hremresearch.com


The research leading to these results has received funding from the European Union Seventh Framework Programme [FP7/2007-2013] under grant agreement n°312483 (ESTEEM2) and the Carl Zeiss Foundation.

Fig. 1: CBED patterns of SrTiO3 acquired using QED. a) [-110] zone axis pattern, b) sum of all patterns acquired in this experiment (data stack), and c) one of the CBED patterns acquired at high beam tilt. Aberrations of the illumination system have been compensated up to 5th order.

Fig. 2: Bright-field (BF) and dark-field (DF) LACBED discs extracted from the data stack shown in Fig. 1. Similar data can be extracted from any of the more than 70 reflections shown in Fig. 1b.

Type of presentation: Poster

IT-9-P-3362 EBSD and EDS Characterisation of Vanadium rich β Phase Lamella in Advanced Titanium Alloys

Stephens C. J.1
1Thermo Fisher Scientific
chris.stephens1@thermofisher.com

EBSD has been the focus of significant interest in recent years, driven by advances in detector technology, computational power and indexing routines. The high spatial resolution of EBSD enables structural characterization of materials on the nanoscale, giving instant quantitative information such as orientation, phase and texture. This work discusses developments in EBSD and EDS for the characterisation of submicron grains within advanced titanium alloys, which have important applications in the aerospace sector. High resolution EBSD patterns are used to characterise the propagation of BCC lamella through HCP subgrains, within a larger HCP matrix. Such a system becomes an ideal tool to characterise the structural properties of advanced materials, where the size and orientation of the lamella are related to the formation processes.

EBSD structural information is cross-correlated with chemical information obtained simultaneously through energy dispersive x-ray spectroscopy (EDS), tracking the migration of trace transition elements towards grain boundaries. The detection of these additives presents a challenge for EDS analysis, due to severe overlaps and low concentrations. These are overcome through intelligent peak deconvolution routines and image filtering, with proprietary principal component phase analysis algorithms used to determine chemically unique phases at overall concentrations below 1%.

Ti alloys can exist in three phases, α, α + β and β. At lower temperatures, pure Ti is stable as a close-packed hexagonal crystal structure and at high temperature it undergoes allotropic transformation to the BCC phase. Al and V are respective stabilisers of the α and β phases, therefore accurate quantification is essential. The alloy described here is stable in the α +β phase at room temperatures and pressures, whereby slow to intermediate cooling rates allow the formation of α colonies within the β phase.

Figure 1 (left) shows a high resolution forescatter image of the α phase alongside α + β phase, revealing the different topographical natures of these regions. Figure 1 (right) reveals the BCC region overlaid onto the image quality (IQ) map, showing a distinct lamella structure. EDS analysis of V segregation within the BCC phase is shown in figure 2 (left), with the V rich beta phase again highlighted. The quantification of V requires peak deconvolution due to the low weight percentage and the overlap with the Ti spectrum. The contoured quantitative weight percentage map of the same grain is shown in figure 2 (right), in 0.8% step sizes up to a maximum concentration of 8.07% showing a migration of V towards the phase boundary. This example demonstrates the benefits that intelligent EDS can bring to EBSD analysis of complex materials.


Fig. 1: left: Forescatter SEM image showing a high degree oftopography in the upper left side associated with the α +β phase and a smoother region in the lower rightside of the image assosciated with the a phase. RightBCC Euler map overlaid onto the pattern quality map highlighting theposition of the BCC lamella within the α + β phase

Fig. 2: Left: Quantitative elemental EDS map of vanadium, associated with the β-phase, overlaid onto the forescatter SEM image. Right: Quantitative EDS map in gradients of 0.8 %, revealing a higher concentration of V at the phase boundary

Type of presentation: Poster

IT-9-P-3490 Composite crystal structures of MxCuO2 cuprates; (Mx = Li2, Ca.83, Sr.73, Ba.67, [(Sr/Ca)2Cu2O3]1/√2, Na)

Milat O.1, Salamon K.1, Demoli N.1
1Institute of Physics, Bijenička 46, HR 1000 Zagreb, Croatia
milat@ifs.hr

The MxCuO2 cuprates belong to class of composite crystals consisting of two subsystems[1]: „CuO2-chains“ and „Mx-cations”. An electron microscopy and diffraction study of a number of rare earth cuprates, is presented (Fig.1.) in relation with the corresponding charge ordering that is induced by variable cation valency and nonstoichiometric composition. Level of cation defficiency and the accompanying self doping affects structure modulation, as well as the average Cu-valency; it can range from Cu+2.30 for Ca.83CuO2, to Cu+2.66 for Ba.67CuO2, and even up to Cu+3.0 for NaCuO2, or down to Cu+2.0 for Li2CuO2. In the case of Mx = Ca.83, Sr.73, Ba.67, [(Sr/Ca)2Cu2O3]1/√2, the two subsystems are mutually incommensurate and modulated along the “chain” direction, while for the end cases of: Mx = Na, Li2, the structural unit cells are commensurate (Fig. 2.)

The underlying lattices of these subsystems have common a and b parameters while the ratio cCh/cM of their c-parameters along the chain-direction varies with x. For a particular case of Mx=[(Ca/Sr)2Cu2O3]1/√2, the so-called “chain-ladder” (Sr/Ca)14Cu24O41 compound is well known for its optical and magnetic properties[2],[3]. In this case, the cation subsystem consists of an extended structure: (Sr/Ca)2Cu2O3 -“ladders”. The building unit of the ladders is a pair of cations plus a pair of zigzag edge-sharing CuO4-squares, Fig. 2, that are connected along “rungs”, so that the cLd period is defined by the CuO4-square diagonal. For the chains, the CuO4 building units share opposite edges and the cCh period is equal to the CuO4-square edge. In the case of “chain-ladder” composite structure, the cLd/cCh ratio is found to vary slightly with cation composition, but is always close to √2, so that the formula [(Ca/Sr)2Cu2O3]xCuO2 (x≈1/√2) correctly represents compound's composite structure.

With increasing Ca-substitution the cLd/cCh ratio varies from 1.44 for pure Sr14Cu24O41, to 1.416 for highly substituted Sr0.6Ca13.4Cu24O41. This is accompanied by charge (hole) redistribution between the CuO2-chains and the Cu2O3-ladders[3].

[1] Milat O., Van Tendeloo G., Amelinckx S., Mebhod M., Deltour R. (1992), Acta. Cryst., A48, 618-625.

[2] Vuletić T., Korin-Hamzić B., Ivek T., Tomić S., Gorshunov B., Dressel M., Akimitzu J. (2006), Phys. Rep. 428, 169-258.

[3] Ilakovac V., Gougoussis C., Calandra M., Brookes N. B., Bisogni V., Chiuzbaian S. G., Akimitsu J., Milat O., Tomic S., Hague C. F. (2012), Phy. Rev. B 85, 075108.


S. Tomić from Zagreb (Croatia), A. Migliori from Bologna (Italy), and G. Van Tendeloo from Antwerp (Belgium) are acknowledged for collacoration and support. Financial support has been received from Croatian Ministry of Sciences, Education and Sport.

Fig. 1: EDP of MxCuO2 composite crystals for Mx= Ca.83 (a), Sr.73 (b), [Sr/Ca2Cu2O3]1/√2 (c), along the [1000] zones perpendicular to the “CuO2-chains”. Indexing in 4-D crystallography notation; the third index is for the “chain-”, the fourth one for the “cation-sublattice”. Two subsystem unit cells are indicated; mismatch corresponds to: (1-x)c*Ch.

Fig. 2: Schematic representtion of the composite crystal structures of Ca.83CuO2, (Sr/Ca/La)14Cu24O41, (with incommensurate “CuO2-chain” modulation), and NaCuO2, Li2CuO2 (with commensurate superlattices); in top view (up), and front view (down).

Type of presentation: Poster

IT-9-P-5870 Error Analysis of the Crystal Orientations and Misorientations obtained by the Classical Electron Backscatter Diffraction Method

Ram F.1, Zaefferer S.1
1Max-Planck-Institut für Eisenforschung GmbH, Düsseldorf, Germany
ram@mpie.de

The orientations obtained by the classical two-dimensional Hough transform-based EBSD method are accompanied by error and imprecision. These measures gain importance when the retrieved orientations and misorientations are used as an input for further analysis --- e.g., in grain boundary analysis, dislocation density analysis, and microstructure-based crystal plasticity modeling. In this contribution, the extent of the error and precision of the retrieved orientations and misorientations will be presented. They are examined using descriptive statistical analysis applied to patterns simulated based on the dynamical theory of electron diffraction. For a ~1 Mpixel pattern reduced to 240×204 pixels, subjected to the two-dimensional weighted Hough transform with 0.25° resolution, and convoluted by a butterfly mask of 13×13 pixels dimensionality, the error in the retrieved orientation is 1°, and the precision of the retrieved orientation is 0.7°. The error in the retrieved misorientation is 1.6° and its precision is 0.2°. The accuracies of the retrieved orientations and misorientations are obtained analytically through the model-based inferential statistics. The error in the detected lattice planes is assumed to have a Fisher-von Mises distribution. To estimate the accuracy, this error is propagated to the retrieved orientation and misorientation. The result is a confidence region for the true orientation or misorientation. The accuracies are defined based on their corresponding confidence regions. It is shown that the obtained accuracies are reliable upper bounds for the corresponding error. The maximum level of over/underestimation is 0.3° for orientation; it is 0.9° for misorientation.

1. Krieger Lassen, N. C., D. Juul Jensen, and K. Conradsen. 1994. “On the Statistical Analysis of Orientation Data.” Acta Crystallographica Section A Foundations of Crystallography 50 (6) (November 1): 741–748.


2. Krieger Lassen, N. C. 1996. “The Relative Precision of Crystal Orientations Measured from Electron Backscattering Patterns.” J. Microsc. (Oxford, U. K.) 181 (1): 72–81.


Fig. 1: Schematic drawing of the error, precision, confidence region, and accuracy of a quantity.

Type of presentation: Poster

IT-9-P-5883 Indexing electron diffraction patterns from randomly-oriented crystals

Wang Y. C.1, Wan W.1, Zou X. D.1
1Stockholm University, Stockholm, Sweden
yunchen.wang@mmk.su.se

Electron diffraction (ED) can be used to study nanometre-sized crystals and provides information about the unit cell, space group and even intensities for a complete structure solution. With a known unit cell, ED patterns can be indexed. The reflection indices are found after indexing and the reflection conditions can be used to derive the space group. Indexing is usually done with in-zone ED patterns from aligned crystals. Alignment of the crystal can be time consuming and for many electron beam sensitive materials it is very difficult, even impossible. Indexing ED patterns from randomly orientated crystals is thus necessary and it gives valuable information about the phase of the material, the space group and possibly quantitative intensities of the reflections. Similar problem has been studied in femtosecond X-ray crystallography using X-ray free electron laser [1]. However the algorithms cannot be adopted for electron data due to the differences in wavelengths. Here we propose a fast and robust approach for indexing ED patterns from randomly orientated crystals, given the unit cell parameters. Rather than measuring the d* values and mutual angles of the reflection vectors and matching them to the calculated values, we use instead the “difference vectors”, which are obtained by subtracting the 2D coordinates of pairs of diffraction spots. Difference vectors typically have shorter d* values than the original reflection vectors and are thus easier to index (less ambiguity). Angles between difference vectors are also included in the indexing to identify a unique solution. Zone axes given by the difference vectors are close (usually within 1-2°) to the actual zone axes along which the ED patterns are taken. They can be used to quickly narrow down the solution for indexing the original reflection vectors, resulting in a quick and reliable solution.

[1] H. N. Chapman, et al., Nature 470, 73(2011)


This project is supported by the Knut & Alice Wallenberg Foundation through the project grant 3DEM-NATUR and a grant for purchasing the TEM.

Fig. 1: Illustration of the “difference vector” approach for indexing ED patterns from randomly orientated crystals. The ED pattern was taken from a silicalite-1 crystal in a random orientation. Three reflections, (-7,1,-3), (-5,0,-1) and (-3,2,-4), generate difference vectors (4,1,-1) and (2,-1,2) which are of lower resolution and easier to index

Type of presentation: Poster

IT-9-P-5973 In situ observation of reverse transformation in steels using EBSD measurement at elevated temperature

Masaaki Sugiyama and Akira Taniyama
Innovative Structural Materials Association (ISMA), Futtsu branch located in Nippon Steel & Sumitomo Metal Corporation
sugiyama.88p.masaaki@jp.nssmc.com

In the advanced high strength steels, the phase transformation from a high temperature phase of austenite with fcc structure to the low temperature one of ferrite with bcc structure, including of bainite or martensite phases must be controlled to get a suitable microstructure with good mechanical properties during the manufacturing process. Since there are several kinds of heating and cooling treatments in the steel product line, the nucleation sites and growth behavior in the forward and reverse phase transformation are important controlling factors, which are strongly affected by a local inhomogeneity such as precipitation, segregation, and crystal orientations.

Using a field emission typed SEM equipped with the Electron Backscatter Diffraction (EBSD) detector, the resolution for the local crystal orientation area is improved up to about 20nm at room temperature. As a next step, the development of the EBSD measurement at elevated temperature is expected in order to carry out the in situ observation for the phase transformation, recrystallization and grain growth [1]. Since the observation temperature at a high temperature is depended by the capacity of the heating stage, we have applied a new heating stage developed by TSL solutions specified for the EBSD measurement. Figure 1 shows photographs of a part of SEM (JSM-7800), in which the heating stage has been installed. The sample size is 5.0 x 7.0 x 0.6 mm in dimensions. The thermocouples are mounted on both of a sample surface and the sample holder. The heating stage can be operated at tilting angles up to 70°and a radiation protected cover is fixed.

The reverse phase transformation from the ferrite to austenite phase in the bainitic steel has been investigated to make clear the austenite grain orientation memory effect [2]. Since the memory effect shows strong heating rate dependence and the heating temperature range, it is considered that the growth rate of retained austenite between the lath boundaries compete with the decomposition of cementite in the microstructure of the bainite phase. Figure 2 is one of examples showing the orientation imaging maps during the reverse transformation measured up to 1100℃. The in situ observation technique developed at elevated temperature with a stable drift control, it becomes possible to discuss the nucleation sites of the austenite and its orientation characteristics. The artifact arising from the vacuum experiment in SEM must be also discussed.

References

[1] S.Wright and M.M.Nowell ; Electron Backscatter Diffraction in Materials Science, A.J.Schwartz et al. (eds) , (2009) 329.

[2] T.Hara, N.Maruyama, Y.Shinohara, H.Asahi, G.Shigesato,M.Sugiyama, T.Koseki, ISIJ Int., 49(2009) 1792.


This work was supported by the New Energy and Industrial Technology Development Organization (NEDO) under the " Innovative Structural Materials Project (Future Pioneering Projects)".

Fig. 1: Photographs showing a SEM column and the heating stage fixed inside the chamber.

Fig. 2: Orientation imaging maps measured at high temperature using the heating stage. (a) Ferrite(BCC) map at 1000C, (b) Austenite(FCC) map at 1000C, (c) Ferrite(BCC) map at 1100C, (d) Austenite(FCC) map at 1100C.

IT-10. Electron tomography

Type of presentation: Invited

IT-10-IN-1695 Colouring atoms in 3 dimensions

Bals S.1, Goris B.1, Altantzis T.1, Van Tendeloo G.1
1EMAT, University of Antwerp, Belgium
sara.bals@ua.ac.be

New developments in the field of nanoscience drive the need for 3 dimensional (3D) quantitative characterization techniques yielding information down to the atomic scale. The 3D resolution of electron tomography was recently pushed to the atomic level [1-3]. One approach is based on compressive sensing (CS), a technique specialized in finding a solution with a sparse representation to a set of linear equations. At the atomic scale, the approach exploits the sparsity of the object since only a limited number of voxels is occupied by atoms. The CS technique was applied to the 3D reconstruction of Au nanorods based on only 4 HAADF-STEM images. The crystal lattice of the nanorods was reproduced without prior knowledge on the atomic structure [3]. As shown in Figure 1, also the 3D visualization of crystal defects at the atomic scale is currently possible using the same technique.

Going further than determining the 3D positions of atoms, a crucial aim is identifying the type of individual atoms in hetero-nanoparticles. We recently investigated core-shell Au@Ag nanorods using the CS methodology [4]. A detailed analysis of the position and the atom type was performed using orthogonal slices through the 3D reconstruction (Figure 2). Individual Ag and Au atoms can be distinguished, even at the metal-metal interface, by comparing their relative intensities. These results demonstrate the feasibility of chemically sensitive 3D reconstructions with a resolution at the atomic scale. However, such experiments are experimentally and computationally still far from straightforward and very time consuming.

An alternative approach to resolve the chemical composition of complex nanostructures in 3D is by using energy dispersive X-ray (EDX) mapping. This is a suitable technique for electron tomography since the number of generated X-rays increases with sample thickness. Early 3D EDX experiments were complicated by the specimen-detector geometry [5], but recent efforts enable 3D EDX in an optimized manner [6]. A 3D EDX reconstruction of a Au@Ag nanocube is presented in Figure 3 and clearly illustrates the potential of 3D EDX mapping, but further challenges include extracting quantitative information from such reconstructions.

[1] S. Van Aert, K. J. Batenburg, M. D. Rossell, R. Erni, G. Van Tendeloo, Nature 470 (2011) 374

[2] M.C. Scott et al., Nature 483 (2012) 444

[3] B. Goris, S. Bals, W. Van den Broek, E. Carbo-Argibay, S. Gomez-Grana, L. M. Liz-Marzan, G. Van Tendeloo, Nature Materials 11 (2012) 930

[4] B. Goris, A. De Backer, S. Van Aert, S. Gómez-Graña, L. M. Liz-Marzán, G. Van Tendeloo, S. Bals, Nano Lett. 13 (2013) 4236

[5] G Möbus, RC Doole and BJ Inkson, Ultramicroscopy 96 (2003) 433

[6] P Schlossmacher et al, Microscopy Today 18 (2010) 14


We acknowledge support from the ERC ( “Countatoms-#24691”, and “Colouratoms-#335078”) and the FWO. We thank Prof. Liz-Marzán for providing the samples and useful discussions.

Fig. 1: Fig. 1 (a) 3D reconstruction of a Au nanodumbbell. At the tip, twin boundaries are present. (b) Orthogonal slices through the reconstruction of the tip of the nanodumbbell, presented in more detail in (c) enable one to determine the stacking across the twin boundary.

Fig. 2: Fig. 2 (a) Orthogonal slices through the reconstruction show the core-shell structure of the nanorod (b) Detailed view of a slice through the reconstruction perpendicular to the major axis of the nanorod (c) Intensity profile acquired along the direction indicated in (b), showing the capability to assign each atom to be either Ag or Au.

Fig. 3: Fig. 3 (a) 2D EDX map of a Au@Ag nanocube. A tilt series of these 2D EDX maps was acquired and a 3D reconstruction shown in (b) could be obtained.

Type of presentation: Invited

IT-10-IN-3013 Functional soft matter

Friedrich H.1
1Department of Chemical Engineering and Chemistry, Eindhoven University of Technology, Den Dolech 2, 5612AZ Eindhoven, The Netherlands
h.friedrich@tue.nl

The design and synthesis of materials with novel functional properties is a major focus of chemical research. This includes soft matter which are formed by the interactions that arise from (self)-organizing molecules, polymers, and clusters over length scales beyond typical small molecule dimensions. To understand and apply the processes that underlie the formation of nanoparticles and their self-organization into larger functional structures requires 3D nanoscale imaging [1,2]. We focus on the (liquid phase) (self)-organization of soft (in)organic materials and composites thereof using cryogenic (scanning) TEM electron tomography (ET). In this presentation I will take you through the complex and beautiful 3D nano and meso landscape of functional soft matter. Examples will include quantitative ET of a Ruthenium loaded carbon nanotube based heterogeneous catalyst (Figure 1a) [3], quantitative ET of the assembly process of organic solar cell bulk heterojunctions composed of P3HT and PCBM polymers (Figure 1b) [4], and cryogenic ET of liquid infiltration and drying processes in ordered mesoprorous silica (SBA-15) crystallites [5]. Since to-date, more frequently a detailed quantification understanding of particle sizes, size distributions, or particle location and distances is required, I will focus on this information contributes to determine the self-organization pathways [6,7]. Furthermore, I will discuss the effects of limited electron dose, applied angular sampling scheme, and reconstruction algorithm on the achievable 3D resolution (Figure 1d-h) [8,9]. Our findings suggest that for cryo conditions fewer images in the tilt-series are advantageous, contradictory to Crowther’s sampling-based resolution estimate [8,9]. Finally, I will conclude with an outlook on trends for 3D imaging.

[1] H. Friedrich et al, Angewandte Chemie International Edition 49 (2010) 7850.

[2] H. Friedrich et al, Chemical Reviews 109 (2009) 1613.

[3] H. Friedrich et al, ChemSusChem 4 (2011) 957.

[4] M. Wirix et al. Nanoletters (2014) accepted.

[5] T. M. Eggenhuisen et al, Chemistry of Materials 25 (2013) 890.

[6] G. Prieto et al, Nature Materials 12 (2013) 34.

[7] J. Zečević et al. ACS Nano 7 (2013) 3698.

[8] D. Chen et al. Journal of Physical Chemistry C 118 (2014) 1248.

[9] D. Chen et al. manuscript in preparation.


The author gratefully acknowledge the contributions of all (co)authors of the referenced manuscripts and especially, J. Zečević, D. Chen, M. Wirix, G. Prieto, T. M. Eggenhuisen, and funding from the NRSCC, NWO and the European Union.

Fig. 1: Quantitative 3D imaging examples: (a) Ru/CNT catalyst; (b) P3HT/PCBM bulk heterojunction; (c) infiltrated and dried SBA-15 crystallite; (d) simulation model to determine ET resolution and reconstructions using (e) SIRT; (f) TVM; , DART (d), and WBP (e) at a total electron does of 104 e/Å2 and tilt increments of 1°.

Type of presentation: Oral

IT-10-O-1637 Combined tilt- and focal series scanning transmission electron microscopy: TFS 3D STEM

Dahmen T.1, Baudoin J. P.2, Lupini A. R.3, Kuebel C.4, Slusallek P.1, de Jonge N.5
1German Research Center for Artificial Intelligence GmbH, Saarbrücken, Germany, 2Department of Molecular Physiology and Biophysics, Vanderbilt University School of Medicine, Nashville, TN, USA, 3Materials Science and Technology Division, Oak Ridge National Laboratory, Oak Ridge, TN, USA, 4KIT –Karlsruhe Institute for Technology, Eggenstein-Leopoldshafen, Germany, 5INM – Leibniz Institute for New Materials, Saarbrücken, Germany
niels.dejonge@inm-gmbh.de

Tilt-series transmission electron microscopy (TEM) tomography is the method of choice to obtain nanoscale three-dimensional (3D) information from samples in biology and materials science. 3D data is acquired by mechanically tilting the sample stage and recording images at each tilt angle. However, the tilt range is limited to ±70° for most samples, and the tomographic reconstruction suffers from missing information and a limited resolution on account of the so-called “missing wedge”. Furthermore, the quality of the reconstruction critically depends on the precision of the alignment of the individual images. Alternatively, scanning TEM (STEM) focal series can be recorded avoiding tilting altogether but this method lacks axial resolution [1]. A new recording scheme to obtain 3D data is presented by combined tilt- and focal series (TFS) STEM. This method significantly reduces the two aforementioned limitations of tilt series tomography. The specimen is rotated in relatively large tilt increments, and for every tilt direction, a through-focal series is recorded (Fig. 1). Both the tilt-series and focal-series data are reconstructed into a 3D tomogram in the same software algorithm. The conical shape of the STEM probe is taken into account via forward- and backward projection operators. The TFS method exhibits reduced “missing wedge” artifacts and a higher axial resolution than obtainable using STEM tilt series [2]. Fig. 2 shows that the missing wedge is still present in the TFS but low spatial frequency signal components are now present (arrow) in the central vertical region. Streaks corresponding to the tilt directions are also less pronounced. The TFS reconstruction results in a superior shape representation and tolerates a much smaller number of tilt angles than tomography, which is beneficial for image stack alignment. A further advantageous application of TFS STEM is the imaging of micrometers-thick samples. With TFS it is possible to limit the overall tilt range while obtaining a higher axial resolution than for a tilt series alone.


We thank L. Marsallek and S.J. Pennycook for discussions, and E. Arzt for support through the INM. Electron microscopy performed at the SHaRE user facility at Oak Ridge National Laboratory, and at the Karlsruhe Nano Micro Facility, Helmholtz infrastructure, Karlsruhe Institute of Technology. Research supported by the U.S. Department of Energy, Basic Energy Sciences, and by NIH grant R01-GM081801.

Fig. 1: Schematic views of the TFS recording scheme with STEM. (left) A thin specimen is imaged pixel-by-pixel in dark field mode STEM using the annular dark field detector (ADF). (right) In a combined tilt- and focal series, images are acquired in a through-focal series at each tilt angle. The specimen stage is titled after each focus series.

Fig. 2: Comparison of tilt series STEM tomography with TFS STEM. (left) Spatial frequency spectrum (Fourier Transform) of a vertical (xz) slice of the conventional tilt series STEM tomography data. (right) Frequency spectrum of an xz slice of the TFS STEM dataset. The red lines mark the border of the missing wedge. With permission from [2].

Type of presentation: Oral

IT-10-O-1863 Towards mass contrast in 3D atomic resolution tomographic reconstructions from HRTEM images with the inverse dynamical electron scattering method

Van den Broek W.1, Koch C. T.1
1Institute for Experimental Physics, Ulm University, Albert-Einstein-Allee 11, 89081 Ulm, Germany
wouter.vandenbroek@uni-ulm.de

We have recently reported a novel reconstruction algorithm, inverse dynamical electron scattering (IDES), that retrieves the object’s electrostatic potential from a high resolution transmission electron microscopy (HRTEM) tilt series [1]. By reformulating the multislice algorithm as an artificial neural network IDES is able to very efficiently recover the scattering object using a gradient based optimization. Prior knowledge about the atomic potential shape and the sparseness of the object is accounted for naturally. In [2] we demonstrated the generality of this method with applications to simulated ptychography and scanning confocal electron microscopy data sets and the optimization of the images’ defocus values along with the object.

Reconstructions so far had been performed on simulations of a small Au nanoparticle that was 1.6nm wide. In this work a larger and more complicated system is treated: A PbSe-CdSe core-shell particle with 1963 atoms. The particle is derived from a CdSe rock-salt structure and has a cubic shape with sides of 3.7nm. The unit-cell parameter a equals 0.61nm. The Cd atoms in the cube’s interior octahedron are replaced by Pb-atoms and shifted by a vector [−a/4, a/4, a/4]; see [3].

The potential is reconstructed from a simulated HRTEM tilt series with the α-tilt varying from -10° to +10° in 2°-steps with zero β-tilt, and the β-tilt varying from -10° to +10° in 2°-steps with zero α-tilt. The images were given a signal-to-noise ratio of 10. See Fig. 1 for 6 typical simulations out of the total of 21.

The forward simulations have been carried out with the new FDES (forward dynamical electron diffraction) software that runs on the graphics processing unit. A solution was achieved with conjugate gradients optimization. The solution’s sparseness was increased through L1-norm regularization and through the use of a generalized potential, taken as the weighted average of the potentials of Se, Cd and Pb.

Although FDES was set to use more accurate approximations than IDES—specimen rotation is treated exactly in FDES, but is approximated by a shifted Fresnel propagator in IDES and FDES used a multislice slice thickness 5 times smaller than IDES—an atomic resolution reconstruction was still possible. The Pb-atoms can be isolated by thresholding the reconstructed gray values, thus demonstrating mass contrast. Fig. 2 shows a side-by-side comparison of the original and the reconstructed particle potential, visualized with IMOD [4].

[1] W. Van den Broek, C.T. Koch. Phys. Rev. Lett. 109 (2012) 245502.
[2] W. Van den Broek, C.T. Koch. Phys. Rev. B 87 (2013) 184108.
[3] S. Bals et al. Nano Lett. 11 (2011) 3420–3424.
[4] http://bio3d.colorado.edu/imod/index.html


The authors acknowledge the Carl Zeiss Foundation as well as the German Research Foundation (DFG, Grant No. KO 2911/7-1).

Fig. 1: Six typical simulations. The acceleration voltage is 80kV, the defocus is -20nm, the spherical aberration is 64μm, the pixel size is 0.025nm and the multislice slice thickness is 0.02nm. The point resolution is 0.17nm. The images are degraded by partial temporal and spatial coherence and the CCD’s modulation transfer function.

Fig. 2: Left: Pb-atoms in the original particle’s core. Right: The Pb-atoms in the reconstructed potential, identified by simple thresholding, thus demonstrating mass contrast. Due to the high level of noise in the measurements there are some false positives and false negatives. Lowering the threshold would also reveal the Cd and Se atoms.

Type of presentation: Oral

IT-10-O-2126 Four-dimensional simultaneous EELS & EDX tomography of an Al-Si based alloy

Haberfehlner G.1, Albu M.1, Orthacker A.1, Kothleitner G.1
1Institute for Electron Microscopy and Nanoanalysis, Graz University of Technology & Centre for Electron Microscopy and Nanoanalysis, Graz, Austria
georg.haberfehlner@felmi-zfe.at

Electron tomography is a powerful technique for 3D characterization at the nanoscale. Recent developments in electron tomography include the use of TEM imaging techniques based on inelastic scattering, such as EFTEM, EELS and EDX [1,2] as well as reconstruction algorithms, including prior information in the reconstruction process, such as total-variation (TV) minimization [3].

Modification of eutectic Si in Al-Si alloys by modifying elements such as Sr, Na or Yb is a frequently used method to improve mechanical properties of alloys [4]. Nevertheless the modification mechanisms are still a matter of debate and require understanding the elemental distribution down to the atomic level.

In this work we investigate an Al-5 wt.% Si alloy with 15 ppm Na and 6500 ppm Yb. On this alloy we demonstrate simultaneous EELS and EDX tomography. We apply a TV minimization reconstruction algorithm to elemental maps extracted from analytical TEM data and we reconstruct spectral EELS and EDX data, to get local spectra in three dimensions.

Tomography experiments were performed on a probe-corrected FEI Titan3 microscope operated at 300 kV, equipped with a Gatan GIF Quantum energy filter and a Bruker Super-X detector. Low-loss (-80 to 920 eV), core-loss EELS (1120 to 2120 eV) and EDX spectrum images of a FIB-prepared needle-shaped sample were acquired every 5° over a range of +/-75°. All analytical tomography data was aligned based on HAADF STEM images, acquired at the same time as the spectrum images. 2D elemental maps were extracted for each tilt angle from both EDX and EELS spectrum image data sets.

Reconstruction of the elemental maps was done with the simultaneous iterative reconstruction technique (SIRT) as well as using TV minimization. TV minimization is an efficient method for reconstruction from few and noisy projections, as is the case for elemental maps. It assumes locally constant regions in the reconstruction, a valid assumption as we expect sharp interfaces between Yb-rich precipitates, Si-particles and the Al-rich matrix. Fig. 1 shows reconstructions of elemental maps extracted from EDX data and Fig. 2 shows the same reconstructions of elemental maps for EELS core-loss data.

Additionally we reconstructed spectral EDX and EELS data. Using SIRT each spectral channel is reconstructed, which provides four-dimensional datasets containing EELS and EDX spectra for each voxel (see Fig. 3).

This work founds a basis for quantitative elemental mapping in three dimensions and can be extended towards 3D chemical fingerprinting or extraction of local electrical and optical properties.

[1] Jarausch et al, Ultramicroscopy 109:326, 2009

[2] Lepinay et al, Micron 47:43, 2013

[3] Leary et al, Ultramicroscopy 131:70, 2013

[4] Li et al, Philos. Mag. 92:3789, 2012


The authors would like to thank Jiehua Li and Peter Schumacher from Chair of Casting Research, University of Leoben for providing the samples, This work has been supported by the FFG OPTIMATSTRUCT project and within the European Union's 7th Framework Programme in the project ESTEEM2.

Fig. 1: Reconstructions of elemental maps extracted from EDX data. (a) Orthogonal slices through reconstructions of the Yb L-line, Si K-line and Al K-line using SIRT and TV minimization. (b) 3D surfaces extracted from the TV minimization reconstructions.

Fig. 2: Reconstructions of elemental maps extracted from EELS core-loss data. (a) Orthogonal slices through reconstructions of the Yb M45-edge, Si K-edge and Al K-edge using SIRT and TV minimization. (b) 3D surfaces extracted from the TV minimization reconstructions.

Fig. 3: Local reconstructed spectra from (a) Yb-rich precipitates, (b) Si-rich region, (c) Al-rich matrix. (1) Masks of the investigated regions, (2) core-loss EELS spectra, (3) EDX spectra. (2)-(3) show single voxel spectra (top) and spectra summed over all voxels of the masks shown in (1) (bottom).

Type of presentation: Oral

IT-10-O-2253 Possibilities and limitations for atom counting using quantitative ADF STEM

De Backer A.1, Martinez G. T.1, MacArthur K. E.2, Jones L.2, Béché A.1, Nellist P. D.2, Van Aert S.1
1Electron Microscopy for Materials Science (EMAT), University of Antwerp, Antwerp, Belgium, 2Department of Materials, University of Oxford, Oxford, United Kingdom
Annick.DeBacker@uantwerpen.be

Advanced statistical methods can be used to count the number of atoms in each atom column of high-resolution ADF STEM images [1-3]. Here we discuss the possibilities and limitations of achieving single atom sensitivity.
Four images of the same Ir/Pt nanoparticle were recorded at different magnifications and electron doses. In order to allow comparison with simulations, the images were normalised with respect to the incoming electron beam intensity [4]. Next, using statistical parameter estimation theory, the total scattered intensities are quantified atom column–by–atom column. An example analysis for the image recorded at the highest magnification and electron dose is illustrated in Fig. 1; the total scattered intensities are visualised in the histogram. The number of significant components and their intensities were retrieved by evaluating the so-called integrated classification likelihood (ICL) criterion in combination with Gaussian mixture model estimation. These results allow us to quantify the number of atoms in each atom column. As shown in [3], the reliability of atom counts depends on the number of atom columns present in an image, the width of the components, and the performance of the ICL criterion. These parameters can be linked with the quality of the recorded images.
In Fig. 2, the intensities of the components resulting from the counting analyses are compared with the total scattered intensities resulting from simulated images using STEMsim. For image 3 an excellent match was found. However, analysing images of lower magnification and/or electron dose worsens the match with simulation. The same effect is observed when analysing an image composed of every second pixel of image 3. In this way, the lower magnification of images 1 and 2 is mimicked. This leads to less precise measurements of the total scattered intensities resulting in insufficient statistics for the determination of the number of components. However, when enhancing the statistics by combining the values of the scattered intensities of the four images collectively, the experimental intensities again match with simulated values. In addition, the statistical approach for atom counting provides us high precision leading to near single atom sensitivity for this combined set of images.
In conclusion, an advanced quantitative method to count the number of atoms is presented together with its possibilities and limitations. Single atom sensitivity may be achieved when the experimental images are of sufficient quality to yield sufficient statistics.

References

[1] S Van Aert et al., Nature 470, p 374 (2011)
[2] S Van Aert et al., PRB 87, 064107 (2013)
[3] A De Backer et al., Ultramicroscopy 134, p 23 (2013)
[4] A Rosenauer et al., Ultramicroscopy 109, p 1171 (2009)


Funding from the FWO Flanders, the EU FP7 (312483 - ESTEEM2), and the UK Engineering and Physical Sciences Research Council (EP/K032518/1) is acknowledged.

Fig. 1: Illustration of the atom counting procedure; a) experimental ADF STEM image, b) refined parameterised imaging model, c) evaluation of ICL as a function of the number of components, d) histogram of estimated scattered intensities together with the estimated Gaussian mixture model, e) quantification of the number of atoms in each atom column.

Fig. 2: Comparison of experimental and simulated total scattered intensities. The inset shows the specific pixel size and dwell time for the individual images of the Ir/Pt particle.

Type of presentation: Oral

IT-10-O-2435 Towards 4-D EEL spectroscopic scanning confocal electron microscopy (SCEM-EELS) optical sectioning on a Cc and Cs double-corrected transmission electron microscope

Wang P.1, Boothroyd C. B.2, Dunin-Borkowski R. E.2, Kirkland A. I.3, Nellist P. D.3
1National Laboratory of Solid State Microstructures and College of Engineering and Applied Sciences, Nanjing University, Nanjing 210093, People’s Republic of China, 2ER-C and PGI5, Forschungszentrum Jülich, D-52425 Jülich, Germany, 3Department of Materials, University of Oxford, Parks Road, Oxford OX1 3PH, UK
wangpeng@nju.edu.cn

The spectrum imaging method of combining scanning transmission electron microscopy with electron energy-loss spectroscopy (STEM-EELS) has been widely used for materials characterization at the atomic-scale. Three dimensional (3-D) optical sectioning using scanning confocal electron microscopy (SCEM) [1], as shown in Fig. 1 a), has been developed as an alternative to tilt-series electron tomography [2]. The confocal imaging mode in STEM has been implemented with spherical aberration (Cs) correctors, which allow the use of substantially increased objective apertures and hence provide a considerably decreased depth of focus, typically below 10 nm. Both the theoretical basis of the image contrast [3] and the experimental implementation of the technique [4-6] have been studied. However, due to Cc-aberrations in the post-specimen optics, inelastically scattered electrons with different energy losses △E are focused at different focal length (Fig. 1a)), which causes an EEL spectrum to be out-of-focus away from the confocal energy loss [5], as shown in Fig. 2a). In order to avoid this problem, a Cc-corrector is need in addition to a double-aberration corrected TEM, as shown in Fig. 1b).
In this work, we propose a novel spectrum imaging mode by combining the SCEM and EELS techniques, which can potentially let one perform 4D EEL spectroscopic SCEM (or so called SCEM-EELS for short) optical sectioning, allowing quantitative chemical characterization over a full 3D specimen volume. Preliminary experiments have been carried out both on a non Cc-corrected Oxford-JEOL 2200MCO instrument with 3rd order double Cs correctors and on a Cc-corrected FEI Titan 60-300 PICO, which has an illumination-side Cs corrector, a Cs-Cc achro-aplanat image corrector and a post-specimen EEL spectrometer. Fig. 2b) shows 2D spectrum images recorded from an amorphous carbon film on the EELS CCD camera in a confocal configuration aligned for energy losses of 0 eV on a Cc-corrected TEM. It demonstrates that the inelastically scattered electrons are simultaneously in-focus on the EELS CCD camera over the entire energy loss range.

References:
[1] P.D. Nellist, P. Wang, Annual Review of Materials Research, 42, (2012), 125-143.
[2] P.A. Midgley, M. Weyland, Ultramicroscopy, 96 (2003) 413-431.
[3] A.J. D'Alfonso et al, Ultramicroscopy, 108 (2008) 1567-1578.
[4] P. Wang et al, Ultramicroscopy, 111 (2011) 877-886.
[5] P. Wang et al, Physical Review Letters, 104 (2010) 200801.
[6] P. Wang et al , Applied Physics Letters, 100 (2012) 213117.


P.W., A.I.K. and P.D.N. acknowledge financial support from the Leverhulme Trust (F/08 749/B), the EPSRC (EP/F048009/1); P.W. acknowledges financial support from the Thousand Talents Program.

Fig. 1: Schematic diagrams of confocal trajectories for SCEM with an EEL spectrometer behind a circular pinhole on a non Cc-corrected TEM (a) and a Cc-corrected TEM (b), respectively. Due to Cc aberration correction in the post-specimen lenses in (b), electron rays with an energy loss difference of △E can still be focused on the pinhole plane.

Fig. 2: 2D spectrum images recorded on the EELS CCD camera in a confocal configuration established for energy losses of 0 eV on a non Cc-corrected TEM from a Si slab (a) and on a Cc-corrected TEM from an amorphous carbon film (b), respectively.

Type of presentation: Oral

IT-10-O-2543 On-axis electron tomography of needle-shaped biological samples

Saghi Z.1, Divitini G.1, Winter B.2, Spiecker E.2, Ducati C.1, Midgley P. A.1
1(1) Department of Materials Science and Metallurgy, University of Cambridge, 27 Charles Babbage Road, Cambridge, CB3 0FS, UK, 2(2) Center for Nanoanalysis and Electron Microscopy (CENEM), Department Werkstoffwissenschaften / VII, Universität Erlangen-Nürnberg, Cauerstraße 6, 91058 Erlangen, Germany
saghizineb@gmail.com

A key challenge in structural biology is to image large volumes while maintaining sufficient resolution to identify small features in their original cellular context [1]. Electron tomography (ET) has contributed greatly to this field, but imaging sections thicker than a few hundred nanometers is difficult because of the sample geometry and microscope configuration: as the specimen is tilted to high angles, the thickness increases and the quality of the images deteriorates. Moreover, for geometric constraints, the tilt range rarely exceeds ±70°, leading to elongation and blurring of features, and an overall challenging volume to segment. Here, we show that preparing a needle-shaped sample rather than a flat section can alleviate many of the limitations encountered in biological ET. The technique is illustrated on a 500nm diameter needle extracted from an epoxy-embedded portion of the nucleus accumbens shell from a rat brain. The sample was prepared in a Helios NanoLab focused ion beam (FIB) machine and transferred to an on-axis tomography holder. An HAADF-STEM tilt series was then acquired from -90° to +90° with 1° increment, using an FEI TITAN microscope operating at 300kV, and Inspect3D was used for the alignment and reconstruction by weighted backprojection. Figure 1(a) shows a low magnification view of the needle and the region selected for tomography. A voxel projection and slice through the reconstructed needle (b,c) provide highly detailed structural information. In Figure 2, we compare the quality of ET results from -90°:2°:+90° and -60°:2°:60° acquisitions. Cross-sections through a portion of exitatory synapse and mitochondrion (positions 1 and 2 in Figure 1(c)) illustrate the advantages of a full tilt range on-axis ET experiment with enhanced signal-to-background ratio and isotropic sharpness of features. Note that the ±60° cylindrical volume shown here is still of better quality than a reconstructed section from similar tilt range, since the thickness remains constant throughout the tilt series.
Combining this novel sample preparation technique with advanced imaging modes (BF-STEM for example [2]) and sophisticated reconstruction algorithms such as compressed sensing [3], we anticipate that ET will provide a complementary method to serial sectioning and FIB-SEM slice-and-view techniques [4].

[1] W. Baumeister, Current Opinion in Structural Biology 2002, 12(5):679.
[2] A.A. Sousa et al., Journal of Structural Biology 2011, 174(1): 107.
[3] Z. Saghi et al., Nano Letters 2011, 11(11):4666.
[4] Samples provided by Andrea Falqui and Roberto Marotta, IIT, Genova, Italy.


The research leading to these results has received funding from the European Union Seventh Framework Programme under Grant Agreement 312483-ESTEEM2 (Integrated Infrastructure Initiative–I3), as well as from the European Research Council under the European Union’s Seventh Framework Programme (FP/2007-2013)/ERC grant agreement 291522-3DIMAGE.

Fig. 1: (a) Needle-shaped biological sample, prepared by FIB. The rectangle indicates the area selected for electron tomography. (b) Voxel projection of the reconstruction from a full range acquistion. (c) A slice through the volume showing a detailed view of an exitatory synapse (1) and a mitochondrion (2).

Fig. 2: Cross-sectional slices through the synapse (a,b) and mitochondrion (c,d) with different tilt ranges. Isotropic resolution is observed for full tilt range on-axis tomography (left), as illustrated in the insets.

Type of presentation: Oral

IT-10-O-2765 A method for quantitative analysis and improvement of 3D electrostatic nanopotentials reconstructed by electron holographic tomography

Wolf D.1, Lubk A.1, Lichte H.1
1Triebenberg Laboratory, Institute of Structure Physics, Technische Universität Dresden, Dresden, Germany
daniel.wolf@triebenberg.de

Electron holographic tomography (EHT), i.e. using off-axis electron holography (EH) as imaging mode for electron tomography (ET) in the transmission electron microscope (TEM), facilitates the 3D mapping of materials on the nanometer scale [1,2]. The phase shift of the electron wave that can be reconstructed by EH contains the projected electrostatic scalar potential and, for magnetic samples, the projected magnetic vector potential of the specimen [3]. Therefore, tomographic reconstruction of phase tilt series results in 3D maps of the electric potential (magnetic case is not considered here).

At nanometer resolution (1-10nm), the major contribution to tomograms reconstructed by EHT is the mean inner potential (MIP). Its value depends on the atomic species, the atomic packing in the unit cell, but also on the distribution of the valence electrons. Thus, the MIP represents a finger print of chemical composition and can be used to detect for example core-shell structures (e.g. AlGaAs-GaAs [2]) or gradients of composition in nanowires (NWs). Recently, the three-dimensional nanosponge structure of Si embedded in SiO2 has been revealed with EHT [4]. Furthermore, functional potentials, such as the built-in potentials across p-n junctions in semiconductors can be measured [1,2]. In this context, also surface and sub-surface effects, e.g. Fermi-level pinning [5], have been studied, quantitatively.

In order to extract quantitative information from the 3D reconstructions, it is indispensable to know their fidelity. Here, we show a procedure to proof the reliability of the tomograms by comparing their re-projections with the original ones (Fig. 1a)). By applying this procedure on an Ag, ZnO and Si NW and evaluating the potential averaged over the entire specimen, we determine the MIP values from the projection data (Fig. 1b)).

Moreover, the 3D reconstruction can be remarkably improved by normalizing it with the tomogram reconstructed from the projected thickness. The latter is obtained after step 3 in the procedure shown in Fig. 1a). Because its reconstruction is done from the same tilt range, the resulting tomogram contains very similar missing wedge artifacts as the original one. Therefore, such artifacts can be corrected to a great amount using this approach (compare in Fig. 2: a,b with c,d).

[1] P.A. Midgley and R.E. Dunin-Borkowski, Nature Materials 8, (2009), p. 271.
[2] D. Wolf, A. Lubk, F. Röder, and H. Lichte, Current Opinion in Solid State and Materials Science 17, (2013), p. 126.
[3] H. Lichte and M. Lehmann, Reports on Progress in Physics 71 (2008), p. 016102.
[4] R. Hübner, D. Wolf, D. Friedrich, B. Liedke, B. Schmidt, K.H. Heinig, at this conference.
[5] D. Wolf, A. Lubk, A. Lenk, S. Sturm, and H. Lichte, Appl. Phys. Lett. 103 (2013), p. 264104.


We thank M. Graf, TU Dresden for providing the Ag nanowire, and Z. L. Wang, Georgia Institute, Atlanta for providing the ZnO nanowire. The research leading to these results has received funding from the European Union Seventh Framework Programme under Grant Agreement 312483 - ESTEEM2 (Integrated Infrastructure Initiative - I3).

Fig. 1: a) Procedure to determine from the projected potential tilt series the averaged potential tilt series on the example of an Ag nanowire. b) Potential averaged over the entire specimen vs. projection angle a. The mean of these values corresponds to the mean inner potential V0.

Fig. 2: 3D reconstruction with reduced missing wedge artifacts on the example of an Ag NW and a ZnO NW. a,b) Slice through 3D potential. c,d) Same slice as in (a,b) but normalized with the reconstruction of the projected thickness. e,f) Line profiles corresponding to the gray arrows in a-d).

Type of presentation: Oral

IT-10-O-2787 EELS and tomography: from EELS Spectrum Images to Spectrum Volumes.

Yedra L.1,2, Eljarrat A.1, Rebled J. M.1,3, López-Conesa L.1, Dix N.3, Sanchez F.3, Estradé S.1,2, Peiró F.1
1Laboratory of Electron Nanoscopies (LENS)- MIND/IN2UB, Dept. d'Electrònica, Universitat de Barcelona, Barcelona, Spain, 2CCiT, Scientific and Technical Centers, Universitat de Barcelona, Barcelona, Spain, 3Institut de ciencia dels materials de Barcelona (ICMAB), Bellaterra, Spain
llyedra@el.ub.edu

In transmission electron microscopy (TEM), 3D tomographic reconstruction can be achieved by acquiring a series of images at different tilt angles. A different approach is obtaining 3D chemical reconstructions from energy filtered images in the TEM (EFTEM)[1-3], and more recently, by acquiring EELS spectrum images (EELS-SI), each pixel containing a complete EELS spectrum [4,5]. However, in both techniques only a limited amount of information is effectively reconstructed. In this paper we aim to derive a full EELS dataset in 4D, where every voxel of a whole volume contains a complete spectrum of energy losses, as schematized in Fig. 1. By analogy to the spectrum image notation, we will name this 4D dataset as EELS spectrum volume (EELS-SV).

Our approach to EELS-SV reconstruction is based upon SI, thus taking a single SI for every tilt angle. It takes advantage of Multivariate Analysis (MVA), and more precisely of blind source separation (BSS)[6], to find a new spectral basis (Fig. 2a) which can describe all the spectra in the dataset as a weighted sum of its components. Therefore only the 3D reconstructions of the weighting components (Fig. 2b) will be necessary to recover the spectra in each voxel (Fig. 2c-e). We will apply this approach to analyze a BFO/CFO nanocomposites, enabling the characterization of a CFO nanocolumn embedded in BFO matrix.

References

[1] G. Möbus, et al, Ultramic, 96 (2003) 433.

[2] M. Weyland, P.A. Midgley, Microscopy and Microanalysis, 9 (2003) 542.

[3] R.D. Leapman, et al, Ultramic, 100 (2004) 115.

[4] K. Jarausch, et al, Ultramic,. 109 (2009) 326.

[5] L. Yedra et al., Ultramic., 122 (2012) 12

[6] N. Dobigeon et al., Ieee Transactions on Signal Processing, 57 (2009), 4355


Fig. 1: Schematic of the 4D dataset, the EELS spectrum volume, consisting of 3 spatial dimensions plus an additional energy loss dimension. Here it is presented along with an extracted xy spectrum image, a spectrum line along z direction and a single spectrum from an inner voxel.

Fig. 2: EELS-SV reconstruction procedure. a) Components of the spectrum. b) 3D reconstructions, c) Schematic representation of two orthoslices and reconstructed SI for transversal and longitudinal orthoslices. d) Single spectrum and e) spectrum line extracted from the slices.

Type of presentation: Oral

IT-10-O-2806 Electron cryo-tomography with a new type of phase plate

Danev R.1, Buijsse B.2, Fukuda Y.1, Khoshouei M.1, Plitzko J.1, Baumeister W.1
1Max Planck Institute of Biochemistry, Martinsried, Germany, 2FEI, Eindhoven, The Netherlands
danev@biochem.mpg.de

Recent years have shown an increased interest in the development and use of phase plates in cryo-EM. The oldest and the most productive type of phase plate is the carbon film Zernike phase plate. Despite its good performance the Zernike phase plate has a few pitfalls. One major practical hindrance is its short lifetime, typically about 10 days. Another disadvantage is the generation of fringes around high-contrast features in the image. Despite its shortcomings the Zernike phase plate has been the main motivation and experience generator in the last years.

We are currently working in collaboration with FEI on the development and testing of a new type of phase plate. It addresses both of the above mentioned shortcomings of the Zernike phase plate. Our tests indicate that the new phase plate lasts for more than six months inside the microscope. This is a big advantage in terms of servicing and up time of the microscope. Another big advantage of the new phase plate is that it produces fringe-free images which resemble in appearance light microscopy phase contrast images. The new phase plate is being developed as a part of a product package which will include new hardware – phase plate holder & phase plate, and new software for alignment, calibration and ease of use.

We tested the new phase plate with two automated acquisition software packages – FEI Tomography and SerialEM. Both packages work well and are able to automatically acquire tilt series with the phase plate. There are only a couple of additional steps that are required for setting up the phase plate before starting the tilt series acquisition. The reconstruction process for the phase plate tomograms is almost identical to that for conventional defocus phase contrast tomograms and can be performed with any existing reconstruction software package, such as IMOD. Because of the strong low frequency components in the phase plate images a simple weighted back-projection reconstruction was in most cases sufficient to produce sharp, high-contrast tomograms. No further processing, such as de-noising, was necessary. In a few thick specimen cases a SIRT reconstruction produced better looking tomograms and again no de-noising or other post-processing was necessary. Overall the new phase plate works very well for cryo-tomography and users with cryo-tomography experience require only a short training on how to use it.

The new phase plate was tested in two microscope models – a 200 kV FEI Tecnai F20 and a 300 kV FEI Titan Krios. The lifetime and image quality performance was equally good with both microscopes.


We thank Matthijn Vos from FEI for providing test samples.

Fig. 1: A slice through a phase plate cryo-tomogram of a vitrified primary culture neuron cell.

Fig. 2:
Type of presentation: Oral

IT-10-O-2812 Fast tomography acquisition for in situ 3D analysis of nanomaterials under variable gas and temperature conditions in Environmental-TEM

EPICIER T.1, 2, ROIBAN L.1, LI S.2, AOUINE M.2, SANTOS AIRES F. C.2, TUEL A.2, FARRUSSENG D.2
1MATEIS, INSA de Lyon, Université Lyon I, 69621 Villeurbanne Cedex, France, 2IRCELYON,Université Lyon I, 2, Av. A. Einstein, 69626 Villeurbanne Cedex, France
lucian.roiban@insa-lyon.fr

In the last two decades, tilted tomography in a transmission electron microscope (TEM) has become a widely used approach in order to quantify the three dimensional (3D) distribution of features in materials and nanomaterials[1, 2]. During the tilt series acquisition, a projection of the area of interest is recorded at each angle over a large angular amplitude, the final resolution along Z axis being directly related to the maximal tilting angle. The tilt series acquisition is usually performed automatically; depending on the employed acquisition method (automatic focusing, and cross-correlation based tracking), the total acquisition time typically ranges between 30 minutes to several hours. Such conditions are totally incompatible with in-situ experiments, where the materials are subject to changes under external mechanical or electrical solicitations as well as variable temperature and gas flow. Following the 3D evolution in such a context can be attempted by a ‘before/after’ strategy, where a first tomography analysis is performed on the object prior to any solicitation, then a second one after the solicitation as performed to track fuel cell nanocatalysts during electrochemical aging [3]. The recent development of commercial Environmental TEM (ETEM) [4] offers a wide range of in situ environmental studies of nanomaterials, such as oxidation / reduction at high temperature: this opens new opportunities to (try to) investigate in situ the 3D structure of nanomaterials. In this context, we are currently optimizing a fast acquisition method for tomography studies, based on video acquisition of tilted series in less than 1-4 minutes. We have applied this approach to the study of metallic Ag nanoparticles (NPs) encaged in silicalite hollow shells (silica-cages) for application in selective catalysis [5]. Single-tilt tomography and ETEM experiments were performed on a Cs-corrected TITAN ETEM, 80-300 kV, recently installed at CLYM in Lyon. Results are illustrated by figures 1 (fast acquisition performed over an angular amplitude of 116° in 3 minutes and 40 seconds) and figure 2 (ETEM experiments up to 700°C and oxygen partial pressure of 2 mbar). References [1] P.A. Midgley, R.E. Dunin-Borkowski, Nature Mat., 8 (2009) 271-280. <span>[2] T. Epicier, chap. 3 ‘Imagerie 3D en mécanique des matériaux’, ed. J.Y. Buffière, E. Maire, Hermès - Lavoisier, Paris, (2014). <span>[3] J. Jinschek, Microscopy and Analysis, Nanotechn. Issue November (2012) 5-10. <span>[4] Y. Yu, H.L. Xin, R. Hovden, D. Wang, E.D. Rus, J.A. Mundy, D.A. Muller, H.D. Abruña, Nano Lett., 12 9 (2012) 4417-4423. <span>[5] S. Li, L. Burel, C. Aquino, A. Tuel, F. Morfin, J.L. Rousset, D. Farrusseng, Chem. Comm. 49 (2013) 8507-8509.

 


Thanks are due to CLYM (www.clym.fr) for guidance of the ETEM project financed by CNRS, Région Rhône-Alpes, ‘GrandLyon’ and French Ministry of Research and Higher Education. The authors thank N. Blanchard and C. Langlois for fruitful discussions and L. Burel for her assistance.

Fig. 1: Fast single-tilt tomography; a-b): video frames extracted at 78° and - 38.5° from a continuous tilting series acquired in bright field mode in less than 4 minutes; c): surface rendering of the silica-cages (green) and size histogram of Ag NPs (red); only 3% are outside of the silica cages. Acquisition conditions: high vacuum, 20°C, 300 kV.

Fig. 2: a): Assembly of silica cages containing Ag NPs at 20°C under high vacuum; b): same area at 700°C under high vacuum: the Ag NPs have grown but are mostly still inside the silica-cages; c): other area at 450°C under 2 10-2 mbar of O2 flow: note that all Ag NPs are out of the cages on the carbon supporting film, contrarily to b).

Type of presentation: Oral

IT-10-O-2946 Advanced 3-D Reconstruction Algorithms for Electron Tomography

Arslan I.1, Sanders T.2, Binev P.2, Roehling J. D.3, Batenburg K. J.4, Gates B. C.3, Katz A.5
1Pacific Northwest National Laboratory, Richland, WA, USA, 2University of South Carolina, Columbia, SC, USA, 3University of California-Davis, Davis, CA, USA, 4University of Antwerp, Antwerp, Belgium, and Centrum Wiskunde & Informatica, Amsterdam, The Netherlands., 5University of California–Berkeley, Berkeley, CA, USA
ilke.arslan@pnnl.gov

Electron tomography in the physical sciences has become a powerful tool for nanomaterial analysis. Recently, electron tomography is finding applications in more beam sensitive materials such as catalysts. For beam sensitive materials, the goal is to acquire the smallest number of images as possible but still maintain an accurate and high resolution 3-D reconstruction. Standard methods of 3D reconstruction, such as weighted back projection (WBP) and simultaneous iterative reconstruction technique (SIRT), are not equipped to handle this lack of information, and create significant blurring.  This gives rise to a search for new methods of reconstruction.  Two of the recent successful algorithms are the discrete algebraic reconstruction technique (DART) and total variation (TV) minimization within compressed sensing (CS).

 

DART uses an algebraic reconstruction method (e.g. ART, SIRT) and pairs it with the prior knowledge that there are only a small number (two or three) of different materials in the sample, each corresponding to a different gray value in the reconstruction.  An initial reconstruction is computed and rounded to the chosen fixed gray values based on some threshold, and iteratively refined using ART.  The method of TV minimization stems from the mathematical theory of compressed sensing and only recently became available due to new algorithms for solving the TV minimization problem.   The method considers the characterization of real images and encourages the reconstruction to minimize the number of jumps in gray values, creating clearer material boundaries than conventional methods (i.e. WBP or SIRT), hence creating a similar effect to that of DART.

 

The advantage of DART is that an accurate selection of the gray values and the rounding procedure for the reconstruction gives very accurate material boundaries, not available through any other reconstruction technique. However, the TV minimization procedure has fewer parameter selections, making initial reconstruction simpler and providing a more stable method for reconstruction. Moreover, the introduction of the TV norm has the potential for creating boundaries alternate to what a DART reconstruction would find.  Both methods are extremely valuable. In this presentation we discuss the pros and cons of each method, and show examples to illustrate when to use one method over the other.  


This research was funded in part by the DOE BES DE-SC0005822 and the LDRD and Chemical Imaging Initiative programs at PNNL. The Pacific Northwest National Laboratory is operated by Battelle under contract DE-AC05-76RL01830.

Type of presentation: Oral

IT-10-O-2962 Incoherent and coherent imaging for tomography of nano particle

Chen F. R.1, Kisielowski C.2, Tsai C. Y.1, Van Dyck D.3
11. National Tsing Hua University, Hsin-Chu, Taiwan, , 22. NCEM & JCAP Berkeley, CA, USA, , 33. University of Antwerp, Belgium
fchen1@me.com

Transmission electron microscopy (TEM) has been well recognized for its power in spatial resolution to the sub-Å level, especially, with aberration-corrected optics. However, the ultimate goal of electron microscopy is not only to obtain nice images but also to advance materials science. This means that EM has to evolve from describing to understanding materials properties. It is well-known that all the structure-property relations are encoded in the positions of the atoms and the shape of particle, specially, in the case of catalysts and biological species. The drawbacks of high resolution TEM are two folds. First, it gives only 2D projected structural information. And second, the passband of the lens transfer at the low spatial frequencies is very poor and such that the information about shape is lost.
In my talk, I will show that how we develop a novel theory and method for incoherent and coherent TEM imaging technique to determine 3D shape of nano-object with atomic resolution. For coherent imaging, our approach is to retrieve the three-dimensional atomic structure of nanocrystals from the electron exit wave function of a single projection image. The method employs wave propagation to determine the local exit surface of a sample together with the mass of each atomic column. Intensities are scaled by the mean inner potential of the sample and single atom sensitivity is expected since aberration-corrected electron microscopes are now available with such extraordinary capabilities. The validity of the approach is tested with a simulated exit wave function of a gold wedge, as shown in the fig. 1(a) and 1(b). The fig. 1(c) shows the reconstructed tomogram for the wedge Au crystal. For incoherent imaging, we present a new route to enhance the contrast by hollow cone imaging technique for biological objects using thermal diffuse scattered (TDS) electrons. Hollow cone imaging is incoherent and thus does not interfere with the central beam therefore it generates the amplitude contrast. Furthermore, the TDS signal is linear to the mass-thickness and easy to interpret and so it is suitable for soft material tomography. Here Fig. 2. we report the first results on the application of TDS to single particle analysis of proteins. The proof of the concept of the method has been demonstrated experimentally for Chaperonin GroEL as a standard protein since it is stable, easy to obtain and the structure is well-known.


CK is supported by the Office of Science, Office of Basic Energy Sciences of the U.S. Department of Energy under Contract No. DE-AC02—05CH11231. D. Van Dyck acknowledges the financial support from the Fund for Scientific Research - Flanders (FWO) under Project nos. VF04812N and G.0188.08 . F.-R. Chen would like to thank the support from NSC 96-2628-E-007-017-MY3 and NSC 101-2120-M-007-012-CC1.

Fig. 1: Fig. 1(a) modulus of he simulated exit wave for Au wedge crystal (b) phase of simulated exit wave of Au wedge crystal (c) reconstructed tomogram from 1(a) and 1(b).

Fig. 2: Fig. 2 (a) Bright field image of GroEL (b) Hollow cone image (HCI) of GroEl (c) Intensity profile across the image in 2(a) and 2(b). (d) Reconstructed tomogram

Type of presentation: Poster

IT-10-P-1482 Accuracy and applications of electron-beam deposited nano-dot fiducial markers in electron tomography of rod-shaped specimens

Hayashida M.1, Kumagai K.1, Malac M.2 3, Bergen M.2
1National Institute of Advanced Industrial Science and Technology (AIST), 2National Institute of Nanotechnology, 3University of Alberta
misa-hayashida@aist.go.jp

Electron tomography is a method employed in a transmission electron microscope (TEM) to reconstruct a three-dimensional (3D) volume from a series of images acquired at suitable tilt increments. An easy, accurate alignment of the series is critical to obtain good quality 3D reconstruction of the sample. For tomography of biological samples, Au nanoparticles are usually used as fiducial markers, which are randomly placed on the sample from a suspension. For precise alignment, markers must be uniformly dispersed over the observed region of the specimen. However, it is difficult to obtain even dispersion because the colloidal Au nanoparticles are usually dense materials. On the other hand, for high resolution imaging of nanomaterials, rod-shaped specimen is usually used, because it allows us to obtain data without missing wedge.. It is, however, difficult to disperse colloidal Au nano-particles near the observing area while not interfering with the objects of interest in such samples. Electron-beam fabricated tungsten nano-dot fiducial markers, deposited in a standard scanning electron microscope with a gas delivery system, allows placing the fiducial markers at nearby locations that do not interfere with the area of interest. In particular we discuss the accuracy of alignment using the nano-dot fiducial markers and demonstrate the application of the method to some rod-shaped specimens with nanoparticles.

<Sample> An example of a typical rod-shaped specimen with Ag nanoparticles is shown in Fig. 1a. The rod-shaped specimens were fabricated by focused ion beam. Nano-dot markers were fabricated onto the specimen using electron beam induced-deposition of tungsten from W(CO)6 precursor.

<Accuracy> As seen in Fig. 2a, the nano-dot markers have nearly parabolic cross section. However, for automatic detection of positions of markers, cross correlation of a radially symmetric template is typically used. To improve alignment accuracy, the shape of the template was changed to better reflect the shape of the markers and their projected shape for various tilts of the specimen, as shown in Fig. 1c. The total error in marker position over the entire tilt range is taken as the minimum root-mean-square (RMS) error between the expected and measured marker position in the projected images. Using the improved shape of template, the RMS error was reduced from 2 to 1.6 pixels compared to radially symmetric template.

<Reconstructed image> Figure 2b shows the cross-sectional image of Ag nanoparticles. Not only the rod shaped sample boundary appears sharp, but the individual Ag nanoparticles with less than 5 nm diameter are clearly visible in the reconstructed images.


We are grateful for financial support of AIST in Tsukuba, Japan and NINT / NRC in Canada. The ongoing support of Hitachi High Technologies Canada contributed to success of this work. We are indebted to Dr. Takashi Nakamura (AIST) for Ag NP sample preparation and Martin Kupsta for support and for assistance on the Hitachi NB 5000.

Fig. 1: (a)Image of the entire specimen. The regions are (starting from the top): amorphous carbon with nano-dot markers, ordered array of Ag nanoparticles, Si wafer substrate.(b) Radially symmetric template image.(c) Selected template image was modified to fit the shape to actual markers' shape at every tilt for every marker.

Fig. 2: Fig.2 (a),(b) a cross section of reconstructed volume of Ag nanoparticles in a plane perpendicular to the tilt axis. The plane of the cross section is marked in Fig1a.

Type of presentation: Poster

IT-10-P-1681 Quantitative 360° electron tomography analysis of mesoporous hematite nanoparticles

Winter B.1, Butz B.1, Distaso M.2, Dudák M.3, Kočí P.3, Klupp Taylor R. N.2, Peukert W.2, Spiecker E.1
1Center for Nanoanalysis and Electron Microscopy, FAU Erlangen-Nürnberg, Cauerstraße 6, 91058 Erlangen (Germany), 2Institute of Particle Technology, FAU Erlangen-Nürnberg, Cauerstraße 4, 91058 Erlangen (Germany), 3Institute of Chemical Technology, Technická 5, CZ 16628 Prague (Czech Republic)
benjamin.winter@ww.uni-erlangen.de

Metal oxide nanomaterials find application in diverse fields of research and industry such as catalysis, drug delivery, sensing, photo-electrochemistry, optical detection or as pigments [1,2].
The materials properties and hence application performance are affected by the particle size, shape, surface characteristics as well as by possible pore or defect structures. Tuning the parameters of the chosen synthesis technique can affect some or all of these properties and thus enables the tailoring of the particle function.
Here, wet-chemically synthesised mesoporous hematite nanoparticles (NPs) are investigated regarding their internal re-organization during calcination. Scanning electron microscopy (SEM) and transmission electron microscopy (TEM) studies showed that the calcination temperature has a pronounced effect on the pore morphology although a quantitative evaluation of the size distribution and connectivity of the pores was not possible (Fig. 1). Therefore, we carried out a detailed study using electron tomography (ET) in scanning TEM (STEM) imaging mode (enhanced mass-thickness contrast) in order to reveal the three-dimensional morphology, size distribution and interconnectivity of the pores inside the hematite NPs (Fig. 2 b,c,d). NPs in aqueous solution were directly drop-casted onto a pillar with an electron-transparent thickness that has been preliminarily thinned using focused ion beam (FIB) milling. Transferring this tip (Fig. 2 a) onto a 360° ET sample holder enables the acquisition of a tilt series with full tilt-angle range (Fig. 2 b). This permits the reconstruction of the particle morphology without a missing wedge of information and therefore improves the quality of the 3D reconstruction.
The tomograms were used to perform a quantitative analysis of the pore space. Virtual capillary condensation (VCC) [3] and maximum sphere inscription (MSI) [4] were applied to determine the pore size distribution of NPs as illustrated for one particle in Fig. 3. The results are compared with independent N2 sorption measurements. NPs calcined at lower temperatures tend to have a more open-porous structure with a higher degree of porosity, whereas with increasing temperature fewer and rather enclosed pores were found to occur more frequently.
The quantitative analysis of the pore morphology using 360° ET will help to get a better understanding of this promising class of material.

1. Wang et al., New J. Chem. 38 (2014), pp. 46.
2. Echigo et al., Jr., Am. Mineral. 98 (2013), pp. 154.
3. Štěpánek et al., Colloids Surf., A 300 (2006), pp. 11.
4. Novák et al., Chem. Eng. Sci. 65 (2010), pp. 2352.


DFG Research Training Group 1161, Cluster of Excellence “Engineering of Advanced Materials” (EXC 315), DFG SPP 1570, Czech Science Foundation (GACR P106/10/1568).

Fig. 1: a) SEM and TEM images of hematite nanoparticles calcined at 400°C and 500°C.

Fig. 2: a) Photo of a 360° ET tip with a droplet of hematite particles in solution; b) STEM image of Fe2O3 nanoparticles attached to this tip - green frames tag particles calcined at 400°C, yellow frames at 500°C; c), d) comparison of STEM images and the 3D reconstruction (volume rendering) of c) 400°C and d) 500°C particles.

Fig. 3: Reconstructed hematite nanoparticle calcined at 400°C showing a) a snapshot of the virtual capillary condensation (VCC) [3] (pores gradually filled with liquid (blue)), b) a slice displaying colour-coded results from the MSI analysis [4]; c) diagram showing the pore size distribution resulting from the MSI analysis and the VCC simulation.

Type of presentation: Poster

IT-10-P-1698 The physical properties of plastic support films for 3-D transmission electron microscopy.

Daraspe J.1, Longo G.2, Kizilyaprak C.1, Humbel B. M.1
1University of Lausanne, Electron Microscopy Facility, Lausanne Switzerland, 2EPFL, Laboratory of Physics of Living Matter, Lausanne, Switzerland
jean.daraspe@unil.ch

In the last 10 years, the acquisition of large volumes of biological specimens at high resolution has regained importance, especially in the field of neurobiology. The analysis of large volumes has become essential to better understand cell-cell relationship and the interaction of subcellular organelles.
Though more modern and automated processes like serial block face scanning electron microscopy [1] and focused ion beam scanning electron microscopy [2] ease the process of gaining large volumes, traditional serial section has come to a revival [3].
Serial sectioning maintains its ability to image almost an unlimited 3D volume, up to mm2 at high resolution (X=Y=~1nm; Z=50-70nm). Further serial thick section TEM tomography can improve the Z resolution to about 2nm. The only limit is the skill and persistence of the operator.
For serial section image acquisition and for serial TEM tomography the stability of the plastic support film of the TEM grids, especially large slot grids, is crucial. The following quality criteria are required:
- Flatness, the films should not bend during pick-up of the sections.
- Resilient and strong, the films have to support mechanical tensions.
- Beam resistance, the films should not drift and disrupt in the electron beam.
- Temporal stability, these qualities should be maintained for a long time to allow storage of prepared grids.
To improve the stability of the section on the support film it is important to review the properties of the different polymers commonly used and to find the formulation, which best matches the quality criteria required.
We prepared support films from 6 different polymers and analysed their behaviour during pick-up and TEM imaging. Plastic films were casted on microscope slides and two films of equal thickness were separated. One was used as a support film on a TEM grid and the other was mounted on a slide for AFM analysis. In the AFM the thickness, stiffness and adhesion force was measured. In the TEM, drift measurements were done by following gold particles at high magnification using time lapse series acquisition. Thickness was also measured by TEM tomography.
In conclusion, two polymers emerged that fulfil the requested criteria for 3D investigations by serial sectioning.

1 - Serial block-face scanning electron microscopy to reconstruct three-dimensional tissue nanostructure. Denk W et al.; PLoS Biol. 2004 Nov;2(11):e329
2 - Serial section scanning electron microscopy of adult brain tissue using focused ion beam milling. Knott G et al.; J Neurosci. 2008 Mar 19;28(12):2959-64
3 - Array Tomography: A New Tool for Imaging the Molecular Architecture and Ultrastructure of Neural Circuits. Micheva KD et al; Neuron. 2007 Jul 5;55(1):25-36


Type of presentation: Poster

IT-10-P-1846 A Remote Control/Observation System and an Operation Support System for the Ultrahigh Voltage Electron Tomography

Yoshida K.1, Nishi R.1, Yasuda H.1
1Research Center for Ultra High Voltage Electron Microscopy, Osaka University
yoshida@uhvem.osaka-u.ac.jp

The full text of the abstract is not available. Please contact the presenting author.


Type of presentation: Poster

IT-10-P-1866 A memory efficient method for 3D object reconstruction with HAADF STEM depth sectioning

Van den Broek W.1, Rosenauer A.2, Van Aert S.3, Sijbers J.4, Van Dyck D.3
1Institute for Experimental Physics, Ulm University, Albert-Einstein-Allee 11, 89081 Ulm, Germany, 2Institut für Festkörperphysik (IFP), Universität Bremen, Otto-Hahn-Allee 1, 28359 Bremen, Germany, 3Electron Microscopy for Materials Science (EMAT), University of Antwerp, Groenenborgerlaan 171, 2020 Antwerp, Belgium, 4iMinds - Vision Lab, Department of Physics, University of Antwerp, Universiteitsplein 1, 2610 Antwerp, Belgium
wouter.vandenbroek@uni-ulm.de

In high angle annular dark field scanning transmission electron microscopy (HAADF STEM) depth sectioning, the object is illuminated with a beam with a small depth of field, caused by a large beam convergence angle, see Fig. 1.

In [Ultram. 110 (2010) 548--554], it is shown that the simultaneous iterative reconstruction technique (SIRT) applies to depth sectioning. The SIRT algorithm is given as

fk+1 = fk + AT [ ( q - A fk ) / A If ] / [ AT Ip ],

where k indicates the iteration number, A is the projection matrix, f is the object vector and q the experimental projection, arithmetic operators between vectors are elementwise and Ip and If each denote a vector of which each element equals 1 with a length equal to that of the projection and the object, resp.

The matrix A scales badly with object size and readily grows too large for the computer memory. Here, we propose to perform the matrix-vector multiplication A fk implicitly through a 2D convolution of each horizontal layer of the object with its corresponding single atom image, followed by a sum in the vertical direction over all horizontal layers. Implementation of the matrix-vector multiplication A If is analogous. The multiplication of AT with the vector ( q - A fk ) / A If can be carried out implicitly by stacking the image in an 3D array and convolving each of the layers with the corresponding single atom image. The matrix-vector multiplication AT Ip can be implemented analogously. The memory load is now reduced to storing the object.

A technique analogous to charge flipping [Acta Cryst. A64 (2008) 123--134] is added: After each iteration values below a certain positive threshold have their sign reversed.

The validity of these implicit matrix-vector multiplications is tested with a simulation of a 1.6nm Au particle. The microscope parameters are: Acceleration voltage: 200kV; C1: -2.67nm; C3: 3.54μm; C5: -1.13mm; C7: 10cm; convergence semi-angle: 86.8mrad. The object measures 1000 x 1000 x 125 voxels of 8 x 8 x 210 pm3. See Fig. 2. The particle is tilted away from the zone-axis. The single atom images encoded in the matrix A are a convolution of the Au atom potential and the probe intensity.

A total of 125 images with a defocus step of 0.21nm is simulated and used as input for SIRT with 64 iterations. In Fig. 3 it is shown that the severe elongation in the vertical direction, so typical for depth sectioning, is overcome. The authors have reported these results in [A memory efficient method for fully three-dimensional object reconstruction with HAADF STEM, Ultram., accepted].


W. Van den Broek: The Carl Zeiss Foundation and DFG, KO 2911/7-1; A. Rosenauer: DFG, AR 2057/8-1; S. Van Aert, J. Sijbers, D. Van Dyck: FWO, G.0393.11, G.0064.10, G.0374.13.

Fig. 1: Features of the object close to the beam crossover are well localized in the image while features further away are smeared out. Depth information is unlocked by varying the defocus, thus bringing other regions of the object in focus. The underlying assumption is that the image formation is incoherent.

Fig. 2: Upper left: Horizontal slice through the middle of the data set. Lower left: Average of all horizontal slices. Right: Average of the vertical slices. The white oval marks the particle's position.

Fig. 3: Depth-profiles of the simulated measurements, the reconstruction and the original object averaged over a disc of radius 1.6nm centered on the particle. The FWHM of the profile of the reconstruction (2.07nm) now approximates that of the original object (1.85nm).

Type of presentation: Poster

IT-10-P-1875 Improving Depth of Focus in STEM Tomography using Focal Series

TREPOUT S.1, MESSAOUDI C.1, BASTIN P.2, MARCO S.1
1Institut Curie / INSERM U759, Campus Universitaire d'Orsay, Bât. 112, 91405 ORSAY cedex FRANCE, 2Institut Pasteur, CNRS URA 2581, Parasitology & Mycology Department, Institut Pasteur, 25, rue du Docteur Roux, 75015 PARIS, FRANCE
sylvain.trepout@curie.fr

Scanning Transmission Electron Microscopy (STEM) is a point to point imaging method which uses a focused electron beam to build a projection image of the sample. One of the advantages of STEM is that the focused electron beam can pass through thicker samples (up to 1 µm-thick on a 200 kV FEG) providing higher signal-to-noise ratio than standard TEM (wide spread beam). This makes STEM more suitable for electron tomography of thick samples. Nevertheless, the smaller is the size of the probe, the more reduced is the depth of focus (DOF). A way of solving this difficulty is to take advantage of raster scanning to get focused images line by line by dynamic focus [1]. However, in our experience, the recovery of full-focused images by dynamic focus is not helpful when the DOF cannot encompass a very thick sample (>0.5 µm). To circumvent this limitation we have developed an acquisition scheme and an image processing method in which we reconstruct full-focused images from STEM images recorded at different defocus.
This process of DOF correction can be used for both 2D studies and tomographic experiments. Briefly, it consists in two steps. For a 500 nm thick sample, i) five images are acquired at different focus values ranging from -300 nm up to 300 nm defocus (using 150 nm steps); ii) the images are processed using an ImageJ macro based on Turboreg [2] and Extended Depth of Field [3]. This macro consists of two steps: i) the images acquired at different focus are aligned; ii) regions of these registered images which are at focus are combined to compute a single image.
The performance, in terms of resolution, contrast and SNR, of the afore mentioned approach has been evaluated in silico by comparing 3D reconstructions computed from the projections of a phantom volume leading to two different datasets. In the first dataset the whole images are at focus whereas in the second one, parts of the images have been modified to mimic the focus changes observed in experimental data. The method was applied to compute the 3D reconstruction of a resin-embedded T. brucei 500 nm-thick section using five focal series. Sections of the reconstructed volumes without (1 focal series), with intermediate (3 focal series) and full (5 focal series) DOF correction on high-tilt images (±40°) are displayed in figure 1 showing the improvement in the observed details.

References:
1. Feng et al. “Automated electron tomography with scanning transmission electron microscopy”. J. Microsc., 2007, 228:406-12.
2. Thévenaz et al. "A Pyramid Approach to Subpixel Registration Based on Intensity". IEEE Tr. Im. Pro., 1998, 7: 27-41.
3. Forster et al. « Complex Wavelets for Extended Depth-of-Field: A New Method for the Fusion of Multichannel Microscopy Images”. Mic. Res. Tech., 2004, 65:33-42.


This work has been funded by the ANR grant 11-BSV8-0016.

Fig. 1: Comparison of 3D information recovered from depth of focus correction of high-tilt images. a, b, c) 1 nm-thick XY planes from reconstructions computed with only high-tilt images (±40°) without, with intermediate and full depth of focus correction respectively. d, e, f) 1 nm-thick YZ planes from the same reconstructions. Scale bar, 100 nm.

Type of presentation: Poster

IT-10-P-1904 Electron Holographic Tomography of Electric and Magnetic Stray Fields around Nanowires

Lubk A.1, Wolf D.1, Lichte H.1
1Triebenberg Laboratory, TU Dresden, Dresden, Germany
Axel.Lubk@triebenberg.de

One important consequence of the support theorem of the Radon transformation is that the so-called outer Radon problem (ORP), stating that the tomographic reconstruction of a scalar function in the exterior of a convex domain from projections along submanifold passing outside that domain, has a unique solution. Cast into the specific circumstances of electron holographic tomography that means that a reconstructed phase tilt series from outside a charged nanostructure, such as a nanowire (NW), is sufficient to reconstruct the corresponding potential in that outer region (Fig. 1). The topic of this contribution is the solution of the ORP, including its technical implementation and benefits for the determination of physical quantities within the framework of electron holographic tomography. Until now, only the interior Radon problem, i.e. the reconstruction of potentials within a convex (circular) domain containing e.g. a pn-junction has been treated. There, a number of issues which can be potentially tackled with the help of the ORP have been encountered: (A) Tilt angles within the TEM are often limited to within approx. -70° to 70°, rendering spatial frequencies in the corresponding missing wedge in Fourier space inaccessible. (B) Phase unwrapping algorithms cannot distinguish between unresolved phase gradients larger then π and phase jumps, producing artefacts at object boundaries (e.g. at FIB prepared specimen). (C) During tilting the proximity of a low-index zone axis can lead to dynamical scattering. (D) Magnetic field reconstruction suffers from the difficult alignment of 2 by 180° flipped tilt series required to seperate electric and magnetic phase shifts. The lowered influence of these issues in the ORP is of course purchased by disregarding any information from the blocked specimen region. However, the laws of electro-(magneto)statics relate the outside potentials to object charges and dipoles, thus rendering the fringing fields a valuable property. We will demonstrate a solution to the ORP including its potential to mitigate the above issues by reconstructing a small beam induced charging field around a core-shell GaAs-AlGaAs NW (Fig. 2) and a magnetic field emerging from the tip of a Co2FeGa Heusler alloy NW (Fig. 2) into vacuum. Both fields are very weak and could only partially recovered by means of the interior Radon problem due to the above issues.


The authors acknowledge financial support from the European Union under the Seventh Framework Program under a contract for an Integrated Infrastructure Initiative. Reference 312483 - ESTEEM2.

Fig. 1: Geometry and coordinate system employed in the ORP. D indicates the reconstruction domain containing the function f. The projection is denoted by F and in both functions f and F the dark gray region indicates the part excluded within the ORP.

Fig. 2: Electrostatic potential and axial component of the magnetic field reconstructed in the exterior of two NWs by solving the ORP.

Type of presentation: Poster

IT-10-P-1980 Energy dispersive X-ray (EDX) tomography of bimetallic nanoparticles

Slater T. J.1, Macedo A. M.2, Burke M. G.1, O'Brien P.1, Camargo P. H.2, Haigh S. J.1
1The University of Manchester, Manchester, UK, 2Universidade de Sao Paulo, Sao Paulo, Brazil
thomas.slater-5@postgrad.manchester.ac.uk

Electron tomography can be used to provide spectroscopic analysis in three dimensions at nanometer resolution through a variety of imaging techniques. EDX tomography promises accurate simultaneous projections of all elements that fully meet the projection requirement. We present the latest results of three-dimensional elemental analysis using EDX tomography to investigate bimetallic nanoparticles.
The design of the Super-X detector [1] as included on the FEI Titan G2 80-200 has a large solid angle of detection (≈0.8sr) and therefore a count rate high enough to allow EDX tomography of many types of sample. EDX tomography of focused ion beam (FIB) prepared samples is now entirely feasible and results in minimal detector shadowing at any angle [2]. However, FIB prepared rods of free standing nanoparticles are not straightforward to prepare. Dispersing nanoparticles on to a standard TEM grid is a simple preparation method but the grid bars and sample holder shadow the EDX detectors at a range of angles. This prevented acquisition of EDX data at a large range of angles when using traditional single detectors but can be overcome when using the Super-X detector. To combat effects of detector shadowing we have used a simple acquisition-time varying sampling scheme that is guided by prior characterisation of the detector. We acquired EDX spectral images of a single AgAu nanoparticle at a range of tilt angles for a fixed acquisition time. Acquisition time at each tilt angle was then adjusted to provide similar Au counts at each angle (Fig. 1).
We have used our novel acquisition scheme to investigate elemental distributions in bimetallic nanoparticles synthesized via the galvanic replacement reaction, primarily AgAu nanoparticles. The elemental distribution in these nanoparticles is particularly important for their catalytic and optical properties. Two dimensional EDX mapping suggested that the extent of surface segregation in the AgAu nanoparticles varied with composition. Particles with low Au content (below 20 at% Au) appeared to display Au surface segregation and particles with high Au content (above 40 at% Au) displayed Ag surface segregation. However, two dimensional maps cannot categorically reveal surface compositions. For example, it is unclear whether intense lines of Au counts in Fig. 2a are situated on the surface of or within the nanoparticle. For this reason, we performed EDX tomography on nanoparticles that showed, separately, Au and Ag surface segregation (Fig. 2c,d). Through EDX tomography we were able to conclusively show a reversal in surface segregation in AgAu nanoparticles prepared via the galvanic replacement reaction.

References
1 von Harrach, H. S. et al. Microsc. Microanal. 15 (2009), 208-209
2 Lepinay, K. et al. Micron 47 (2013), 43-49


SJH thanks the USA Defense Threat Reduction Agency (grant number HDTRA1-12-1-0013) and Gates Foundation for funding support. PHCC and AM thank FAPESP and CNPq for funding support (grant numbers 2011/06847-0, 2013/19861-6 and 471245/2012-7, respectively).

Fig. 1: Figure 1. a) Counts of Au and Ag peaks as a function of tilt angle when using a fixed time acquisition scheme. (b) Varied acquisition-time scheme used in order to correct for detector shadowing by the sample holder. (c) Counts of Au and Ag peaks as a function of tilt angle when the varied acquisition-time scheme is used.

Fig. 2: Figure 2. a,b) Two dimensional EDX spectral images showing Au and Ag distributions within nanoparticles with Au and Ag surface segregation. c,d) Surface visualisations of reconstructed tomograms of Au and Ag in nanoparticles displaying Au and Ag surface segregation.

Type of presentation: Poster

IT-10-P-2061 Whole-Cell Imaging of the Budding Yeast Saccharomyces cerevisiae by High-Voltage Scanning Transmission Electron Tomography

Murata K.1, Esaki M.2, Ogura T.2, Arai S.3, Yamamoto Y.3, Tanaka N.3
1National Institute for Physiological Sciences, Okazaki, Aichi, 444-8585, Japan, 2Institute of Molecular Embryology and Genetics, Kumamoto University, Kumamoto, 860-0811, Japan, 3Ecotopia Science Institute, Nagoya University, Nagoya, Aichi, 464-8603, Japan
n-tanaka@esi.nagoya-u.ac.jp

High-voltage electron tomography provides three-dimensional (3D) information about cellular components in thicker sections beyond 1 μm, but image degradation caused by multiple inelastic scattering of transmitted electrons limits the attainable resolution. Scanning transmission electron microscopy (STEM) is believed to give enhanced contrast compared to conventional transmission electron microscopy (CTEM), and thicker samples up to around 1 μm can be analyzed with an intermediate-voltage electron microscope, because the depth of focus and the inelastic scattering are not the critical limitations.
      We have applied STEM to 1 MV high-voltage electron tomography to extend the limitation of the specimen thickness, and seamlessly investigated the whole-cell structure of the budding yeast, Saccharomyces cerevisiae, size of which was ~3 μm width. The high-voltage STEM tomography, especially with a bright-field mode, demonstrated sufficient enhanced contrast and more-intense signals compared to regular TEM tomography (Fig. 1), permitting segmentation of major organelles in the entire cell (Fig. 2). The technique also showed less specimen shrinkage. The current spatial resolution is limited with the specimen preparation and the relatively large convergence angle of the scanning probe, but the present new technique has a potential to solve longstanding problems of image blurring in thick biological specimens beyond 1 μm, and to open a new research field in cell structural biology.
[1]K. Murata et al., to be submitted (2014).


This study was supported by the program of Joint Usage/Research Center for Develpment Medicine at IMEG, Kumamoto University. H-1250M at NIPS and JEM-1000K RS at Nagoya University were used for observation.

Fig. 1: Tomogram slices of an whole budding yeast cell at xy (a) and xz (b) planes were calculated from the tilt series collected by TEM-BF(Blight field), STEM-BF, and STEM-ADF(Annular dark field), respectively. Scale: 1 μm.

Fig. 2: The major organelles of a budding yeast cell were successfully segmented in a 3D tomogram calculated from the STEM-BF tilt series. Scale: 1 μm.

Type of presentation: Poster

IT-10-P-2098 Cryo-STEM Tomography for 3D Analysis of Cell Structure

Aoyama K.1,2
1Application Laboratory FEI Japan, Tokyo, Japan, 2Osaka University, Osaka, Japan
kazuhiro.aoyama@fei.com

Recently, several studies for observation of biological specimens as plastic section have been performed by using STEM, and the potential has been indicated. STEM tomography offers several important advantages including: (1) it is effective even for thick specimens, (2) ‘dynamic focusing’, (3) ease of using an annular dark field (ADF) mode and (4) linear contrasts. It has become evident that STEM tomography offers significant advantages for the observation of thick plastic specimens. In this study, the technique applied for Cryo-specimens. Of course, even in Cryo-Tomography, the advantages of STEM above mentioned are valid. Because STEM has advantage to resist specimen thickness, it is expected to be powerful method for observing whole cell structure in Cryo-microscopy without thin sectioning.

The insufficient contrast is one of the serious problems in Cryo-electron microscopy. Therefore, the image contrasts by TEM and STEM have to be compared carefully and quantitatively. Figure 1 shows the comparison of the image qualities. The specimen was vitreous ice on Quantifoil made by standard procedure of Vitrobot. Titan Krios equipped with 2 cameras and STEM system was used for the experiment. TEM images were taken by CCD camera (Gatan US4000) and direct detected CMOS camera (FEI Falcon). STEM image was taken in bright field mode. The imaging conditions, image pixel size and the number of irradiation electrons, were normalized. The average counts of the pixels and the standard deviation (SD) were measured for each image, and then SD/mean was calculated. The result was clear that the STEM image had very low back ground noise. This character can be explained theoretically, there are several seasons; (1) Short operation time for pixel (dual time vs. exposure time), (2) Large physical size of the detector, (3) Very small collection angle (same as very small objective aperture in TEM imaging).

By applying Cryo-STEM tomography, clear membrane structure of organelle appeared without staining and without sectioning.


Fig. 1: A comparison of the image quality in Cryo-EM

Type of presentation: Poster

IT-10-P-2123 3D electron tomography analysis of silicon nanoparticles in SiC matrices by quantitative determination of EELS plasmon intensities

Xie L.1, Jarolimek K.2, Van Swaaij R.2, Thersleff T.1, Zeman M.2, Leifer K.1
1Uppsala University, Department of Engineering Sciences, Applied Materials Sciences, Box 534, SE-751 21 Uppsala, Sweden., 2Photovoltaic Materials and Devices, Delft University of Technology, P.O. Box 5031, 2600 GA Delft, The Netherlands.
ling.xie@angstrom.uu.se

Silicon nanoparticles (NPs) embedded in insulating or semiconducting matrices has attracted much interest for the third generation of photovoltaics, “all-Si” tandem solar cells. This study is to show how silicon NPs are distributed in 3D on a silicon carbide thin film using the electron tomography technique in the transmission electron microscopy (TEM). [2]

We first have assessed Si NPs distributions in such SiCx sample with a low degree of crystalline using bright field (BF) TEM tomography (figure 1) and found an average nearest neighbour spacing of two NPs of about 12nm. For more crystalline NPs, the projection requirement is no more fulfilled and only those Si NPs that are both crystalline and oriented to a Bragg reflection are detectable. [3] Therefore, in this case, conventional BF TEM signal is unsuitable for electron tomography and we applied spectrum imaging (SI) techniques: EELS SI imaging and EFTEM SI imaging. Since Si and SiCx have different plasmon energies, [4] we can extract Si plasmon and SiCx plasmon images from the spectrum images. We observed that only a proper fit of the plasmon spectrum with subsequent extraction of Si and SiCx plasmon images results in the correct Si ad SiCx distribution (figures 2 and 3), whereas just EFTEM images taken from windows around the Si and the SiC plasmon energy resulted in overlaps in the image.

For both, STEM and EFTEM SI signals, in figure 2 and 3, we are able to detect the entire population of NPs. In figure 3, the stripes like contrast inside of crystalline NPs shown in the BF TEM image persist in plasmon images. This is due to parallel beam illumination in EFTEM SI mode thus making the STEM SI imaging more suitable for tomography of these NPs. In Figure 2, for STEM SI, the contrast evolution during the tilting is thickness dependent, thicker part of the sample gives stronger contrast in the extracted plasmon images, and this nonlinear thickness effect can be corrected by introducing attenuation coefficient. [5]

In summary, to study the 3D distribution of Si NPs in SiCx matrix, we compared three signals from BF TEM, STEM and EFTEM SI signals. In order to overcome the non-linearity of contrast change during the tilting process, STEM-SI signal in combination with quantitative treatment of the plasmon spectra shows clear Si NP contrasts and overcomes limits set by the projection requirement.

[1] S. Perraud et al., Phys. Status Solidi A, 1–9 (2012).

[2] J. Frank, Electron Tomography: Three Dimensional Imaging with the Transmission Electron Microscope, Plenum, New York, London, 1992.

[3] P. A. Midgley et al., Ultramicroscopy 96 (2003) 413.

[4] R.F. Egerton, Electron Energy-Loss Spectroscopy in the Electron Microscope, 420, 2011.

[5] W. Van den Broek et al. Ultramicroscopy 116 (2012) 8–12


The authors acknowledge the support from the EU founded FP8 project “SNAPSUN”.

Fig. 1: Figure 1. (Left) 2D BF TEM image of Si nanoparticles embedded in amorphous Si riched SiC:H matrix, (Right) 3D model view of Si nanoparticles distributed in matrix, green spheres indicate Si nanoparticles, and reconstructed X, Y and Z images are also shown in the volume. Scale bar is 10 nm.

Fig. 2: Figure 2. (Left) 2D STEM-ADF image of Si nanoparticles embedded in SiCx matrix, (Middle) Si plasmon image, (Right) SiC plasmon image. Both plasmon images are extracted from STEM SI data set.

Fig. 3: Figure 3. (Left) 2D BF TEM image of Si nanoparticles embedded in SiCx matrix, (Middle) Si plasmon image, (Right) Reconstructed tomogram and the formation of Si networks were shown in the volume by using the isosurface. Plasmon images are extracted from EF TEM SI data set which is acquired with a 2 eV energy slit at 17 eV (Si).

Type of presentation: Poster

IT-10-P-2238 EELS CS tomography of FeO-Fe3O4 core-shell nanoparticles. An approach to recover 3D oxidation state distribution.

Torruella P.1, Arenal R.2,5, Saghi Z.3, Yedra L.1,4, Eljarrat A.1, de la Peña F.3, Midgley P. A.3, Estradé S.1,4, Peiró F.1, Estrader M.6, López-Ortega A.7, Salazar-Alvarez G.8, Nogués J.9,10,11
1LENS–MIND-IN2UB, Dept.d’Electrònica, Universitat de Barcelona, Barcelona, Spain , 2LMA, Instituto de Nanociencia de Aragon, Universidad de Zaragoza, Zaragoza, Spain, 3Department of Materials Science and Metallurgy, University of Cambridge, Cambridge, United Kingdom, 4CCiT, Universitat de Barcelona, Barcelona, Spain, 5Fundacion ARAID, Zaragoza, Spain, 6Departament de Química Inorgànica, Universitat de Barcelona, Barcelona, Spain, 7INSTM and Dipartimento di Chimica “U. Schiff”, Università degli Studi di Firenze, Firenze, Italy, 8Department of Materials and Environmental Chemistry, Arrhenius Laboratory, Stockholm University, Stockholm, Sweden, 9Departament de Física, Universitat Autònoma de Barcelona, Bellaterra (Barcelona), Spain, 10ICN2 – Institut Catala de Nanociencia i Nanotecnologia, Campus UAB, Bellaterra (Barcelona), Spain, 11Institució Catalana de Recerca i Estudis Avançats (ICREA), Barcelona, Spain
pautorruellabesa@gmail.com

The aim of this work was to characterize a sample consisting on FeO-Fe3O4 core-shell cubic-shaped nanoparticles. Because of the similarities in the composition and effective atomic number of the core and the shell, high angle annular dark field (HAADF) imaging could not be used to resolve the structure.

As an alternative, EELS fine structure can be used to obtain information on the oxygen and iron oxidation state, thus making it possible to distinguish between FeO and Fe3O4. However, there is the limitation that EELS projects the information of the 3D nanoparticle into a 2D map.

To overcome this limitation there is the possibility to consider EELS spectrum image data-sets as suitable for 3D tomographic reconstruction, not only containing information on the chemical composition of the sample (as in [1]) but also on the oxidation state of Fe at each voxel.

A tilt series of spectrum images (SI) was acquired on a probe corrected FEI Titan. Then the images were treated with Hyperspy to obtain independent spectral components from the iron edge, with could be correlated with the different iron oxides. In order to improve the quality of the reconstruction, a new reconstruction algorithm based on the mathematical theory of compressed sensing (CS) was used. To our knowledge this is the first time that the CS algorithm has been used to reconstruct an EELS core-loss spectrum image data-set.

The CS reconstructions show a shell thickness of 9nm around the core. The 3D reconstruction proves a total shell coverage of the core and that there has been no appreciable phase mixing.

[1] Ll. Yedra et al., Ulramicroscopy 122 (2012), pages 12-18.


The measurements were performed in the Laboratorio de Microscopias Avanzadas (LMA) at the Instituto de Nanociencia de Aragon (INA) - Universidad de Zaragoza (Spain). We acknowledge the support received from the European Union Seventh Framework Program under Grant Agreement 312483 - ESTEEM2 (Integrated Infrastructure Initiative – I3) and under Grant Agreement 291522-3DIMAGE.

Fig. 1: Obtained spectral components from the iron edge after performing PCA and ICA with Hyperspy.

Fig. 2: Central orthoslice from CS reconstruction corresponding to ‘core’ component in figure 1 showing core thickness measurement.

Fig. 3: Central orthoslice from CS reconstruction corresponding to ‘shell’ component in figure 1 showing shell thickness measurement.

Type of presentation: Poster

IT-10-P-2280 Tomography in Analytical Transmission Electron Microscopy of Nanomaterials

Orthacker A.1, Haberfehlner G.1, Tändl J.3, Poletti M. C.3, Kothleitner G.1,2
1Center for Electron Microscopy, Graz, Austria, 2Institute for Electron Microscopy and Nanoanalysis, Graz University of Technology, Graz, Austria, 3Institute for Materials Science and Welding, Graz University of Technology, Graz, Austria
angelina.orthacker@felmi-zfe.at

The engineering of specific material properties requires both a structural and chemical understanding, often obtained with electron microscopic techniques. While analytical scanning transmission electron microscopy (STEM), including energy dispersive x-ray (EDX) spectroscopy and electron energy loss spectroscopy (EELS), can offer the necessary chemical information, the integrative character of the signal acquired through transmission might hide important structural details of the material. Those details can be revealed by using electron tomography, where the data is acquired at different tilt angles and, after alignment, reconstructed to form a full 3D model of the investigated material. This technique on its own, however, lacks the important chemical information. The content of this work is the combination of both techniques, analytical STEM and tomography, which offers a more complete understanding of the material structure and composition.
Combining analytical signals in the form of spectrum images with a tomographic acquisition however, represents a major challenge. First of all there is no possibility to acquire a tilt series including EELS and EDX spectra for each data point automatically. Secondly, there is no simple way of reconstructing a four dimensional object, containing two spatial, one energy and one tilt angle coordinate. Having tackled this problem, the next issue arises due to the often very limited statistical quality of the analytical signals. In order to minimize acquisition times, dose and sample drift, dwell times per pixel are often in the order of a few milliseconds, necessitating special reconstruction algorithms that can handle noisy data.
The material studied in this work is an alloy containing scandium and zirconium rich nanoparticles embedded in an aluminum-magnesium matrix (fig.1 and fig.2). These nanoparticles increase the mechanical resistance of the material. Their sizes and chemical compositions can vary, depending on the aging process. Previous work reported that these particles can exhibit a core-shell structure.
As the acquisition of EDX and EELS spectra takes more time than non-spectroscopic imaging techniques special care needs to be taken concerning the stability of the sample. This can be achieved by reducing the number of tilt angles, which is possible if special reconstruction algorithms are used. Total variation minimization mathematically assumes minimized gradients which can reduce artefacts and noise. While this can be problematic if the sample itself exhibits gradients of concentrations, it can lead to smooth reconstructions if the sample consists of clearly separable phases.


We thank the Austrian Cooperative Research Facility, the European Union (7th Framework Programme: ESTEEM2), and the Austrian Research Promotion Agency FFG (TAKE OFF project 839002) for funding.

Fig. 1: aluminum alloy; left: EFTEM overview image at 40eV; right: HAADF image of Sc-rich nanoparticle

Fig. 2: extracted spectra of an EELS (left) and EDX (right) spectrum image of a nanoparticle in its matrix, supporting that the particle is Sc-rich

Type of presentation: Poster

IT-10-P-2420 Tilt-less Electron Tomography

Oveisi E.1, Letouzey A.2, Lucas G.1, Cantoni M.1, Schäublin R.3, Fua P.2, Hebert C.1
1Interdisciplinary Centre for Electron Microscopy, Ecole Polytechnique Fédérale de Lausanne (EPFL), Switzerland, 2Computer Vision Laboratory, Ecole Polytechnique Fédérale de Lausanne (EPFL), Switzerland, 3Centre de Recherches en Physique des Plasmas, Ecole Polytechnique Fédérale de Lausanne (EPFL), Switzerland
emad.oveisi@epfl.ch

Accurate three-dimensional (3D) knowledge of dislocation structures is more and more vital for the understanding of complex deformation mechanisms at the nanometer scale. In this respect, tomography in transmission electron microscopy developed in the area as a fruitful technique, but is demanding and sometimes inapplicable, as it requires the acquisition of multiple images within a large tilt range [1-4]. The major challenge to electron microscopists now is the development of efficient techniques to overcome these limitations and thus facilitate 3D reconstructions.
We have developed an efficient method that provides a highly reliable insight into the 3D reconstruction of dislocations, in which both image acquisition and reconstruction are addressed. This technique makes use of the convergent electron beam in scanning transmission electron microscopy (STEM) and provides from a single viewing direction a stereoscopic pair of micrographs. Our newly reconstruction algorithm allows us to derive the true structure of dislocations in three dimensions on the sole basis of the acquired stereoscopic micrographs. The reconstruction algorithm firstly extracts, from the STEM stereo images, the dislocation lines using state of the art curvilinear structures detection algorithm. These 2D representations of dislocations are then automatically matched between images. Finally, the algorithm, given the projection parameter and the corresponding virtual tilt angle of each image, can then reconstruct the 3D structure of the dislocations. The method is successfully demonstrated on the 3D visualization of dislocation arrangements in a Fe-10Cr model alloy.
With a proper input of the crystallographic axes of the specimen in the algorithm, the visualization tool allows determining the habit planes of the curves in the dislocation lines. The missing wedge effects are reduced in the 3D structure reconstructed via this algorithm, and “replaced” by a small uncertainty along the Z-axis, that thanks to purposely-designed smoothing techniques applied in the algorithm, can be kept in the range of few pixels.
In summary, this efficient and straightforward technique will markedly facilitate the 3D reconstruction and, if prior knowledge about the object can be ascertained, can be extended to a wide range of fields in which tilting associated problems, specimen thickness, and sensitivity to electron radiation are the limiting factors.

[1] P.A. Midgley, R.E. Dunin-Borkowski, Nature Materials 8 (2009) 271.
[2] J.S. Barnard, J. Sharp, J.R. Tong, P.A. Midgley, Science 313 (2006) 319.
[3] M. Tanaka, M. Honda, S. Hata, K. Higashida, Materials Transactions 49 (2008) 1953.
[4] M. Weyland, P.A. Midgley, Materials Today 7 (2004) 32.


Prof. P. Stadelmann, Dr. D. Alexander, and Q. Jeangros are gratefully acknowledged for stimulating discussions. This work was financially supported by Swiss National Science Foundation for scientific projects.

Fig. 1: Bright field image of dislocations in a Fe-10Cr model alloy.

Fig. 2: Volume-rendered image of dislocations viewed along various directions. The green, blue, and red arrows respectively correspond to the principal crystalogrphic [100], [010], and [001] axes.

Type of presentation: Poster

IT-10-P-2426 Nonlinear intensity attenuation in bright-field TEM images and its influence on tomographic reconstruction of micron-sized materials

Yamasaki J.1, Mutoh M.1, Ohta S.2, Yuasa S.2, Arai S.1, Sasaki K.1, Tanaka N.1
1Nagoya University, Nagoya, Japan, 2JEOL Ltd., Akishima, Japan
yamasaki@esi.nagoya-u.ac.jp

    Currently, tomography in transmission electron microscopes (TEM) is widely applied to three-dimensional (3D) analyses of nanometer-sized and sub-micron-sized materials. One of the next methodological targets should be quantitative 3D reconstructions in which not only the shape but also the internal density are correctly reproduced. This is hindered generally by the nonlinearity between projection thickness and image intensity. In the case of mass-thickness contrast in bright-field TEM (BF-TEM) images, the ideal exponential attenuation of the image intensity with increasing thickness is disturbed by multiple scatterings.
    In the present study, the nonlinear attenuation in BF-TEM images was analyzed using amorphous carbon microcoils (CMCs) [1] shown in Fig. 1(a). Their well-defined shapes and compositional homogeneity are quite useful for estimating the mass-thickness [2]. The intensity attenuation was measured along the line in Fig. 1(b), which was taken by the high-voltage electron microscope in Nagoya University [3]. The results measured at the acceleration voltages of 400, 600, 800 and 1000 kV were converted to the plots of the electron transmittance T in Fig. 2. At a glance, T at any voltages seems to undergo the linear attenuation. However, the least squares fitted line for the data at 400 kV exhibits a considerably negative intercept value at zero thickness. Such nonlinear attenuation should induce failures in conversion from intensity to thickness and thus inhibits correct 3D reconstructions of the internal density.
    The influence of the nonlinearity on tomographic reconstructions was examined using a 360°-tilt sample holder, which was specially developed for eliminating the missing-wedge effect [2]. Figure 3 shows the results of the reconstructions from the tilt series taken at 400 kV and 1000 kV. Although the 3D shape of the CMC has been reconstructed well in both cases, the internal density is not uniform but has a gradient from the center at 400 kV. Moreover, there is a slight increase in the vacuum level in the interior of the coil. The inaccurate density reconstruction should result from the nonlinearity shown in Fig. 2. Judging from the plot for 600 kV electrons in Fig. 2, the linearity is valid at least down to lnT = −0.4, which corresponds to the electron transmittance of about 2/3. This information should be beneficial in practical tomography experiments because one can foresee quality of the reconstruction from the minimum transmittance in a single BF-TEM image prior to the tilt series acquisition.

References
[1] S. Motojima et al., Appl. Phys. Lett. 56 (1990) 321.
[2] J. Yamasaki et al., submitted to Microscopy.
[3] N. Tanaka et al., Microscopy 62 (2013) 205.


The authors are grateful to Mr. M. Ohsaki in JEOL Ltd. for discussions about designing the sample holder and System In Frontier Inc. for discussions on precise 3D reconstruction procedures. We also thank Mr. Y. Yamamoto and Dr. C. Morita of HVEM laboratory in Nagoya University for their assistance with the experiments. One of the authors (N.T.) thanks Dr. S. Motojima for useful discussions.

Fig. 1: Carbon microcoils. (a) SEM image (Microphase Co., Ltd.) and (b) BF-TEM image taken by the HVEM.

Fig. 2: Attenuation of electron transmittance T in BF-TEM images with increasing thickness.

Fig. 3: 3D reconstructions of the CMC from the tilt series of BF-TEM images taken at (a) 400 kV and (b) 1000 kV.

Type of presentation: Poster

IT-10-P-2565 Using transmission electron tomography to unravel the structure of hybrid active layers in non volatile memory elements

Girleanu M.1,2, Nau S.3, Sax S.3, List-Kratochvil E.3,4, Soliwoda K.5, Celichowski G.5, Grobelny J.5, Brinkmann M.1, Ersen O.2
1Institut Charles Sadron, UPR-22 CNRS, 23 rue du Loess, BP 84047, 67034 Strasbourg cedex 2, France, 2Institut de Physique et Chimie des Matériaux de Strasbourg, UMR 7504 CNRS-UdS, 23 rue du Loess BP43, 67034 Strasbourg cedex 2, France, 3NanoTecCenter Weiz Forschungsgesellschaft mbH, Franz-Pichler-Straße 32, A-8160 Weiz, Austria, 4Institute of Solid State Physics, Graz University of Technology, Petersgasse 16, 8010 Graz, Austria, 5Department of Materials Technology and Chemistry, Faculty of Chemistry, University of Lodz, Pomorska 163, 90-236 Lodz, Poland
girleanu@ipcms.unistra.fr

The interest in plastic electronics has grown significantly over the last twenty years with the development of new electronic devices, particularly solar cells and field effect transistors. More recently, non-volatile memories (NVMEM) based on hybrid materials gained particular interest [1,2]. Typically such devices consist of an insulating or semi-conducting layer sandwiched between a bottom ITO electrode and an upper metallic electrode (silver or aluminum). Bias application on these devices switches them from a low current OFF state to an ON state at high current. The switching mechanism involves the formation of "conducting filaments" [3] in the active layer but no direct observation of such filaments in a hybrid device was demonstrated so far. In hybrid layers, the formation of filaments is eased by the presence of metallic nanoparticles. It is therefore important to understand how such nanoparticles are dispersed in a polymeric matrix. In this study, we have analyzed the morphology of hybrid layers of polystyrene loaded with functionalized Au nanoparticles (NPs). TEM tomography allows to determine the 3D distribution of the nanoparticles in the devices. Evidence is found for the formation of large clusters of Au NPs that are rejected to the top surface of the active layer during the film spin-coating. Moreover, a fraction of Au NPs is dispersed in the PS film but only within an interfacial layer close to the bottom ITO substrate whereas the largest part of the PS matrix does not contain any Au NPs. It is shown that TEM tomography is a valuable tool to unravel the 3D structure of the active layers in NVMEM.

[1] J.C. Scott et al., Adv. Mater. 2007, 19, 1452.

[2] B. Cho et al., Adv. Funct. Mater. 2011, 21, 2806.

[3] S. Nau et al., Adv. Mater. 2014, DOI: 10.1002/adma.201305369.


Fig. 1: Cross-sectional view of the active layer of a hybrid non volatile memory element obtained by tomography. The active layer is made of a polystyrene matrix loaded with Au nanoparticles.

Type of presentation: Poster

IT-10-P-2603 Analysis of bainitic transformation process in Cu-Al-Mn Alloy by using an orthogonally arranged FIB-SEM for precise 3D microstructure analysis

Motomura S.1, Hara T.2, Nishida M.3, Omori T.4, Kainuma R.5, Asahata T.6, Fujii T.7
1Interdisciplinary Graduate School of Engineering Science, Kyushu University, Kasuga, Fukuoka 8168580, Japan, 2Advanced Key Technology Division, National Institute for Materials Science, 1-1 Namiki, Tsukuba, Ibaraki 3050044, Japan, 3Faculty of Engineering Science, Department of Engineering Science for Elec-tronics and Materials, Kyushu University, Kasuga, Fukuoka 8168580, Japan, 4Department of Metallugy, Materiale Science, and Materials Processing, Grad-uate School of Engineering, Tohoku University, Aoba-yama 02, Aoba-ku, Sen-dai 9808579, Japan, 5Department of Metallugy, Materiale Science, and Materials Processing, Graduate School of Engineering, Tohoku University, Aoba-yama 02, Aoba-ku, Sen-dai 9808579, Japan, 6Hitachi High-Tech Science, Corp. 36-1 Takenoshita, Oyama-cho, Sunto-gun, Shizuoka 4101393, Japan, 7Hitachi High-Tech Science, Corp. 36-1 Takenoshita, Oyama-cho, Sunto-gun, Shizuoka 4101393, Japan
MOTOMURA.Shunichi@nims.go.jp

In order to investigate a 3D microstructure of complex materials precisely, we have developed an orthogonally-arranged FIB-SEM instrument which is specially designed to obtain a high-quality serial sectioning SEM image-set. The most characteristic point of this instrument is that the SEM and the FIB are arranged orthogonally. Fig. 1 shows the concept of this instrument. The advantages of this arrangement are that high-resolution and high-contrast SEM images can be obtained with low accelerating voltage such as less than 1kV because of the uniform background intensity and the short working distance (2mm). Furthermore, since the analytical instruments (EDS, EBSD and STEM, etc. ) can be located ideally , multiscale analyses can be performed in the single instrument. Fig. 2 shows the arrangement of apparatuses around a specimen viewed from the top along the SEM axis. We applied this technique on the analysis of the microstructure of bainite phase in non-ferrous noble metal based alloys. Bainitic transformation has both characteristics of a diffusionless and a diffusional transformation. Many studies on bainitic transformation have been conducted in various alloy systems such as steel, a noble metal (Cu, Ag, and Au) based alloy systems. However, the mechanism of the bainitic transformation has still been unclear. In this study, in order to reinvestigate the bainite in Cu-Al-Mn alloy, several samples with varying aging condition are prepared and observed by the orthogonally-arranged FIB-SEM and other recent SEM and TEM techniques. Fig. 3 show the SEM secondary electron (SE) image of the sample aged at 503K for 10 min. We can see some capillary binite precipitations. As a result of serial sectioning, however, it was revealed that the shape of bainite is plate-like crystal. The results of the analysis of the transformation process with these new techniques will be discussed.


Fig. 1: Schematic illustration showing the configuration of the SEM and FIB in the orthogonally arranged system.

Fig. 2: Schematic illustration showing the arrangement of apparatuses around a specimen.

Fig. 3: SEM SE-image of Cu-Al-Mn alloy aged at 503K for 10 min sample.

Type of presentation: Poster

IT-10-P-2620 Tune Stem-ADF/HAADF conditions improving dislocations tomography

Ibarra A.1, Cuestas C.1, San Juan J.2, Nó M. L.2, Arnaudas J. I.1,3
1Laboratorio de Microscopías Avanzadas, Instituto de Nanociencia de Aragón, Universidad de Zaragoza, 50018 Zaragoza, Spain , 2Dptos. Física Aplicada II y Física Materia Condensada, Fac. Ciencia y Tecnología, Universidad del Pais Vasco Apdo. 644, 48080 Bilbao, Spain, 3Dpto. Física de la Materia Condensada, Universidad de Zaragoza, 50009 Zaragoza, Spain
aibarra@unizar.es

The 3D geometry of dislocations and their interactions govern the properties of many materials. Mechanical and electrical properties are controlled by the density, distribution and dynamics of defects at the micrometer level: dominant slip mechanisms, Lomer-Cottrell locks of dislocations… Transmission electron microscopy (TEM) enables image individual dislocations and understands the relationship between the defect structure of materials and their macroscopic properties. Conventional electron micrographs are, however, two-dimensional (2D) projections of three-dimensional (3D) structures. A solution to this problem is electron tomography, which has mainly been used for reconstructing the shapes of samples using mass-thickness contrast [1]. However diffraction contrast has been considered too complex to be used in tomography due to the change of the contrast when tilt the sample.
Recently, dislocation tomography was attempted using weak-beam dark-field (WBDF) contrast [2]. A successful reconstruction was produced but the process was needlessly difficult. An alternative approach was tried by using annular dark field (ADF) STEM image [3], which is less sensitive to small specimen misalignments, making it more suited to materials research. In addition, ADF STEM images are less dynamical than WBDF images and consequently other strong features such as bend contours, thickness fringes and striped dislocation contrast affect to ADF-STEM image in a minor way than in WBDF condition.
The aim of this work is to tune the microscope conditions in order to enhance the contrast and resolution of the image avoiding some artifacts. The first results have been carried out in CuAlNi Shape Memory Alloys (SMA) because of the interest of dislocations as nucleation points for the martensitic transformation and consequently on the thermo-mechanical properties of the alloys. We can observe in Fig. 1 dislocations loops parallel to the sample surface (001), carried out by STEM-HAADF technique adjusting the conditions in the microscope to improve the contrast when tilts the sample. A tomogram of the observed dislocations has been taken each degree between +60 and -60 and reconstructed by Simultaneous Iterative Reconstruction Technique (SIRT). A picture of the final reconstruction is showed in Fig. 2. Large edge dislocations with [001] vector line and mixed ones with [101] direction have been analyzed.

References:
[1] Weyland M and Midgley P 2003 Ultramicroscopy 96 413
[2] Barnard J S, Sharp J, Tong J R and Midgley P A SCIENCE VOL 313 21 JULY 2006
[3] Sharp J, Barnard J S, Kaneko K, Higashida K and Midgley P A Journal of Physics: Conference Series 126 (2008) 012013


Acknowledgments
This work has been supported by Gobierno de Aragón, Grants E81 and Fondo Social Europeo. Authors thank Spanish Ministry of Economy and Competitivity, MICINN projects MAT2012-36421 and Consolider-Ingenio CSD2009-00013.

Fig. 1: Fig. 1: STEM-HAADF image, the conditions in the microscope have been adjusted to enhance the contrast and avoid some artifacts. We can observe a prismatic loop dislocation, lines [100] and [010] with b = (1/2)[001] in plane (001) and dislocations out of plane.

Fig. 2: Fig. 2: Snapshot of the 3-D reconstruction. We can observe different steps in the prismatic loop and check the out-of-plane dislocations. Edge dislocations with vector line [001] and mixed dislocations with [101] direction.

Type of presentation: Poster

IT-10-P-2637 Combined 3D characterization of porous zeolites by STEM and FIB tomography

Beltrán A. M.1, Przybilla T.1, Winter B.1, Knoke I.1, Machoke A.2, Schwieger W.2, Spiecker E.1
1Center for Nanoanalysis and Electron Microscopy (CENEM), Department of Materials Science and Engineering, Friedrich-Alexander-Universität Erlangen-Nürnberg, Cauerstrasse 6, 91058 Erlangen, Germany, 2Lehrstuhl für Chemische Reaktionstechnik, Friedrich-Alexander-Universität Erlangen-Nürnberg, Egerlandstrasse 3, 91058 Erlangen, Germany
Thomas.Przybilla@ww.uni-erlangen.de

Conventional transmission electron microscopy (TEM) techniques are usually limited to acquire information of specimen in two dimensions. In many cases, a three-dimensional (3D) characterization is required, as is the case for porous materials used in catalysis, for which a detailed knowledge of the 3D-morphology, size distribution and interconnectivity of the pores is crucial.
In this work we compare the 3D characterization of micro- and macroporous zeolite particles used as catalyst support (Fig. 1) with the aim of obtaining a detailed analysis of its porous structure by two different and complementary techniques, namely electron tomography (ET) based on annular dark-field (ADF) scanning TEM (STEM) and focused ion beam (FIB) tomography. The size of the particles is in the range of several micrometres, so both techniques are applicable with their specific advantages and disadvantages. ADF-STEM ET has been performed in a FEI Titan3 80-300 microscope at 200 kV, with a spatial resolution of 2.1 nm and a convergence semi-angle of 5 mrad for an increased depth of focus. Tilt series have been acquired in a tilt range from -72° to 72° (1.5° tilt increments). For the 3D reconstruction, the simultaneous iterative reconstruction technique (SIRT) algorithm has been applied with 50 iterations. ADF-STEM ET has revealed a porous structure (Fig. 2a) with mostly interconnected pores. However, due to the missing wedge, artefacts in the reconstruction can be observed due to the lack of information for the highest tilt angles (Fig. 2b).
In order to reconstruct larger volumes FIB tomography is a well suited alternative. In this case, a FEI Helios NanoLab 660 Dual Beam FIB is used to perform the sequential milling and imaging with a depth resolution of 20 nm by using a milling voltage of 5 kV and beam currents in the range of 10 - 40 pA. Fig. 3 exemplarily shows the SEM image of one slice recorded during a FIB tomography series. The pores are nicely resolved. However, in this case the well-known curtaining effect leads to artificial striations in the shadow of pores.
The combination of both tomography techniques is well suited for a more complete 3D characterization of such medium-sized structures. Further work is focusing on the combined application of 360° ET (full tilt-angle range) and (subsequent) FIB tomography to one and the same particle and a detailed comparison of the reconstructed volumes. The application of 360° ET prevents missing wedge artefacts and, therefore, improves the quality of the reconstruction.


Financial support from the German Research Foundation through the Priority Program 1570 and the Cluster of Excellence EXC 315 “Engineering of Advanced Materials”.

Fig. 1: a) SEM and b) STEM images of the studied zeolites.

Fig. 2: a) Reconstruction (surface rendering) and b) ortho slice view of the reconstruction showing the artefacts due to missing wedge and interconnections between pores of a zeolite particle by ADF-STEM ET.

Fig. 3: SEM image (tilted view) showing one single slice of the zeolite particle shown in Figure 1a) during FIB tomography series (milling was performed from top to bottom). Please note that milling artefacts occur due to curtaining below the pores (indicated by arrows).

Type of presentation: Poster

IT-10-P-2720 Prospects of Electron Holographic Tomography at Atomic Resolution : Linear Reconstruction of Dynamic Scattering using Simulated Tilt Series

Krehl J.1, Lubk A.1, Lichte H.1
1TU Dresden, Dresden, Germany
Jonas.Krehl@triebenberg.de

Electron Holographic Tomography (EHT) has been shown to be a powerful tool for directly measuring electric potentials at medium resolution in three dimensions (3D). Although Electron Holography enables the retrieval of waves also at atomic resolution, it has not yet been possible to reconstruct the 3D information from a specimen with atomic resolution. That is because tomographic reconstruction schemes generally assume a linear transfer of specimen information (e.g. the potential) into the recorded signal. At medium resolution and orientation out of zone axis this is valid for the phase of electron waves (in the Phase Grating Approximation). However, as dynamic scattering becomes dominant at atomic resolution this linear approximation becomes invalid. The aim of this work was to examine and clarify the artefacts and errors that are created by this disparity of applying standard tomographic reconstruction methods. The obtained results are important for the development of tomographic schemes suited for atomic resolution.

To that end tilt series of a single gold nanocrystal were simulated and subsequently reconstructed; the thusly acquired electric potential is analysed and compared to the original specimen potential. Care has been taken to avoid low-index zone axes during tilt. Furthermore, special attention was paid to the influence of regularisation on the reconstruction. In order to characterise the reconstruction quality several generic defects have been simulated apart from a pure crystal: a lattice vacancy, a substitute atom and a shifted atom.

The results show that the neglection of Fresnel diffraction of the wave, while transmitting through the specimen, in standard tomography is the dominating artefact in the reconstructed potential. It leads to broadening of the reconstructed atomic potentials and a characteristic dip at their centre (see Fig. 1); both artifacts depend on the distance to the focal planes of the individual waves of the tilt series. Apart from that, the reconstructed atomic potential information is well located around the original atomic position, as shown by the well-localized effect of the atom removal / substitution in the reconstruction (see Fig. 2).

Consequently, linear reconstruction schemes are not disqualified per se at the atomic level: If they could be augmented to include the Fresnel Propagation they may become a viable method for the reconstruction of experimental data. However, the numerous problems of the experimental acquisition of atomic resolution tilt series will prove to be additional hurdles on the path to 3D atomic resolution.


This work is funded by the European Union (ERDF) and the Free State of Saxony via the ESF project 100087859 ENano.

Fig. 1: A stripe of a cross section through the reconstructed potential of a gold monocrystal (diameter about 4 nm). The simulated projections, used in the reconstruction, were at 1° intervals from -90° to 89° and for each orientation the focus was set to the object exit plane.

Fig. 2: Equivalent areas of a cross section through the reconstructed potentials of gold monocrystals. (A) exclusively made up of gold atoms, whereas (B) one atom replaced with a silver. Their difference (C) = (A-B) shows the influence of the different Z of Gold and Silver.

Type of presentation: Poster

IT-10-P-2985 Autofocus method with high-definition TV camera for ultrahigh voltage electron microscope tomography

Nishi R.1, Kanaji A.1, Yoshida K.1, Kajimura N.1, Nishida T.1, Isakozawa S.2
1Osaka University, Osaka Japan, 2Hitachi High-Technologies Corporation, Ibaraki, Japan
rnishi@uhvem.osaka-u.ac.jp

The 3 MV ultra-high voltage electron microscope (UHVEM) H-3000 at Osaka University has capability of observation for micrometers' thick-sliced biological samples. This fea¬ture is suitable for tomographic three-dimensional imaging [1]. While taking obtain a tilt series of electron tomography, acquiring a hundred images, their image position and focus must be accurately aligned automatically. We proposed the Auto-Focus system using image Sharpness (AFS) [2] is suitable for acquisition of UHVEM tomography series [3]. The method is that values of image sharpness corresponding to defocus values become to be maximized as shown in Fig.1. To find the maximized image sharpness, we use fitting five points with a different defocus value to quasi-Gaussian function [3]. Acquisition of images by the slow scan CCD (SS-CCD) camera is good image quality but the acquisition time is taken more than one minute for one autofocus operation getting five defocused images.
In this study, we use a high-definition TV camera (HDTV camera; effective image area is 1.2k × 1k size) instead of the SS-CCD camera (4k × 4k) for fast acquisition of images. The HDTV camera captures one image for only 1/30 second. However, S/N of the image and the resolution are lower than the SS-CCD camera. To improve poor S/N, we integrated the images for 22 frames so that each image sharpness is enough to fitting. For lower resolution than SS-CCD image, we selected the defocus step made to be larger to discriminate difference of sharpness with each defocused image. By using HDTV camera for autofocus process, it took 6 seconds during one autofocus procedure, which became shorter by one order. It took 30 seconds to record one image by SS-CCD after autofocus and position alignment. We can obtain the series of 61 images during 30 minute. So, we successfully decreased total acquisition time of tomography series in half.

[1] A. Takaoka, et al, Ultramicroscopy 108 (2008) 230-238.
[2] H. Inada, et al, Proc. of 8APEM (2004) Kanazawa, Japan, 60-61.
[3] R. Nishi, et al, Microscopy 62(5)(2013) 515–519.


This work was supported by the Ministry of Education, Culture, Sports, Science and Technology (MEXT), Japan, under a Grant-in-Aid for Scientific Research (Grant No. 23560024, 23560786).

Fig. 1: Image sharpness S(x) to relative object lens current. The unit of the abscissa axis is manually adjusted minimum focusing step. Red solid circles are measured image sharpness and solid curve is fitted curve with the inset equation.

Fig. 2: Relative objective lens current change with a tilt angle during acquisition of tomography series. Starting angle is -60 degree and end angle is +48 degree. Two series were acquired in the same area. The deviation was smaller than the minimum step by manual.

Type of presentation: Poster

IT-10-P-3102 Compressed-sensing EDX Tomography of Composite Nanowires

Yeoh C. S.1, Saghi Z.1, Midgley P. A.1
1Department of Materials Science and Metallurgy, University of Cambridge, UK
csmy2@cam.ac.uk

The introduction of high solid angle EDX detectors has seen a renaissance in interest in EDX-based electron tomography [1]. The confined geometry of the TEM however makes the EDX spectrum prone to artefacts especially for samples tilted to high angles or mounted on copper grids [2]. Here we examine these issues further by analysing an EDX spectrum-image tilt series of a composite nanowire structure.

The sample, a Cobalt phthalocyanine (Co Pc) zinc oxide core-shell nanowire, has been studied using an FEI Tecnai Osiris equipped with a large solid angle (>0.9 srad) silicon drift detector (SDD). The ZnO nanowire shell is polycrystalline and the level of detail reconstructed in the tomogram should allow comparative assessments of spatial resolution and uniformity of grey levels in the reconstruction.

A series of EDX spectrum-images together with STEM HAADF images was recorded at 5° tilt increments with equal acquisition time (Fig 1a). The total counts in each spectrum-image for each element of interest are plotted as a function of tilt in Fig 1b. As has been seen previously the counts at high angles increase dramatically. At low angles there is a small drop in counts up to around ±20°; this is likely to be attributable to geometric shadowing of the EDX detector by the specimen holder.

The trend in the count increase is typified by the Cu signal which we believe arises primarily through the excitation (via scattered electrons) of x-rays originating from the copper support grid. As the grid is tilted the area of copper in line of sight of the scattered electrons increases by a simple geometric factor equal to approximately 1/cos(tilt angle), resulting in an increase in Cu signal. This function is plotted in Fig 1b and the fit to the Cu signal is good.

We decided to use the Zn signal in the nanowire as a test case; the increase in the signal mirrors to a large extent that seen in the Cu (and Co). As a first approximation, in order to use this tilt data for a reconstruction, we normalised the Zn signal at each tilt increment, given each map had the same acquisition time. We used compressed sensing (CS) methods [3] to reconstruct the Zn tomogram and compared the result with a more conventional SIRT reconstruction (see Fig 2). Two advantages of CS reconstruction are apparent: i) the morphology, seen in the cross-sectional slice, is more faithfully reproduced (c.f. STEM HAADF reconstruction) and ii) greyscales within the ZnO phase are more uniform.

Further work is underway to confirm the origins of the x-ray signal variation with tilt in order to move towards a true quantitative compositional tomogram.

[1] Möbus et al. Ultramicroscopy 2003,96(3-4),433-451
[2] Slater et al. Proceedings of EMAG Conference 2013
[3] Leary et al. Ultramicroscopy 2013,131,70-91


We thank Ana Borras, ICMS, Seville, Spain, for providing samples, Rowan Leary and Pierre Burdet for their help including CSET and EDX spectrum-image processing. We acknowledge support received from David Brown and Sasol Technology UK, and also the European Union Seventh Framework Program under Grant Agreements: 312483–ESTEEM2 (Integrated Infrastructure Initiative - I3) and 291522-3DIMAGE.

Fig. 1: (a) 0° tilt STEM HAADF image of Co Pc-ZnO core-shell nanowire mounted on 5 nm C-film, the tilt axis is vertical, (b) Total counts summed over EDX maps for Co, Cu and Zn Kα peaks for varying tilt angle with function 1/cos(tilt angle) fitted to Cu distribution. Acquired with a probe current of 0.7 nA, 45x45 pixels and dwell time of 40 ms.

Fig. 2: (a) and (b): Slices through tomographic reconstructions of the Zn Kα peak (integrated over 8.49 – 8.79 keV): (a) 30 iterations SIRT reconstruction performed in Inspect3D, (b) CSET reconstruction. (c) Slice through tomographic reconstruction of STEM HAADF tilt series using CSET.

Type of presentation: Poster

IT-10-P-3208 Electron tomography in the scanning electron microscope

Ferroni M.1, Migliori A.2, Morandi V.2, Ortolani L.2, Pezza A.2, Sberveglieri G.1
1Department of Information Engineering, University of Brescia and CNR-INO, Via Valotti 9, 25123 Brescia - Italy, 2CNR-IMM Section of Bologna, via Gobetti 101, 40129 Bologna, Italy
matteo.ferroni@unibs.it

The achievements in the implementation of electron tomography in the scanning electron microscope (SEM) and the potential of this 3-D imaging technique are summarized and discussed.
In SEM, the 3D imaging strategies consist in slice-and view assisted by FIB or microtomy, for the investigation of large specimen volumes. Differently the best resolution is pursued for relatively small volumes through TEM at high beam voltage.
The proposed implementation of electron tomography in the SEM is appropriate to the investigation at nanometric resolution of specimen volumes in the intermediate range, namely 2000 (w) x 2000 (l) x 200 (thickness) nm, as it combines the reconstruction algorithm with the signal corresponding to incoherently scattered electrons in the Scanning-Transmission (STEM) imaging mode. STEM imaging takes advantage from some peculiar characteristics of the experimental set-up [3]. This approach attains nanometric resolution and is free from aberrations caused by post-specimen imaging lenses; it also allows to collect transmitted electrons over a wide angular range [4][5]. The optimization of detector design and performance makes the contrast comply with local variations of composition or projected thickness. The bright-field component of the transmitted electrons can be effectively separated from the dark-field one, by varing the detection strategy [4].
The STEM mode preserves the monotonic variation of the signal with specimen thickness and meets the basic projection requirement for the 3-D analysis of nanowires, carbon based nanostructures or ultrastructures of biological specimens. In addition, the large value for the maximum detection angle ensures a complete detection of the scattered electrons, even in case of relatively large specimen thickness. In the case of tomography, these features are essential to maintain the proper image contrast when the specimen is rotated.
Fig. 1 shows the reconstruction of a ZnO crystalline nanostructure from a 110° tilt series at 1° step. ImageJ [6] with the TomoJ plug-in was used [7]. The disposition of the wires, their uniform section and the tapered termination are properly retrieved. Similarly, carbon-based tubes, filled with cobalt nanoclusters, were reconstructed as shown in Fig. 2. The tomogram from the STEM tilt series featuring compositional contrast, clearly shows the cobalt clusters inside the tubes.
These results demonstrate the potential of the method and optimization of the experimental set-up is under development to consolidate this technique in the set of 3-D methods of electron microscopy.

REFERENCES

1 Merli et al. Ultramic. 88 (2001) 139.
2 Morandi et al , JAP 101 (2007) 114917.
3 Morandi et al. APL 90 (2007) 163113.
4 http://rsbweb.nih.gov/ij/
5 Messaoudii et al, BMC Bioinf. 8 (2007) 288.


The authors acknowledge the finacial support from TomoSEM (F97I12000120007)

Fig. 1: Up-Left) SEM image of ZnO nanowires. The ROI is boxed. Up-Right) TEM shows the regular hexagonal shape and the pyramidal termination of the ZnO single-crystal nanowires. Bottom) Visualization of the reconstructed volume (4.5 micron large) corresponding to a primary magnification of 50.000.

Fig. 2: (Left and Center) STEM Bright- and Dark- field compositional images of carbon tubes filled with Co nanoclusters - Beam energy 30 keV. – (Right) Tomogram of the 
cobalt clusters inside the carbon tubes. The inset shows one Co particle, demonstrating the chemical sensitivity and the monotonically variation of the contrast upon tilting.

Type of presentation: Poster

IT-10-P-3212 Phase Contrast Cryo-Electron Tomography with a New Phase Plate

Khoshouei M.1, Danev R.1, Gerisch G.2, Ecke M.2, Plitzko J.1, Baumeister W.1
1Max Planck Institute of Biochemistry, Department of Molecular Structural Biology, Martinsried, 82152, Germany. , 2Max Planck Institute of Biochemistry, Department of Cell Dynamics, Martinsried, 82152, Germany.
maryamkh@biochem.mpg.de

There has been more than 60 years working on the concept of combining phase contrast method with electron microscopy. Many laboratories are involved in development of the phase contrast electron microscopy to improve the performance of transmission electron microscopes and to reduce the beam damage to frozen-hydrated biological specimens [1].

There are different types of phase plates e.g. thin film, electrostatic, Photonic, magnetic and anamorphotic phase plates. Each method has its own pros and cons but overall the most successful phase plate has been thin carbon film Zenrike phase plate. Thin carbon film Zenrike phase plate has its own drawback such as charging effect and having a short lifetime [2]. Recently, we developed a new type of the phase plate, which has much longer lifetime and in general better performance.

The aim of the current work is development and applications of the phase contrast method. The work is being carried out at Max-Planck Institute of Biochemistry in Germany in collaboration with FEI in the Netherlands.

The structure of the whole vitrified worm sperm from Lumbricus terrestris species has been studied using new generation of phase plate. Sperm motility is a critical factor for fertilization and highly depends on the ultrastructure of different parts of the mature sperm [3].

This large and filiform mature cell comprises of acrosome (Fig.1a), nucleus (Fig.1b), mitochondria (Fig.1c) and flagellum (Fig.1d). In the pertinent literature, some research has taken place based on ultrastructure of worm sperm using tissue fixation or plastic sectioning but not in cryo. This research is carried out in cryo in combination with phase contrast method.


References:

[1] R. Danev, S. Kanamaru, M. Marko and K. Nagayama, Journal of Structural Biology 171 (2010), p. 174-181.
[2] R. Danev and K. Nagayama, Journal of Structural Biology 161 (2010), p. 211-218.
[3] A. Rolando et al, International Journal of Morphology (2007), p. 277-284.


We would like to thank Julia Mahamid for her assistance in Cryo-electron tomography.

Fig. 1: Slices of tomograms from Earth worm sperm with the new generation of phase plate. Acrosome (a), Nucleus (b), Mitochondria (c) and Flagellum (d) [Titan Krios 300kV, energy filter, direct detector, def: -500nm, mag:26000]                                                                                                                              

Type of presentation: Poster

IT-10-P-3225 Dedicated and innovative system for tomography in the Scanning Electron Microscope

Morandi V.1, Migliori A.1, Ortolani L.1, Pezza A.1, Maccagnani P.1, Masini L.1, Ferroni M.2, Sberveglieri G.2, Rossi M.3, Vittori-Antisari M.4, Vinciguerra P.5, Pallocca G.5, Del Marro M.5
1CNR-IMM Bologna Section, Bologna, Italy, 2Dept. of Information Engineering, Brescia University, Brescia, Italy, 3Dept. of Basic and Applied Sciences for Engineering, Sapienza University, Roma, Italy, 4Unità Tecnica Tecnologie dei Materiali, ENEA Casaccia, Roma, Italy, 5Assing S.P.A., Roma, Italy
morandi@bo.imm.cnr.it

For 3D non-destructive materials characterization, two are the leading tomography techniques: X-Ray Computed Tomography and Electron Tomography implemented in TEM operated in STEM mode [1]. The first one is undoubtedly the most important for industrial applications, providing resolution of few tens of μm, for cm- to mm-scale objects, while, the second one is of great interest in many research fields, and is capable of 3D reconstruction of sub-μm-scale objects with a resolution up to 0.24 nm [2].
In this paper we will highlight the implementation and the capabilities of an alternative electron tomography system in a SEM operated in STEM mode, composed by dedicated sample holder, STEM detector and analogue/digital signal processing system.
This system aims to cover the range between the previously mentioned techniques, taking advantages of the flexibility of the SEM platform, and of the resolution and image quality capabilities of the STEM mode implemented [3]. We will show that the STEM-in-SEM tomography approach opens up the perspective for the 3D analysis of volumes up to 100 μm3, such as nanowires, carbon based nanostructures or biological specimens, with resolution up to 10 nm.
Fig. 1 shows the main building blocks of the of the innovative tomographic acquisition system. The principal constraint for the success of tomography using STEM imaging is to maintain the monotonic variation of the mass-thickness contrast over the whole tilt range. Therefore, the use of STEM for tomography requires an acquisition system capable of collecting the transmitted electron over a large and adaptable collection angles. The detector geometry with five independent circular active sectors permits to optimize the efficiencies of signal collection, as a function of the tuneable specimen-detector distance, energy and beam current. The dedicated specimen holder with rotation capability ensures the eucentricity of the observed detail and a reduced missing wedge, which are fundamental for 3D reconstruction purposes (Fig. 1a). Moreover, the dedicated signal processing system (Fig. 1b) performs faster than conventional STEM systems in the acquisition of the tilt series, preserving the amplification, conditioning and managing of the signals as required by tomographic reconstruction.
Finally, we will demonstrate the potential of the method with examples of 3D reconstruction of micro and nano-structures. In Fig. 2 is reported the tomographic reconstruction and the 3-D rendering of a bundle of human skin collagen fibers, as obtained with the first release of our dedicated system [4].

References

1. P.A. Midgley et al. J. of Ultram. 223 (2006) 185
2. M.C. Scott et al. Nature 483 (2012) 444
3. V. Morandi et al. App. Phys. Lett. 90 (2007) 163113
4. TomoSEM Project: F97I12000120007


Fig. 1: (a) Setup inside the SEM-chamber: specimen-holder on motorized stage and STEM detector on a specific holder with variable working distance capability. (b) SEM platform and external mixed signal boards for amplification and managing of the signals.

Fig. 2: Acquisition steps for a complete tomogram of a collagen fibers bundle. (a) STEM image of a human skin thin section. (b) Tilt series acquisition. (c) Reconstructed tomogram with highlighted the ROI shown in the 3-D rendering in (d).

Type of presentation: Poster

IT-10-P-3301 Improved Electron Tomography Image Reconstruction usingCompressed Sensing based Adaptive Dictionaries.

AlAfeef A.1, Cockshott W. P.1, MacLaren I.2, McVitie S.2
1School of Computing Science, University of Glasgow, Glasgow G12 8QQ, UK, 2SUPA School of Physics and Astronomy, University of Glasgow, Glasgow G12 8QQ, UK
a.al-afeef.1@research.gla.ac.uk

Electron tomography (ET) is an important technique for studying the 3D morphologies of nanostructures using the electron microscope. ET involves the collection of a series of 2D projections over a wide tilt range, which are subsequently aligned and processed to obtain a 3D volume reconstruction. It is well known that the quality of the reconstruction obtained using established algorithms is significantly affected by artifacts when the maximum angular range (the “missing wedge” artefact) or the number of acquired projections is limited. The reconstruction quality can be enhanced by including additional prior knowledge about the specimen in the reconstruction process and this is the key point of the compressive sensing (CS) family of techniques[1]. Such approaches have recently been applied to ET[2] and showed excellent results with higher fidelity and reduced artefacts even with subsampled datasets. Such features give CS major advantages for ET such as reducing total irradiation dose. The key prior knowledge employed in CS is that the signal (i.e. images), needs to be sparse in a transform domain. If a suitable transform enables a sparse representation of the dataset, then the original signal can be accurately reconstructed from a significantly smaller set of measurements than that required by the classical sampling theorem. As sparsity is a key requirement for an accurate reconstruction, researchers have investigated a range of sparsifying transforms, including for ET. In spite of their success in some cases, such transforms may not apply for all cases (nanostructured objects), and real signals are not always compressible (sparse) in such transforms. Also, some sparsifying transforms have a limited ability to remove artifacts. One common sparsifying transform is Total Variation (TV). It is only effective for those samples that are well described in terms of sharp, discrete boundaries. Other drawbacks of using TV include over-smoothing of fine structures and the inability to separate true structures from noise. Consequently, it is essential to seek superior transforms. In this work, we propose an alternative image reconstruction algorithm for ET that learns the sparsifying transform adaptively (in a similar manner to how our visual cortex processes natural images[3]). This new technique ET data with higher fidelity than analytically based CS reconstruction algorithms. The proposed technique is tested using a simulated phantom, which is known to be difficult to reconstruct using the popular CS-TV techniques, together with an experimental tilt series from a polymer solar cell.

References
[1] Donoho, D.2006. IEEE T Inform Theory, 52 1289–1306.
[2] Saghi, Z. et al.2011 Nano Lett. 11 4666–4673
[3] Olshausen, A.et al. 1996 Nature. 381 607–609


This research was supported by a Lord Kelvin Adam Smith Scholarship of the University of Glasgow.

Fig. 1: Reconstructed images from simulated under-sampled tilt series. A) CS-Phantom- The reference image of the numerical simulation. B) Reconstruction using TV based approach (CSTV) and C) proposed dictionary learning based approach (DLET) from noisy 28 projections.

Fig. 2: Quality curves of different reconstruction experiments using WBP, CSTV and DLET with higher degree increment steps between projections. Used metrics are: Peak Signal-to-Noise Ratio (PSNR), visual signal-to-noise ratio (VSNR) and Entropy Correlation Coefficient (ECC).

Fig. 3: Adaptively learned dictionary consisting of 100 atoms of 55 patches used in the DLET.

Type of presentation: Poster

IT-10-P-3455 Use your smartphone to calibrate your TEM's goniometer

Wollgarten M.1, Stapel H.1, Garcia-Moreno F.1
1Helmholtz Zentrum Berlin für Materialien und Energie GmbH, Hahn-Meitner-Platz 1, 14109 Berlin, Germany
wollgarten@helmholtz-berlin.de

For reconstruction of tomographic data sets the precise knowledge of the experimental parameters is mandatory. Besides the incident intensity [1], the tilt angles have to be known precisely. In a recent paper, Hayashida and co-workers [2] measured the accuracy of a TEM goniometer and found a total deviation of 4° over an angular range of about 180°. This shows that the goniometer might be a source of flawed input data with serious consequences for the reconstruction work. Thus, calibrating the TEM's goniometer can be essential for high quality tomography work.

For calibration, a digital protractor can be used[2]. However, smartphones provide a number of sensors, among them accelerometers.

To determine the tilting accuracy of our LIBRA 200 FE TEM the acceleration sensors were used. An android application was written which reads out the acceleration raw data along the SP´s x-, y-, z-axis and stores it to a file.

While orienting the SP such that two axes show zero acceleration, the third axis is expected to be parallel to the earth´s gravity field vector. In these positions accelerations different from 9.81 m/s² where found. As a first approach, we slowly rotated the device to find for each axis the maximum and minimum acceleration value and used this pair to linearly scale the reading to the interval [-1,1]. In a second step, the three component acceleration vector was normalized to length 1.0 g (= 9.81 m/s²).

To measure the tilting angles of the TEM, the SP was mounted on a Fischione (model 2040) tomography holder. A self written script within Gatan´s Digital Micrograph was used to tilt the holder from -76° to 76° with 1° increments. Each tilting step was followed by a 10 seconds rest.

The acquired data set was processed as outlined above resulting in a stepped curve (tilt versus time).

The difference between the measured and the set value is plotted in Fig. 1.

An error with a periodicity of about 15° of the goniometer's worm gear is evident. However, whereas the reproducibility within each group is very good, an offset is observed between the groups which we attribute to a insufficient calibration of the accelerometers.

Nevertheless the sensitivity of the sensors turns out to be enough to detect tilting deviations significantly smaller than 0.1°. Therefore, we aim at improving the sensor calibration. This will allow for precise calibration of the goniometer.

References
[1] Wollgarten, M., Habeck, M., Micron, in print, doi: 10.1016/j.micron.2014.02.005 (2014).
[2] Hayashida, M., Terauchi, S., Fujimoto, T., Rev. Sci. Instrum. 82, 103706, doi: 10.1063/1.3650457 (2011).


Fig. 1: Measured deviation from the nominal goniometer tilt angle. The squared markers represent the configuration for which the smartphone was vertically mounted at a nominal tilt of zero degree. For the other set it was fixed in horizontal orientation.

Type of presentation: Poster

IT-10-P-5857 3D imaging of Si FinFETs by combined HAADF-STEM and EDS tomography

Qiu Y.1, 2, Van Marcke P.1, Richard O.1, Bender H.1, Wilfried V.1, 2
1Imec, Kapeldreef 75, B-3001 Leuven, Belgium , 2Instituut Kern-en Stralings Fysika, K.U.Leuven, B-3001, Leuven, Belgium
Yang.Qiu@imec.be

In last few decades electron tomography is intensively studied to recover the 3D shape of a wide range of nano-scale materials. High angle annular dark-field scanning transmission electron microscopy (HAADF-STEM) tomography attracts also increasing attention to study the morphology of semiconductor nano-devices with dimensions below 10 nm scale and a truly 3D morphology as instead of 2D projection images it provides 3D reconstruction with sensitivity along the beam direction. Recent advances of TEM systems with high brightness field emission gun (FEG) coupled with high count rate X-ray energy dispersive spectrometers (XEDS) opens the possibility to analyze the 3D volume both in imaging and chemical analysis modes.
Samples with dense Si fins (45 nm pitch) after Si etch, oxide fill and recess and thin epi-Si growth are explored. Pillar shaped specimens with various diameters are deposited on top of standard TEM grids and analyzed with Fischione conventional tomography holder. The work is done without gold markers, cf. fig 1, in a Titan cube 60-300 double aberration corrected system with SuperX EDS detector. Compared to lamella samples, the pillar configuration allows to increase the maximum tilt angle range from ~±65º to ±80º so that missing wedge effects are minimized. HAADF-STEM images are acquired in 1º steps and EDS maps each 5º. The STEM images are aligned and reconstructed by Inspect3D. Fig 2 presents the volume rendered 3D visualizations of the HAADF-STEM reconstruction and the orthoslices indicated in the volume. Si, SiO2 and FIB-damaged Si can be easily differentiated by the intensity. Moreover, Si and O X-ray maps are reconstructed separately with the same alignment. Fig 3 shows the superposition of both Si and O reconstructed volumes and the same corresponding slices as shown in fig 2. The orthoslices from the EDS reconstruction agree well to the HAADF-STEM reconstruction and also reveals the oxide grown on the Si fins. Compared to the 2D projection images, the orthoslices from the reconstruction indicate that the Si etch depth varies as indicated by blue arrows in fig 2c, 2d and fig 3c, 3d.
We show that the alignment can be done without marker tracking. The reconstructed volume based on HAADF images and EDS maps can bring useful fully interpretable information in composition and morphology unlike the 2D images that suffer from projection effects. Aberration corrected TEM improves the spatial resolution. Further analysis will involve using 360º tilt holder to fully eliminate the missing wedge artifacts and quantification of the elemental reconstructed volume, as well as application to next FinFET device processing steps involving gate and metallization steps.


Fig. 1: Fig. 1 a) HAADF-STEM images of the pillar shaped specimen (diameter around Si fins area ≈ 240nm) on standard copper grid zoomed in b) and c)

Fig. 2: Fig. 2 a) Volume rendered 3D visualization of HAADF-STEM reconstruction and its corresponding orthoslices in xy plane (b), yz plane (c) and xz plane (d) respectively

Fig. 3: Fig. 3 a) 3D chemical rendering of the Si fins and its corresponding orthoslices in xy plane (b), yz plane (c) and xz plane (d) respectively (gray : Si, yellow/red : SiO2)

Type of presentation: Poster

IT-10-P-6047 Multi-Axis Electron-Holographic Tomography

Sturm S.1, Wolf D.1, Lubk A.1, Lichte H.1
1Triebenberg Laboratory, Institute of Structure Physics, Technische Universität Dresden, Germany
Sebastian.Sturm@Triebenberg.de

Electron holography has proven to be a suitable method for measuring electrostatic and magnetic fields of nanostructures in the TEM. In electron holographic tomography (EHT) [1], the intrinsic fields can be mapped in all three dimensions, thus providing quantitative access to the true inner potentials and not only their projections. However, single-axis (SA) electron tomography (ET) often suffers from loss of information in one dimension, due to the limited tilt range of common tomographic specimen holders. In Fourier space, this limitation leads to an unsampled area, the so-called “missing wedge”. Performing multi-axis ET by combining several SA tilt series of the same object, one can minimize this missing volume. In case of Dual-Axis Tomography (DA) for example (two series acquired with tilt-axes perpendicular to each other) the volume can be reduced to a “missing pyramid” [2]. Fig 1 shows the corresponding missing volumes in Fourier space, depending on the available tilt range. For usual tilt ranges of about +-70°, DA is expected to provide a much better resolution in the third dimension than SA.
Another potential benefit of Multi-Axis EHT is the possibillity to reconstruct more then only one component of the B-vector field of magnetic samples.
Here we report on the development and implementation of Multi-Axis EHT.
For the necessary alignment of the residual displacements within the tilt series, the centre of mass for the projected potential at each tilt has been determined [3]. This method has two major advantages: all tilt series are inherently aligned with respect to each other, and the (projected) tilt axis is automaticaly alligned.
For tomographic reconstruction, a self-written software program has been developed. It is based on the concept of a weighted simultaneous iterative reconstruction technique (WSIRT [4]) but is especially adapted to the peculiarities of multi-axis geometries (Fig 2). Combining all tilt series within the process of a 3D back-projection, before application of a weighting filter in 3D-Fourier space and applying the iteration loop, takes into account the projections of all tilt series at once, instead of applying the weighting and iteration loop for each 2D-slice independently with separate SA tilt series.
As an example in Fig 3 the 3D potential of barium titanate nanoparticles [5], reconstructed by DA EHT, is compared with the two corresponding 3D potentials, reconstructed from only one of the two SA tilt series of the complete dataset. In z-direction the DA tomogram exhibits sharper edges then its SA counterparts.

[1] D Wolf et al. Ultramicroscopy, 110(5) (2010), 390-399

[2] P Penczek et al. Ultramicroscopy, 60(3) (1995), 393-410

[3] S Sturm. Diploma thesis, TU Dresden (2011)

[4] D Wolf. Dissertation, TU Dresden (2010)


[5] BTO nanoparticles provided by D. Szwarcman and G. Markovich (Tel-Aviv University).

[6] Funded by the EU (ERDF) and the Free State of Saxony via the ESF project 100087859 ENano.

Fig. 1: Fig. 1: Missing wedge (a) and missing pyramid (b) in Fourier space and information loss according to the corresponding volumes.

Fig. 2: Fig. 2: 3D-WSIRT algorithm reconstructing several tilt series synchronously.

Fig. 3: Fig. 3: Dual Axis EHT reconstruction of BaTiO3 nanoparticles. The measured object edges in the tomogram are compared with single axis reconstructions of the two data subsets.

IT-11. Electron holography and lens-less imaging

Type of presentation: Invited

IT-11-IN-1935 Atomic Resolution Electron Diffractive Imaging and 3D

Zuo J. M.1, Lyu X. W.1, Gao W. P.1, Meng Y. F.1
1Department of Materials Science and Engineering, University of Illinois at Urbana-Champaign, Urbana, IL, 61801, USA
jianzuo@illinois.edu

Electron diffractive imaging promises sub-angstrom resolution imaging in 3D. Key to electron diffraction imaging is coherent electron diffraction using a parallel beam for selected area diffraction. Lateral coherent length as large as ~500 nm in FEG TEM has been reported [1].
The principle of phase retrieval is based on finding solutions based on a set of constraints. One of the constraints is the object support. The phase retrieval is carried out iteratively. For electron diffractive imaging, use of phases recorded in electron images at the beginning of iteration helps with a number of experimental issues[2]. Sub-Å resolution imaging has been demonstrated for a number of materials, including carbon nanotubes[3], CdS[2], CeO2 [4], Si [5] and TiO2 [6]. Accurate phase retrieval at nm resolution was recently demonstrated by Yamasaki et al [9]. Electron diffractive imaging can also be easily extended to medium and low energy electrons [7, 8]. Dronyak and his co-workers experimentally determined the morphology of a single MgO nanocrystal using the measure 3D Bragg diffraction peak [10]. 3D reconstruction resolving atoms was reported by Chen et al. [12] using experimental STEM image data [13].
Here we report a new method of 3D reconstruction using Fienup’s hybrid input-output (HIO) algorithm. Electron diffraction patterns are centered in 3D reciprocal space in a single axis tilt series. For the 3D sample data, the computational cost increase dramatically in order to achieve higher resolution result. We overcome this challenge by GPU-acceleration. Simulation.3D structure of Au icosahedron is reconstructed from calculated diffraction patterns including missing angles and noise in order to test the algorithm performance. Experimental implementation and its challenge will be discussed.

Reference
[1] S. Morishita, J. Yamasaki, N. Tanaka, Ultramicroscopy 129, 10-17 (2013)
[2] W. J. Huang, J. M. Zuo et al., Nature Physics 5, 129-133 (2009).
[3] J.M. Zuo, J. Zhang, W.J. Huang, K. Ran, and B. Jiang, Ultramicroscopy 111, 817-823 (2011).
[4] A. J. Morgan et al., Phys. Rev. B 87, 094115 (2013)
[5] S. Morishita et al. Applied Physics Letters 93(18), 183103 (2008)
[6] L. De Caro et al., Nature Nanotechnology 5, 360-365 (2010)
[7] O. Kamimura et al., Ultramicroscopy 110(2), 130 (2010)
[8] T. Latychevskaia et al., (2103), arxiv.org/pdf/1305.1897
[9] Yamasaki, J et al, Appl. Phys. Lett, 101, 234105 (2012)
[10] Dronyak R et al., Appl. Phys. Lett., 96 , 221907 (2010)
[11] Chen, C.-C. et al., Nature 496, 74 (2013)
[12] Rez, P. & Treacy, M. M. J. Nature 503, http://dx.doi.org/10.1038/nature12660 (2013)
[13] J. Miao et al., Nature 503, E1–E2 doi:10.1038/nature12661, (2013)
[14] This work is supported by DOE BES DE-FG02-01ER45923


This work is supported by DOE BES DE-FG02-01ER45923.

Fig. 1: Figure 1, left, a schematic illustration of sampling in reciprocal space as used in a single axis tilt series of electron diffraction patterns. Right, reconstructed 3D image using GPU accelerated HIO algorithm from simulated diffraction patterns of an icosahedron nanoparticle with 1.2 Å information limit and 25º missing wedge and simulated noise.

Type of presentation: Invited

IT-11-IN-5757 Towards atomic resolution and three-dimensional mapping of electrostatic potentials and magnetic fields using off-axis electron holography

Dunin-Borkowski R. E.1
1Ernst Ruska-Centre for Microscopy and Spectroscopy with Electrons and Peter Grünberg Institute, Forschungszentrum Jülich, D-52425 Jülich, Germany
rafaldb@gmail.com

The ability to achieve high phase sensitivity with close-to-atomic spatial resolution in off-axis electron holographic measurements is offered by the latest generation of ultra-stable transmission electron microscopes, which are equipped with high brightness electron sources and aberration correctors. In this talk, I will discuss recent developments in the quantitative and three-dimensional characterization of electrostatic potentials and magnetic fields with close-to-atomic spatial resolution using electron holography. I will begin by describing two complementary approaches that can be used to measure the electrostatic potentials and electric fields of electrically-biased metal needles as a function of applied voltage in the electron microscope. The phase shift can be analyzed either by fitting the recorded phase distribution to a simulation based on lines of uniform charge density or by using a model-independent approach involving contour integration of the phase gradient to determine the charge enclosed within the integration contour. Both approaches typically require evaluation of the difference between phase images recorded at two applied voltages, in order to subtract mean inner potential (and magnetic) contributions to the phase. I will then describe recent progress in the development of a model-based approach that can be used to reconstruct the three-dimensional magnetization distribution in a specimen from a series of phase images recorded using electron holography. The approach involves the projection of three-dimensional magnetization distributions onto two-dimensional grids to simulate phase images of three-dimensional objects from any projection direction. This forward simulation approach is then used in an iterative model-based algorithm to solve the inverse problem of reconstructing the three-dimensional magnetization distribution in the specimen from a tilt series of phase images. Such a model-based approach avoids many of the artifacts that result from the use of classical tomographic techniques. Finally, I will consider challenges associated with the use of chromatic aberration correction of the Lorentz lens in the TEM to achieve higher spatial resolution in magnetic characterization. When considering experiments aimed at the retrieval of weak phase shifts, it is important to remember that the sample must remain clean, that electron-beam-induced charging can contribute to the measured phase shift and that the quantitative interpretation of phase shifts measured from crystalline specimens can require comparisons with dynamical simulations. If time permits, I will conclude with recent progress in the application of off-axis electron holography to obtain results during ultrafast switching processes in situ in the electron microscope.


M. Beleggia, T. Kasama, V. Migunov, J. Caron, J. Ungermann, A. Kovacs, A.H. Tavabi, P. Diehle, A. London, T.F. Kelly, D.J. Larson and M. Farle are thanked for their valuable contributions to this work.

Type of presentation: Oral

IT-11-O-1464 Split-illumination electron holography

Tanigaki T.1, Aizawa S.1, Park H. S.1, Matsuda T.2, Harada K.3, Murakami Y.1,4, Shindo D.1,4
1Center for Emergent Matter Science (CEMS), RIKEN, Saitama, Japan, 2Japan Science and Technology Agency, Saitama, Japan , 3Central Research Laboratory, Hitachi Ltd., Saitama, Japan, 4Institute of Multidisciplinary Research for Advanced Materials, Tohoku University, Sendai, Japan
tanigaki-toshiaki@riken.jp

  Off-axis electron holography [1] has been used for observing microscopic distributions of magnetic fields, electrostatic potentials and strains at nanoscale level and for aberration-corrected electron microscopy by detecting phase shifts of electron waves. An off-axis electron hologram is formed by overlapping an object wave transmitted through a sample with a reference wave passed through the reference area. The inherent problem with this method is that the distance D between the object and reference waves, or the hologram width W, is limited by the lateral coherence length R or by the brightness of the illuminating electron waves.

 

  We solved this long-standing problem by developing split-illumination electron holography (SIEH). Experiments were performed using a 300-kV cold field emission transmission electron microscope (TEM) (HF-3300X, Hitachi High-Technologies Co.).

 

  In our SIEH (Fig. 1), we can illuminate a sample by using two highly separated and yet coherent electron waves without reducing the density of electron and form high-contrast holograms at regions far from the sample edge. The separation distance D can be controlled by a condenser biprism in the illuminating system. The fringe spacing s and the width W of the hologram can be independently controlled as in double-biprism electron interferometry [2]. Using SIEH, a fringe contrast of 50% can be attained even if the object wave is as far as 17 μm from the reference wave in the sample plane [3].

 

  Recently, in order to improve precision of phase measurement in SIEH, double condenser biprism type SIEH without Fresnel fringes was developed (Fig. 2) [4]. Since demanded phase shifts to be measured in nanoscale are becoming smaller and smaller, it is important to improve precision of phase measurements to broaden the applications of the off-axis electron holography. The developed methods are used for varieties of applications and will be used for revealing electromagnetic phenomena in atomic scale.

 

References:

[1] A. Tonomura, “Electron holography”, (Springer-Verlag, 1999).

[2] K. Harada et al., Appl. Phys. Lett. 84 (2004) 3229.

[3] T. Tanigaki et al., Appl. Phys. Lett. 101 (2012) 043101.

[4] T. Tanigaki et al., Ultramicroscopy 137 (2014) 7.


The authors are grateful to the late A. Tonomura for his valuable discussions. This research was supported by a grant from JSPS through the “FIRST Program”, initiated by CSTP.

Fig. 1: Schematic diagrams of electron-optical method and fringe contrasts of holograms. (a) Conventional electron holography. (b) Split-illumination electron holography in which a coherent electron wave is split into two coherent partial waves. (c) Measured fringe contrasts C of holograms as function of distance D between object and reference waves.

Fig. 2: Schematic diagram of double condenser biprism (CB) type split-illumination electron holography without Fresnel fringes (a) and holograms (b, c) and phase images (d, e) of charged latex particles. (b, d): double CB, (c, e): single CB.

Type of presentation: Oral

IT-11-O-1669 Electrostatic potential of single-layer graphene measured using electron holography and ab-initio calculations

Chang S.1, Dwyer C.1, Nicholls R.2, Boothroyd C. B.1, Bangert U.3, Dunin-Borkowski R. E.1
1Ernst Ruska-Centrum, Forschungszentrum Juelich, Germany 1, 2Department of Materials, University of Oxford, UK 2, 3Department of Physics and Energy, University of Limerick, Ireland 3
shery.chang@fz-juelich.de

Graphene, a single-layer, hexagonally-coordinated carbon material has attracted huge attention in a wide range of fields due to its unique structural properties [1]. For example, it has found applications in electronic devices, energy storage, and electrocatalysis [2]. Characterisation of graphene imposes a requirement for high sensitivity to image a thickness of one atom. High-resolution TEM and ADF-STEM have been used to study the atomic arrangements and defects in graphene (and its related materials). The electrostatic potential of a single layer of graphene, a fundamental quantity of a materials structure property, is however less explored.

Here we use electron holography and density functional theory calculations to accurately measure the electrostatic potential of a single-layer of graphene. A Cs and Cc aberration-corrected TEM (Pico), operated at 80kV, was used to take holograms of graphene. The biprism voltage was set to be 175V, giving interference fringes of spacing 0.04nm. The graphene was grown using chemical vapour deposition on a SiO2 substrate and then transferred onto a TEM grid.

Figure 1 shows the phase of a typical area of the graphene sheet. It can be seen that there is a band near the edge of the graphene and some patches across the graphene sheet with larger phase shifts, which are typical features of silicon oxide and other hydrocarbon contamination left on graphene from TEM specimen preparation. The edges of the graphene sheet are more than a one-layer thick although patches of single-layer graphene can be found.

Figure 2 shows a region of single-layer graphene near the edge, after 2-hours of electron beam illumination to form a hole for the reference wave. The modulus of the Fourier transform of the complex wave-function (shown in the inset of figure 2) shows that high spatial resolution information is present in the phase (with the 1-210 reflection visible). The phase shift from the single layer graphene was measured to be 58 mrad (with respect to the vacuum) and the phase profile is shown in the inset of figure 2.

In order to compare the experimental measurement of the electrostatic potential with theory, both all-electron (Wien2K) and density functional theory calculations (VASP) were used. The theoretical calculation gives good agreement with the experimental measurement. Further implications from the theoretical calculations will be discussed in the presentation.

References:
[1] K. S. Novoselov et. al., Science, 2004, 306, 666 [2] Y. Sun et. al., Energy Environ. Sci., 2011, 4, 1113


Fig. 1: Phase of the holographam of a typical region of a graphene sheet.

Fig. 2: Phase of the hologram of the single-layer graphene. Inset (right) shows the modulus of the fft of the complex wave, and the inset (left) shows the phase profile from the vaccum to the graphene region. 

Type of presentation: Oral

IT-11-O-1798 Low-voltage electron diffraction microscopy of multi-layer graphene

Kamimura O.1, Dobashi T.1, Maehara Y.2, Kitaura R.3, Shinohara H.3, Gohara K.2
1Central Research Laboratory, Hitachi, Ltd., 2Division of Applied Physics, Faculty of Engineering, Hokkaido University, 3Department of Chemistry & Institute for Advanced Research, Nagoya University
osamu.kamimura.ae@hitachi.com

The full text of the abstract is not available. Please contact the presenting author.


Type of presentation: Oral

IT-11-O-1905 Electron Holographic Tomography of Mean Free Path Lengths with nm-Resolution

Lubk A.1, Wolf D.1, Röder F.1, Lichte H.1
1Triebenberg Laboratory, TU Dresden, Dresden, Germany
Axel.Lubk@triebenberg.de

In off-axis Electron Holography a Möllenstedt biprism is introduced slightly above an intermediate image plane (image coordinates R) in the TEM to generate an interference pattern (“hologram”) with the following sinusoidal intensity distribution Ihol(R) = I0+I(R)+2μA0A(R)cos(QR+φ(R)). Here, I0, I, μ, A0, A, Q and φ are the reference intensity, conventional image intensity, contrast damping factor due to camera MTF and partial coherence, reference amplitude, reconstructed amplitude, carrier frequency and reconstructed phase. Reconstructed phases have been successfully analyzed in terms of various electrostatic and magnetostatic potential characteristics. In this contribution we will show how to tomographically reconstruct elastic and inelastic mean free path lengths (MFPL) from the concomitantly reconstructed conventional image intensity I and amplitude A. Starting points are the following exponential attenuation laws ln(A(R)/A0) = 0.5∫1/λA(r)dz and ln(I(R)/I0) = ∫1/λI(r)dz with corresponding attenuation coefficients λ-1 holding under out-of-zone axis conditions employed in medium resolution Electron Holography. Based on fundamental electron scattering principles we relate the λs to elastic and inelastic MFPLs correcting a serious misinterpretation preventing quantitative analysis in the past. Noting that the attenuation laws represent a Radon transformation when performed over a π-range of tilt angles, we then develop adapted tomographic reconstruction schemes. That involves dedicated normalization and regularization in order to reduce the influence of the generally low SNR. We demonstrate the MFPL reconstruction at a GaAs-Al1/3Ga2/3As core shell nanowire grown by low pressure metal-organic vapor phase epitaxy (MOVPE) method using colloidal Au nanoparticles (NPs) as metal catalysts. The tilt series was recorded at a Cs-corrected FEI TITAN TEM at 300 kV. Amongst other features its reconstructed potential reveals the core-shell structure as well a potential slope of yet unknown origin towards the Au tip (Fig. 1). The reconstructed elastic MFPL data (Fig. 1) also reveals the core-shell structure as well as a chemical composition variation from AlAs to GaAs in the tapered region thereby facilitating an unambiguous interpretation of the above noted potential slope in terms of chemical composition change. The inelastic MFPL (Fig. 1) on the other hand only vaguely hinds the existence of the core-shell which is a result of the strong delocalization of the dominating bulk plasmon excitation in the inelastic MFPL. Both elastic and inelastic MFPL agree very well with theoretic predictions.


We thank N. Lovergine (University of Salento, Lecce) for providing the GaAs/AlGaAs nanowire and T. Niermann for helping with recording the tilt series. The authors acknowledge financial support from the European Union under the Seventh Framework Program under a contract for an Integrated Infrastructure Initiative. Reference 312483 - ESTEEM2.

Fig. 1: Short compendium of 3D data of GaAs-AlGaAs core-shell nanowire including reconstructed potential and both, elastic and inelastic, mean-free-path-length (MFPL) cross-sections with corresponding linescans.

Type of presentation: Oral

IT-11-O-2018 Mapping the number of graphenes for whole micron-size flakes by mean of low voltage electron holography

Castro C.1, Ortolani L.2, Arenal R.3,4, Monthioux M.1, Masseboeuf A.1
1CEMES, University of Toulouse, Toulouse, France, 2CNR IMM-Bologna, Bologna, Italy, 3Laboratorio de Microscopias Avanzadas, Instituto de Nanociencia de Aragon (INA), Universidad de Zaragoza, Calle Mariano Esquillor, Zaragoza, Spain, 4Fundacion ARAID, Zaragoza, Spain
celia.castro@univ-rouen.fr

Synthesis of graphene with controlled properties is a utopian goal to reach if a multiscale analyzing method is not developed. One of the challenges is to allow accurate counting of monoatomic layers at the nanoscale over a flake area. Therefore, the number of layers as well as their stacking configuration have been related to optical and electrical properties of few-layer graphene.1
The recent developments of aberration-corrected transmission electron microscopes (AC-TEM) working at low-voltage (LV) conditions, which limit the knock-on damage, make possible to obtain atomic-resolution information on carbon-based materials.2-3
Counting edges of graphene stacks or peeling them under the electron beam provide very local information and cannot be applied to thick stacks. Quantitative thickness mapping can be obtained by combining high angle annular dark field imaging (HAADF) and electron diffraction. HAADF intensity is thickness-dependent and electron diffraction provides a calibration tool by determining the signal of a monolayer related to the TEM settings.4
Another way for mapping the number of graphene layers is LV transmission electron holography. The phase shift of electrons induced by the surface electrostatic potential is proportional to the thickness. This phase shift is intrinsic to the mean inner potential of the individual graphene layer and directly represents the local number of layers.5
In the present study, this method is emphasised in the I2TEM machine, a new AC-TEM dedicated to electron holography developed by Hitachi with CEMES Lab. We take advantage of much larger holograms, free of Fresnel fringes keeping irradiation damages limited and still achieving nanometer scale resolution thanks to the unique combination of a double biprism configuration, a second stage unit located upper in the column (Lorentz mode), and the LV with cold field emission. First maps of quantized graphene layers over micronic field of view will be presented. A variety of graphene flakes obtained from CVD or exfoliated graphitewill be analyzed through parameters including stacking type, sample preparation and artifacts of carbon based materials on TEM. Two examples taken from graphite flakes are provided as figure 1 and 2, in which the number of graphenes is large. The method is however sensitive enough for mapping thickness variations related to single graphene.
1 Koshino M. New J. Phys. 15 2013 15010
2 Sasaki T., et al. J. Electr. Microsc. 59 2010 S7
3 Suenaga K. et al. Nature, 468 2010 1088
4 Meyer J.C. et al. Solid State Comm., 1–2 2007 10
5 Ortolani L. et al. Carbon, 49 2011 1423


The authors acknowledge for financial support the EU-7Framework Program 312483-ESTEEM2, the "Conseil Regional Midi-Pyrénées" and the European FEDER within the CPER program, the Transpyrenean Associated Laboratory for Electron Microscopy (TALEM) and the French National Research Agency for the ANR-10-EQPX-38-01 and the ANR GRAAL. 

Fig. 1: a. Electron hologram of a multi-graphene and folded flake b. Phase contour map every 10 graphene layers. Reported values represent local measurements.

Fig. 2: a. Phase contour map every 10 graphene layers. Reported values represent local measurements. b. Profile of number of layers extracted from figure 2a.

Type of presentation: Oral

IT-11-O-2326 Quantitative comparison of experimental and calculated image waves at atomic resolution

Niermann T.1, Lehmann M.1
1Institut für Optik und Atomare Physik, Technische Universität Berlin, Straße des 17. Juni 135, 10623 Berlin, Germany
niermann@physik.tu-berlin.de

Nowadays, the image wave function within the transmission electron microscope can be experimentally obtained by off-axis electron holography with high quality. This becomes possible by recent progresses made in instrumentation as-well-as in the reconstruction of holograms [1].

By matching experimentally obtained wave functions with specialized empirical models, several structural information, like atomic column positions, can be obtained with high precision, e.g. [2]. Instead of using such empirical models, we report here on quantitative comparisons of the full reconstructed wave function with the full image wave functions as calculated by the Bloch-wave method and the normally applied formulation of wave propagation by the objective lens. The matching between experiment and simulation was done by least-square fitting, i.e. using the complex ℓ2-norm pixelwise.

As experimental (nuisance) parameters the origin of the lattice, the absolute amplitude, the global phase as well as a linear phase change over the field of view were optimized. Beside the global phase, these parameters should, in principle, already be experimentally determined by empty holograms that were taken as reference. However, a linear phase change easily happens due to slight charging of the specimen. Furthermore, the global amplitude might change due to drifts of the illumination.

Optimized specimen parameters are thickness (including a linear change of thickness over the field of view), specimen tilt, and strain (linear change of column distance over the field of view). Considered imaging parameters were focal spread, two-fold astigmatism, axial coma, and defocus as-well-as a linear change of defocus over the field of view (corresponding to an inclined exit surface of the specimen).

Fig. 1 shows an experimentally obtained wave function from a GaAs-wedge recorded in [110] zone axis and the corresponding matched calculation. The left hand sides of Fig. 1 shows the comparison with aberrations applied to calculations. The right hand side shows the same comparison with the experimental wave function corrected by these aberrations, which exhibits the familiar dumbbell contrasts of GaAs. Furthermore, we report on the properties of the minimum and investigate the variances and correlations of the parameters (Fig. 2 and 3), which manifest in the shape of the minimum, and investigate specimen regions of different thickness.

[1] Niermann & Lehmann, Micron (2014), DOI: 10.1016/j.micron.2014.01.008
[2] S. Bals et al., PRL 96 (2006) 096101


Support by the DFG within SFB787 is kindly acknowledged.

Fig. 1: Comparison between experimental wave function and calculation. The upper and lower rows shows the wave functions in amplitude and phase, respectively. In the left four panels, the residual lens aberrations are applied on the model, in the four panels on the right hand side, the experimental wave is corrected for the lens aberrations.

Fig. 2: Evaluation of mismatch in dependence of specimen tilt. The tilt is expressed by the coordinates of the center of the Laue circle, zone axis is [110]. The orientation of the reciprocal space is indicated in Fig. 1. The minimum is close to (0.5,-0.5,1.5). (Here only a constant thickness/defocus over the field was assumed).

Fig. 3: The shape of the minimum of error function in dependence of defocus and thickness shows the slight correlation between both parameters. The fit is more sensitive to thickness than defocus, since the minimum is more steeper in the former case. (Here only a constant thickness/defocus over the field were assumed).

Type of presentation: Oral

IT-11-O-2383 Coherent imaging beyond detector area and Abbe limit, towards atomic resolution

Latychevskaia T.1, Fink H-W1
1Physics Institute, University of Zurich, Winterthurerstrasse 190, 8057 Zurich, Switzerland
tatiana@physik.uzh.ch

In a typical imaging experiment, data analysis relies on the recorded data during the experiment. In coherent imaging, this could be a hologram or a diffraction pattern obtained with light, electrons, X-rays, or any other type of radiation with wave nature. The achievable resolution is determined by the numerical aperture of the experimental setup limited by the size of the detector area.
We present a method that allows extrapolation of an experimental record beyond the area detected during the experiment by using intrinsic wave properties, namely their continuity in space. The amplitudes of the scattered waves are mapped onto the detector area and allow retrieval of the phase distribution. Once the complex-valued distribution of the scattered waves are retrieved, we extrapolate them to the full space extend, far beyond the detector area. As a result, the object reconstructed from such extrapolated interference pattern exhibit a higher resolution than provided by the initial experimental record [1-2]. The most attractive feature of our technique is that it does not require a new experiment as it can be applied to an already existing experimental record. The interference pattern can numerically be post-extrapolated to the full 2π hemi-sphere leaving the wavelength as the only resolution limiting factor.
An example is shown in Fig. 1, where a small section of a noisy interference pattern created by two point sources displaying less than three interference fringes is extrapolated to a much larger interference pattern. As a result, the two point sources can be resolved in the reconstruction. Fig. 2-3 show application of the technique to experimental optical in-line holograms and diffraction patterns.
The post-extrapolation technique is especially interesting when applied to electron or X-ray interference patterns as it can reveal atomic resolution from low-resolution images. Moreover, the extrapolation can also be applied to crystalline structures where diffraction patterns exhibit distinct Bragg peaks, such as graphene, see Fig.4. Diffraction patterns of graphene [3], could be extrapolated to reveal higher-order Bragg peaks and achieve enhanced resolution in the reconstruction, as shown in Fig.4.
We will present the application of this extrapolation method towards holograms and diffraction patterns of both, non-crystalline and crystalline structures, demonstrating its application for different types of waves: electrons, X-rays and THz waves, and we will also address the possibility of three-dimensional extrapolation.

1. Latychevskaia, T. and H.-W. Fink, Applied Physics Letters, 2013. 103(20): p. 204105.
2. Latychevskaia, T. and H.-W. Fink, Optics Express, 2013. 21(6): p. 7726-7733
3. Longchamp, J.-N., et al., Phys. Rev. Lett., 2013. 110(25): p. 255501.


The work presented here is financially supported by the Swiss National Science Foundation (SNF).

Fig. 1: Fig.1. Extrapolation of an interference pattern. (a) Fraction of an interference pattern created by two point-scatterers. (b) Extrapolated interference pattern. (c) and (d): Image of two point sources reconstructed from (c) the fraction of the diffraction pattern and (d) from the extrapolated interference pattern [1].

Fig. 2: Fig.2. Extrapolation of an in-line hologram. (a) Scanning electron micrograph of the sample and (b) its experimental optical hologram. (c) Extrapolated hologram. (d) and (e): Object reconstruction from hologram (b) and the extrapolated hologram (c), respectively. The insets shows the intensity profiles [2].

Fig. 3: Fig.3. Extrapolation of a diffraction pattern. (a) Scanning electron micrograph of the sample. (b) Piece of its optical diffraction pattern. (c) Extrapolated diffraction pattern. Reconstructions of the sample obtained from (d) the piece of diffraction pattern and (e), the extrapolated diffraction pattern [1].

Fig. 4: Fig.4. Extrapolation of a simulated diffraction pattern of a crystalline structure. (a) Graphene patch containing 1003 carbon atoms with two divacancies and (b) its diffraction pattern. (c) Extrapolated diffraction pattern. (d) and (e): Reconstructions of the sample obtained from (b) and (c), respectively.

Type of presentation: Oral

IT-11-O-2391 Electron exit wave reconstruction in Gaussian basis: from high resolution image to diffraction pattern

Borisenko K. B.1, Allen C. S.1, Warner J. H.1, Kirkland A. I.1
1Department of Materials, University of Oxford, Parks Road, Oxford, OX1 3PH, UK
konstantin.borisenko@materials.ox.ac.uk

Accurate electron exit wave reconstruction can offer not only increased resolution and improved signal to noise ratio but it can also provide some 3D information about the sample in high resolution transmission electron microscopy (HRTEM). Restoration of the exit wave from experimental data usually involves collecting a series of images with varying focus (a focal series). The exit wave can be also recovered by recording overlapping diffraction patterns (DP) and applying specially designed iterative numeric algorithms in a ptychographic approach. Both these approaches require a number of images or diffraction patterns to be collected, which increases the total electron dose that the sample needs to withstand. Such an approach can be difficult to implement for radiation sensitive materials, due to possible sample damage during acquisition of focal or diffraction series.

In the present work we suggest and test an approach that in principle allows reconstructing exit wave from a single image or diffraction pattern for a weak-phase object. One of the obstacles to using only one image or diffraction pattern for the reconstruction is that in this case the number of variables in the exit wave to be determined is comparable to the number of data recorded. Presence of the aberrations of the objective lens and experimental noise further complicates the analysis and can lead to multiple or unstable solutions. By representing the electron exit wave in Gaussian basis we greatly reduce the number of variables needed to find the solution. This approach also has an advantage of analytic representation of the DP and also derivatives needed for the reconstruction process. We test the suggested method on experimental HRTEM image of graphene and simulated diffraction pattern of a carbon nanotube. The reconstruction algorithm involves solving an overdetermined system of non-linear equations with either numeric or analytic derivatives and appears robust to noise. We compare the reconstructed experimental exit wave phase from graphene with the multislice simulations.


We thank the financial support from the European Union under the Seventh Framework Program under a contract for an Integrated Infrastructure Initiative (Ref 312483-ESTEEM2).

Fig. 1: Experimental image of graphene a) and the reconstructed exit wave phase b).

Fig. 2: Simulated input exit wave and corresponding target diffraction pattern (represented as a logarithm of the square root of the diffracted intensity) and the reconstructed exit wave.

Type of presentation: Oral

IT-11-O-2513 Observation of electric field using electron diffractive imaging

Yamasaki J.1, Ohta K.1, Sasaki H.2, Tanaka N.1
1Nagoya University, Nagoya, Japan, 2Furukawa Electric Co., Ltd., Yokohama, Japan
p47304a@nucc.cc.nagoya-u.ac.jp

    Information on electromagnetic fields in and around nanometer-sized semiconducting or magnetic devices is obtained from phase shifts of illumination electron waves. Although the most established method for phase imaging is presently off-axis electron holography, another choice could be electron diffractive imaging (EDI). In the method, a complex wave field is reconstructed from a diffraction pattern through numerical iterations under some constraints in real space. So far we have succeeded in reconstructions of atomic structures of crystals [1, 2] and thickness maps of wedge-shaped Si [3]. Figure 1 shows an example of the results, in which phase image of the transmission electron wave undergoing dynamical diffractions is reconstructed from the primary spot in a selected-area diffraction (SAD) pattern. In the present study, we performed reconstructing electric fields around MgO nano particles and a p-n junction in GaAs.
    A 200kV thermal field-emission TEM (JEOL: JEM-2100F) was used for taking SAD patterns with spatially coherent illumination. A post-column energy filter (Gatan: GIF tridium) was utilized for removing inelastic background intensity from samples and also for achieving a camera length large enough for precise sampling of low-angle scattering intensity. Energy-filtered bright field TEM images were also recorded to use as the real-space constraints. Figure 2 shows the phase reconstruction around MgO particles isolating in the vacuum. In Fig. 2(c), the electric field, which radiates from the particles positively charged by electron beam irradiation, is clearly observed. The reconstruction of the p-n junction in GaAs is shown in Fig. 3. Although the junction is invisible in the TEM image (Fig. 3(a)), the potential change deforms the primary spot (Fig. 3(b)), which results in visualization of the junction in the reconstructed phase image (Fig. 3(c)). The width of the depletion layer and the offset across the junction agree well with the doping concentration and measurements by off-axis electron holography.
    Unlike off-axis electron holography, the present method needs neither electron biprisms nor a vacuum area adjoining to the field of view of interest. The present study exhibits the future possibility that EDI will become an alternative to electron holography in some cases for observing electromagnetic fields relating to nanometer-sized materials.

References
[1] S. Morishita, et al., Appl. Phys. Lett. 93 (2008) 183103.
[2] S. Morishita, et al., AMTC Lett. 2 (2010) 116.
[3] J. Yamasaki, et al., Appl. Phys. Lett. 101 (2012) 234105.


We thank Dr. S. Morishita in JEOL Ltd. for valuable discussions. The present study was partly supported by JSPS KAKENHI (Grant No. 21760026), The Public Foundation of Chubu Science and Technology Center, and Toyoaki Scholarship Foundation.

Fig. 1: Reconstruction of the phase image of the wedge-shaped Si crystal by electron diffractive imaging. (a) Bright-field TEM image, (b) the primary spot in the SAD pattern, and (c) the reconstructed phase image.

Fig. 2: Visualization of electric field (arrows) around the charged MgO particles. (a) Bright-field TEM image, (b) the primary spot, and (c) the phase image.

Fig. 3: Visualization of the p-n junction in GaAs. (a) Bright-field TEM image, (b) the primary spot, and (c) the phase image.

Type of presentation: Oral

IT-11-O-2667 Caustics and diffraction from two oppositely biased metallic tips imaged in the coherent transmission electron microscope

Tavabi A. H.1, Migunov V.1, Dunin-Borkowski R. E.1, Pozzi G.2
1Ernst Ruska-Centre for Microscopy and Spectroscopy with Electrons and Peter Grünberg Institute, Forschungzentrum Jülich, Jülich, Germany , 2Department of Physics and Astronomy, University of Bologna, Viale B. Pichat 6/2, 40127 Bologna, Italy
a.tavabi@fz-juelich.de

The coherence of a modern field emission transmission electron microscope (TEM) allows fascinating electron-optical phenomena to be observed, such as the fine structure of umbilic foci outlining the caustic of an astigmatic probe, the hyperbolic umbilic catastrophe produced by a coma aberration function [1] and cusped fan-like structures in defocused images of electrically biased nanotube bundles [2].
Here, we study bright-field TEM images of two oppositely-biased metallic tips, which show a rich structure that depends sensitively on applied bias and defocus as a result of a combination of electrostatic field topography and electron-optical phase shift and is strongly reminiscent of the elliptic umbilic diffraction catastrophe that occurs when visible light is refracted by a water droplet with a triangular perimeter [3].
An FEI Titan 60-300 field emission gun TEM was used to study two metallic tips that had been thinned electrochemically and mounted in a specimen holder equipped with piezo-electric drivers and electrical contacts. The tips were placed in front of each other at a separation of ~1 micron and a potential difference of up to 130 V was applied between them. The positively charged wire was found to act like a terminating convergent electron biprism, producing an overlapping region of intensity containing two-beam fringes, whereas the negatively charged wire acted like a terminating divergent biprism. The combined effect of the fields resulted in a highly complex interference pattern, which is shown in Fig. 1 for a nominal defocus of 9.5 mm and a potential difference of 130 V. The overlapping region has a triangular structure that is similar to the elliptic umbilic diffraction catastrophe. Figure 2 shows this region at a higher magnification, with the hexagonal structure of the spots and their modulation by the two-beam biprism fringes visible.
We have interpreted the key features in these images by using a simple model of two uniformly and oppositely charged lines placed in front of each other [4]. The shapes of the tips can then be approximated by suitably choosing two of the equipotential ellipsoidal surfaces. The resulting simulations shown in Figs 3 and 4 are in good qualitative agreement with the experimental results.

References
[1] T. C. Petersen et. al., Phys. Rev. Lett. 110 (2013) 033901.
[2] M. Beleggia et. al., Appl. Phys. Lett. 98 (2011) 243101.
[3] M. V. Berry, J. F. Nye, and F. J. Wright, Phil. Trans. Roy. Soc. London, Series A, 291 (1979) 453.
[4] M. Muccini et .al., Ultramicroscopy 45 (1992) 77.


We are grateful to C. Dwyer for valuable discussions.

Fig. 1: Figure 1. Bright-field TEM image recorded at a defocus of 9.5 mm from two metallic needles that have a potential difference of 130 V between them.

Fig. 2: Figure 2. Central region of Fig. 1 displayed at a higher magnification, revealing spots that have a hexagonal-like structure.

Fig. 3: Figure 3. Simulated image corresponding to the experimental conditions used to acquire Fig. 1.

Fig. 4: Figure 4. Central region of Fig. 3 displayed at a higher magnification.

Type of presentation: Oral

IT-11-O-2781 Towards electron holography of 3D magnetization distributions in nanoscale materials using a model-based iterative reconstruction technique

Caron J.1, Ungermann J.2, Dunin-Borkowski R. E.1, Riese M.2
1Ernst Ruska-Centre for Microscopy and Spectroscopy with Electrons and Peter Grünberg Institute, Research Centre Jülich GmbH, Jülich, Germany, 2Institute of Energy and Climate Research – Stratosphere (IEK-7), Research Centre Jülich GmbH, Jülich, Germany
j.caron@fz-juelich.de

Electron holography is a powerful technique for recording the phase shift of a high-energy electron wave that has passed through a thin specimen in the transmission electron microscope. The phase shift is, in turn, sensitive to the magnetic field and electrostatic potential in the specimen. Here, we introduce an approach that can be used to reconstruct the three-dimensional magnetization distribution in a magnetic specimen from a series of phase images recorded using electron holography. We generate simulated magnetic induction maps by projecting the three-dimensional magnetization distribution onto a two-dimensional Cartesian grid. We use known analytical solutions for the phase shifts of simple geometrical objects to pre-compute contributions to the phase from individual parts of the grid, in order to simulate phase images of arbitrary three-dimensional objects from any projection direction, with numerical discretization performed in real space to avoid artifacts generated by discretization in Fourier space without a significant increase in computation time. This forward simulation approach is then used in an iterative model-based algorithm to solve the inverse problem of reconstructing the three-dimensional magnetization distribution in the specimen from a tomographic tilt series of two-dimensional phase images. The model-based approach avoids many of the artifacts that result from using classical tomographic techniques such as filtered back-projection, as well as allowing additional constraints and known physical laws to be incorporated.


The authors are grateful to the European Research Council for an Advanced Grant.

Fig. 1: Illustration of the simulation process: The projected two-dimensional magnetization distribution is sub-divided into pixels which are represented by simple geometries (e.g., a disc). The contribution to the phase shift from every pixel is calculated in the form of two pre-computed components, which are oriented along the axes of the grid.

Fig. 2: a) Simulated magnetic phase shift of a uniformly magnetized sphere with a radius of 64 nm in a 128nmx128nmx128nm volume. The magnetization direction is indicated by the arrow. b) Corresponding magnetic induction map (20x phase amplified). The colors represent the direction and magnitude of the phase gradient, according to the color wheel shown.

Type of presentation: Oral

IT-11-O-3089 Electrical charge quantification by electron holography

Gatel C.1, Lubk A.2, Pozzi G.3, De Knoop L.1, Snoeck E.1, Hytch M. J.1
1CEMES-CNRS and Université de Toulouse, 29 rue Jeanne Marvig, 31055 Toulouse, France, 2Institute of Structure Physics, Technische Universität Dresden, Mommsenstr. 9, 01069 Dresden, Germany, 3Department of Physics and Astronomy, University of Bologna, Viale Berti Pichat 6/2, 40127, Bologna, Italy
gatel@cemes.fr

The distribution and movement of electrical charge are fundamental to many physical phenomena, particularly for applications involving nanoparticles, nanostructures and electronic devices. However, there are very few ways of quantifying charge at the necessary length scale. Beyond providing structural and chemical information at the atomic scale, TEM can also determine the electrostatic field at the nanometer scale with a dedicated technique known as electron holography (EH).
We recently developed a new quantitative method to count the elementary charges with a precision of one elementary unit of charge using aberration-corrected EH [1]. We achieve this by applying at the nanoscale the elegance and power of Gauss’s Law to phase images extracted from holograms. This method provides direct access to the total charge enclosed by a given contour without assuming further details about neither the position of the charges within or outside the field of view nor the material investigated, contrary to a model-based approach where the whole electrostatic potential has to be computed. The extra sensitivity is achieved by the high signal-to-noise of aberration-corrected instruments and our new methodology. We performed a statistical analysis to reveal the relationship between the size of the contours and the precision of the charge measurement. A dedicated software has been developed for performing the charge evaluation based on line integration.
We will present different examples to illustrate the principle and the precision of this method. Among them, we will show the charge measurements on different MgO nanocubes where we determined a surface distribution of these charges with the corresponding value due to the surface states or adsorbates acting as charge traps (Figures 1 and 2). Another example will concern the in-situ field emission of a biased carbon cone nanotip (CCnT) [2]. The CCnT was placed to a defined distance from an Au-anode plate. We then ramped up the voltage between the nanotip and the anode from 0 to 95 V until the electric field around the tip was strong enough to allow the electrons to tunnel through the barrier and a field emission current could be acquired. During the voltage ramping and the field emission, holograms were recorded at each voltage step (Figure 3). After extracting the phase images, we applied this method to determine the numbers of accumulated charges and the charge density on different place of the tip as a function of the applied voltage (Figure 4). We will then discuss of these values, particularly the charge density at the beginning and during the field emission process.

[1] C. Gatel et al. Phys. Rev. Lett. 111, 025501 (2013)
[2] L. de Knoop et al. Micron (2014) - Accepted


The authors acknowledge the European Union under the Seventh Framework Programme under a contract for an Integrated Infrastructure Initiative Reference 312483-ESTEEM2

Fig. 1: Reconstructed phase image of a MgO nanoparticle.

Fig. 2: By contour enclosed charge as a function of the short side a as indicated on the Figure 1; the linear fit of contours within the particle and the constant fit outside of the particle are indicated by red and black lines respectively.

Fig. 3: Hologram of a biased CCnT for field emission. In white is represented the contour used to count the number of charges in the enclosed area.

Fig. 4: Enclosed charge as a function of the length of the enclosed area and the applied voltage between the CCnT and the Au-anode plate.

Type of presentation: Poster

IT-11-P-1472 Restoration of Singularities in Reconstructed Phase of Crystal Image in Electron Holography

Li W.1,3, Tanji T.2,3
1Graduate School of Engineering, Nagoya University, Nagoya, Japan, 2EcoTopia Science Institute, Nagoya University, Nagoya, Japan, 3Global Research Center for Environment and Energy Based on Nanomaterials Science, Nagoya, Japan
liwei00jp@yahoo.co.jp

Off-axis electron holography, which can be used to measure the inner potential of a specimen from its reconstructed phase image, has been widely used recently for characterizing materials. Under severe conditions such as in-situ observation in gas atmospheres, steep or large phase changes such as crystal lattice images or bulk specimen edges, and low signal-to-noise ratio conditions, abrupt reversals of contrast from white to black may sometimes occur in a digitally reconstructed phase image, resulting in inaccurate information. This phase distortion is due mainly to the digital reconstruction process and weak electron wave amplitude in some areas of the specimen. Hence, a posterior image processing that correct imperfections are indispensable for obtaining accurate phase information. In this study, we apply digital image processing to the phase image of a crystal for the restoration of such abrupt phase contrast changes, and obtain relatively accurate phase information for the crystal structure from the same electron hologram. Figure1 show the restoration of W8Nb18O69 structure phase images obtained by electron holography. The phase image (Fig.1) which is simply reconstructed and corrected the aberration of the microscope includes many singularity points as shown indicated by arrowheads. Restoring such singularity points improves the quality of reconstructed image as shown in Fig.2. Figures 3 and 4 show them in wire frame mode.Further work is required to be accomplished in the practice. The present method of phase image restoration for simulation with Poisson and Gaussian noises contributes to the correctly phase reconstruction of the hologram with quit weak electron-wave amplitude and noisy circumstance.


This study was partially supported by the Global
Research Center for Environment and Energy Based on Nanomaterials Science

Fig. 1: Directly reconstructed phase image.

Fig. 2: Phase image after restoration.

Fig. 3: Wire-frame image of Fig.1.

Fig. 4: Wire-frame image of Fig.2.

Type of presentation: Poster

IT-11-P-1514 Observation of the magnetic flux and three-dimensional structure of skyrmion lattices by electron holography

PARK H. S.1, Yu X.1, Aizawa S.1, Tanigaki T.1, Akashi T.2, Takahashi Y.2, Matsuda T.3, Kanazawa N.4, Onose Y.5, Shindo D.1, 6, Tokura Y.1, 4
1RIKEN Center for Emergent Matter Science (CEMS), Wako, Saitama, Japan, 2Central Research Laboratory, Hitachi, Ltd., Hatoyama, Saitama, Japan, 3Japan Science and Technology Agency, Saitama, Japan , 4Department of Applied Physics and Quantum-Phase Electronics Center (QPEC), University of Tokyo, Tokyo, Japan, 5Department of Basic Science, University of Tokyo, Tokyo, Japan, 6Institute of Multidisciplinary Research for Advanced Materials, Tohoku University, Sendai, Japan
hspark@riken.jp

Topological spin textures have been attracting increasing interest for use in studying quantum magneto-transport and for possible application to spintronics. Skyrmions are particularly attractive for use as information carriers in memory and logic devices because of the emergence of spin transfer torque at extremely low current densities (~106 A/m2) [1]. Several challenges must be addressed before the skyrmion can be applied to actual devices. They include realization of skyrmions at room temperature, clarification of their three-dimensional (3D) structures, and fabrication of thin films containing skyrmions. Despite recent theoretical studies, the 3D structures of skyrmions remain elusive. Observing the 3D structures of skyrmions at the microscopic level is a prerequisite for applications of skyrmions to spin-electronic devices.

Electron holography, using the wave nature of electrons, provides opportunities for directly detecting and visualizing, in real space, the phase shifts of the electron waves due to the electromagnetic fields [2]. However, precise phase measurement of weak phase objects such as skyrmions is very challenging because procedures are needed for averaging the phase images and separating the electric and magnetic vector potentials. Nevertheless, the advantage of electron holography compared to Lorentz electron microscopy and magnetic force microscopy, under just-focused condition, makes it possible to visualize a quantized magnetic flux with nanometer resolution, in addition to determining its density in the vicinity of skyrmions. Here we investigated the 2D magnetic flux distributions (Fig. 2) of skyrmion lattices in helimagnet Fe0.5Co0.5Si thin samples with a stepped thickness as shown in Fig. 1 and estimated the 3D structures of the helical and skyrmion phases by using high-voltage holography electron microscopes [3].

References:

[1] N. Nagaosa, Y. Tokura, Nat. Nanotech. 8, 899-911 (2013).

[2] A. Tonomura, Electron holography, 2nd ed., (Springer-Verlag, Tokyo, 1999).

[3] H. S. Park et al., Nat. Nanotech., in press (2014).


The authors thank the late Dr. A. Tonomura for his valuable discussions. This research was supported by the grant from the JSPS through the “FIRST Program” initiated by the CSTP.

Fig. 1: Fig. 1. Lorentz micrographs. (a) A thin sample produced by FIB technique and its illustration. Thickness differences are represented by different levels of contrast.  (b) Thickness dependence of skyrmion lattices along sample with field cooling at 25 mT and 12 K. The scale bar is 300 nm.

Fig. 2: Fig. 2. Handedness reversal of magnetic flux flow with change in direction of applied field. (a,b) Surface plots of phase image. Sign reversal of phase shift with change in applied field direction is clearly visible. (c) Enlarged surface plot in vicinity of skyrmion. Red and white arrows represent direction of lines of magnetic flux.

Type of presentation: Poster

IT-11-P-1535 Strain fields in the vicinity of SiGe nanopyramids evidenced by focus series phase retrieval imaging

Donnadieu P.1, Neisius T.2, Liu K.3,4, Aqua J. N.4, Ronda A.3, Berbezier I.3
1SIMAP CNRS, Université de Grenoble Alpes, Domaine Universitaire BP 75 38402 Saint Martin d'Hères France, 2Université Paul Cézanne - Campus St Jérôme - 13397 Marseille France, 3IM2NP – CNRS, Université Paul Cézanne 13397 Marseille France, 4INSP - UPMC - 4 place Jussieu, 75252 Paris France
patricia.donnadieu@simap.grenoble-inp.fr

Nanoscale characterization is a major issue for the development of nanostructures. Local composition and morphology are at stake as well as strain within the nanoobjects and in the near substrate since strains might be monitoring self organization. While local chemistry can be known using EDS or EELS methods, efforts are still required to determine morphology and strains by TEM. Off axis holography is able to provide such information but remains a quite dedicated technique. Hopefully among the numerous forms of holography, focus series can be used to derive phase image. The advantage of this in line holography is to be easy to carry out in the course of a classical TEM study and to provide phase image on a large scale (up to several microns). The method consists in acquiring images at different defocus (-Δ, 0, Δ) and deriving the phase image using a filtering type processing that realizes the inversion of the phase transfer function.

In the present work, the focus series method has been applied to SiGe nanostructures deposited on a Si substrate. These SiGe nanostructures are expected to have a pyramidal shape and an average composition Si20Ge80. Figure 1 displays the focus series obtained on a plane view sample.Using the image analysis described in (1), phase image with very unrealistic values were obtained while the method had been validated by the measurement of inner potential on gold nanoparticles. Actually it appears that depending on the nanoobject scale, the image analysis may require further optimization. In particular strain fields are long range effect that prevents from applying filters appropriate for a nanoparticle assembly. The filtering process has to be modified: instead of a Gaussian edge filter, a Tikhonov regularisation was applied (i.e. a q2/(q2+a2)2 filter). The phase image (Fig.2) results from this modified filtering. The nanopyramids are characterized by a central positive phase with four lobes of negative phaseshift located at the corners of the pyramids. This negative phase cannot be explained by chemical segregation since the inner potential VGe is higher than VSi but by a strain effect changing the local inner potential. However the relation between strain and phase requires more investigation.

At this point, we wish to emphasize on the phase retrieval method optimization and on the qualitative information like location of strains and the first stages of self organization, namely the striking observation of non-random islands assemblies. This behaviour is representative of preferential nucleation of islands that is now investigated with computations of both total energy and kinetic nucleation barrier as a function of the distance between islands.

(1) P. Donnadieu et al, Applied Physics Letters 94, 263116, 2009


The French CNRS-CEA network METSA is acknowledged for providing the access to TEM facilities

Fig. 1: Focus series TEM images of SiGe nanostructures deposited on Si (plane view sample prepared by chemical polishing on the Si side). Images were recorded at defocus -200 nm, 0, 200 nm. This defocus value gives the best compromise between image contrast and spatial resolution. For Δ=200 nm, the spatial resolution on the phase image is better than 5 nm.

Fig. 2: Phase image retrieved from the focus series using the modified image processing. The SiGe nanopyramids show a positive phase central part with lobes of negative phase at the corner. Remarkably the negative lobes of neighbouring pyramids are close to contact along <200> directions or overlapping along <110> directions.

Fig. 3: Sketch of a cross section of a SiGe nanopyramid deposited on Si substrate summarizing the interpretation of the phase images. The positive phase central part is consistent with a Si20Ge80 pyramid while the negative phase lobes at the pyramid corners can be interpreted by local strains in the Si substrate.

Type of presentation: Poster

IT-11-P-1545 A thin film Zernike phase plate micro fabricated using MEMS technology.

Konyuba Y.1, Iijima H.1, Abe Y.2, Suga M.1, Ohkura Y.1
1JEOL Ltd., 2Yamagata Research Institute of Technology
ykonyuub@jeol.co.jp

An image of biological and polymer samples with conventional transmission electron microscopy (TEM) shows relatively low contrast without staining because they are composed of light elements such as carbon, nitrogen and oxygen. Zernike phase contrast TEM (ZPC-TEM) is an answer to this problem [1]. So far, many types of phase plates (i.e., thin film, electrostatic, magnetic, etc.) for TEM have been proposed [2]. Above all, a thin film phase plate has produced promising results [3]. Accordingly, ZPC-TEM attracts increasing attention to Cryo-TEM applications such as cryo-electron tomography and single-particle analysis, since a specimen in the field of interest is composed of light elements and preferred to be unstained to avoid artifacts.

The thin film phase plate has however encountered a few problems. The most crucial one is charging that makes its reliability poor and its lifetime short. The second problem is that the fabrication of the phase plate has been made by human hands. For these problems we have tried to produce a phase plate through a mass production process of micro fabrication using micro electro mechanical system (MEMS) technology. At the first onset, we made a sandwich plate composed of silicon nitride (SiN) coated with metallic titanium (Ti) on both sides to reduce the charging. The entire process is shown in Fig.1. In the process the center hole, which is the most essential part for phase plate performance, was also produced by electron beam lithography and dry etching process. By virtue of these processes, the center hole is close to a perfect circle as shown in Fig. 2. And constant hole diameters in the mass-produced phase plates were assured. The left part of Fig. 3 shows a ZPC-TEM image of carbon thin film (Quantifoil) obtained with a Ti/SiN/Ti thin film phase plate using a field emission TEM (JEM-2200FS) and the right part shows its Fourier transform. According to the Fourier transform, we confirmed that the contrast increases in a low spatial frequency region.

Reference
[1] Danev, R. and Nagayama, K., Ultramicroscopy 88, 4, 243 (2001).
[2] Glaeser, M, Rev. Sci. Instrum. 84, 111101 (2013).
[3] Dai, W. et al, Nature, 502, 707 (2013).


Fig. 1: The micro fabrication process of a Zernike phase plate made of a Ti/SiN/Ti thin film. Micro electro mechanical system (MEMS) technology was utilized for the process.

Fig. 2: An SEM image (secondary electron) of the center hole in the thin film Zernike phase plate of Ti/SiN/Ti. The hole diameter is approximately 200nm.

Fig. 3: A ZPC TEM image of a carbon film using the phase plate made of a Ti/SiN/Ti thin film and its Fourier transform.

Type of presentation: Poster

IT-11-P-1659 Factors affecting phase noise in off-axis electron holography

Boothroyd C. B.1, Dwyer C.1, Chang S.1, Dunin-Borkowski R. E.1
1Ernst Ruska-Centrum und Peter Grünberg Institut, Forschungszentrum Jülich, D-52425 Jülich, Germany
ChrisBoothroyd@cantab.net

The amount of noise in the reconstructed wave of an electron hologram depends on the visibility of the interference fringes, which in turn depends on the coherence and intensity of the incident beam and the stability of the microscope[1]. Here we investigate the dependence of phase noise on condenser lens strength for holograms taken on a 300kV FEI Titan with two biprisms and no specimen present. We use a lower biprism voltage of 150V with no extra lens giving a fringe spacing of 0.08nm. The magnification was 450k and round illumination was used for reproducibility. A 4s exposure time, giving negligible biprism drift, was used. We did not change the gun extraction voltage[2]. The gun lens, spot size (C1 lens) and intensity (C2 & C3 lenses) were each varied starting from gun 3, spot 3 and intensity set so the beam filled the screen at 160k magnification.

For each illumination condition a hologram was taken with no specimen present, reconstructed in the standard way and the mean intensity within the hologram overlap region, the fringe contrast and the standard deviation of the reconstructed phase measured. Here the mean intensity is used as a simple measure of the coherence of the beam, a lower mean intensity is associated with a higher coherence. Fig. 1a shows a plot of the fringe contrast vs mean intensity. As the beam is made more coherent (lower mean intensity) using any of the gun lens, spot size or intensity the fringe contrast increases, as expected. The phase noise derived from the same holograms is shown in fig. 1b, from which it can be seen that the lowest phase noise is for a mean intensity of about 700 counts. When the coherence is increased so as to reduce the mean intensity below 700 counts, the increased fringe contrast is offset by increased noise due to fewer counts [3].

It can be seen from both figures that it does not matter whether the coherence is increased by increasing the spot size or the intensity, the resulting phase noise is the same. Increasing the gun lens has almost the same effect except that the fringe contrast and the phase noise are slightly worse for the highest gun lens than for the same coherence set with either the spot size or the intensity. While this observation is to be expected for a perfect microscope with no instabilities[4], it is an important result to demonstrate experimentally.

We thus conclude that for adjusting the beam coherence it makes no difference whether the intensity or spot size are used and that using the gun lens produces only slightly higher phase noise.

[1] H Lichte, KH Herrmann and F Lenz, Optik 77 (1987) 135
[2] A Lenk and H Lichte, Proc EMC 2012 ed DJ Stokes and JL Hutchison (RMS, 2012) 515
[3] WJ de Ruijter and JK Weiss, Ultramic 50 (1993) 269
[4] H Lichte, Ultramic 108 (2008) 256


Fig. 1: (a) Hologram fringe contrast (Imax-Imin)/(Imax+Imin) and (b) standard deviation phase noise plotted against mean intensity within the overlap region for different settings of the gun lens, spot size and intensity.

Type of presentation: Poster

IT-11-P-1673 Observation of Fraunhofer Diffraction Pattern with Electron Vortex Beam using Fork-Shaped Grating with Various Opening Shapes

Harada K.1, Kohashi T.1, Iwane T.1
1Central Research Laboratory, Hitachi Ltd., Hatoyama, Saitama 350-0395, Japan
ken.harada.fz@hitachi.com

      Electron vortex beams are considered as probes for next-generation electron beam machines, especially for transmission electron microscopes (TEM), because the vortex beams carry intrinsic orbital angular momentum. We expect that it will bring with it an unprecedented measurement capability.

      In order to generate a vortex beam, we fabricated a fork-shaped grating [1] made from a Si3N4-membrane with a 200-nm-thickness by using a focused ion beam machine (FB-2100, Hitachi High-Technologies Corp.). The maximum grating size in one direction was 10 μm. Electron diffractions from the gratings were observed with a 300-kV field emission electron microscope [2]. The optical system was constructed for small angle diffraction with a camera length of 150 m and was similar to the twin-Foucault imaging system [3].

      During the experiment, we noticed the shape of the grating openings was superimposed on the ring of diffraction spot, which is a typical shape of the vortex beam. We also noticed the opening size is inversely proportional to the diameter of the diffraction ring. This phenomenon is considered to be due to a combination of Fraunhofer diffractions from the grating and the opening. Figure 1 shows electron micrographs of circular-fork-shaped gratings (left panels) and their electron diffractions (right panels). The smaller the grating opening size, the larger the diameter of the diffraction rings.

      The left panels of Fig. 2 show TEM images of fork-shaped gratings with triangular, square, and pentagonal openings. The right panels show electron diffractions whose spot-shapes reflect those of the openings.

      The combination of the fork-shaped grating and its opening allowed us to observe the rotational phenomena of the diffraction rings in the through-focus condition. Figure 3 shows diffraction patterns from a fork-grating with a diamond-shaped opening for three different focuses. The diffraction spots on the right are rotated in the opposite azimuth direction to those on the left. Figure 4 plots the rotation angles of the first, second, and third diffraction rings (spots) versus the defocusing distance, Δf. They show the lower order rings rotated more. The phenomenological picture of this rotation is consistent with the phase distribution of the vortex beam. The rotation itself can be explained by the Gouy phase [4, 5].

 

References:
  [1] B. J. McMorran et al., Science, 331, 192 (2011).
  [2] T. Kawasaki et al., Jpn. J. Appl. Phys., 29, L508 (1990).
  [3] K. Harada, Appl. Phys. Lett., 100, 061901 (2012).
  [4] G. Guzzinati et al., Phys. Rev. Lett. 110, 093601 (2013).
  [5] B. J. McMorran et al., Laser Science (Frontiers in Optics, San Jose, CA, USA, 2011), LWL 1.


Fig. 1: Ring diameter of diffraction is related to opening size. Contrast of the diffraction pattern of the botom panel is enhanced.

Fig. 2: Opening shape reflects the shape of each diffraction ring.

Fig. 3: Diffractions from grating with diamond-shaped opening is rotated by defocusing.

Fig. 4: Rotation angles of the first, second and third diffraction rings on both sides versus defocusing, Δf.

Type of presentation: Poster

IT-11-P-1676 Noise estimation for off-axis electron holography

Röder F.1, Lubk A.1, Wolf D.1, Niermann T.2
1Triebenberg-Labor, Technische Universität Dresden, 01062 Dresden, Germany, 2Institut für Optik und Atomare Physik, Technische Universität Berlin, Straße des 17. Juni 135, 10623 Berlin, Germany
Falk.Roeder@Triebenberg.de

Off-axis electron holography provides access to the phase of the elastically scattered wave in a transmission electron microscope by formation of an interference pattern at the image plane (hologram) encoding amplitude and phase therein [1]. Quantitative interpretations of experimental phase shifts retrieved from these holograms additionally require the knowledge of the noise transferred through the detection and holographic reconstruction process. Only for the special case that assumes homogeneous samples, uncorrelated Poissonian distributed noise and a special reconstruction aperture corresponding to the real-space reconstruction scheme [2], noise transfer formulas were derived by F. Lenz [3]. Here, we present a general noise transfer formalism for off-axis electron holography providing access to the final covariance matrix of amplitude and phase for arbitrary objects and reconstruction apertures. As an initial condition, we need the covariance matrix of the detected hologram, which is determined by the noise transfer properties of the detector. This covariance is estimated by the recently developed noise spread function (NSF) [4, 5] using suitable approximations. To the general reconstruction formulas, we apply error propagation describing the transfer of the estimated covariance matrix of the acquired hologram into the reconstructed amplitude and phase images. We show that our derived formulas agree with the Lenz model [3], if the corresponding conditions are assumed. For the general case, we experimentally verify the presented noise transfer formulas for two different cameras (Gatan 1024x1024 CCD cameras of model MSC 794 equipped with different scintillators) with and without object. We compare the theoretically determined noise with experimentally measured noise, which is obtained by statistical evaluation of various hologram series. In Figure 1 the variances of amplitude (a) and phase (b) of empty holograms in dependence on the size of the reconstruction aperture q0 are depicted for two different cameras and show good agreement between experiment (dashed red) and theory (solid blue) within the errors. The off-diagonals of the covariance matrices of amplitude (c) and phase (d) are represented for q0 = 1/2 qc (carrier frequency), which are mainly determined by the size of the reconstruction aperture. Also these results exhibit good agreement within the errors of the measurements.

[1] H. Lichte, M. Lehmann, Rep. Prog. Phys. 71 016102 (2008)

[2] M. Lehmann, E. Völkl, F. Lenz, Ultramicroscopy 54 (1994) 335-344

[3] F. Lenz, Optik 79 (1988) 13-14

[4] T. Niermann, A. Lubk, F. Röder, Ultramicroscopy 115 (2012) 68-77

[5] A. Lubk, F. Röder, T. Niermann, C. Gatel, S. Joulie, F. Houdellier, C. Magénd, M. J. Hÿtch, Ultramicroscopy 115 (2012) 78-87


The authors gratefully acknowledge critical and inspiring discussions with Prof. Dr. Hannes Lichte (TU Dresden, Germany). The research leading to these results has received funding from the European Union Seventh Framework Programme under Grant Agreement 312483 - ESTEEM2 (Integrated Infrastructure Initiative - I3).

Fig. 1: Sample variances for amplitude (a) and phase (b) depending on reconstruction aperture size q0 (squares and circles for different cameras). The corresponding off-diagonals for q0 = 1/2 q(carrier frequency) are shown in (c) and (d)  (Δm as distance between two detector pixels). Experimental values are in red and calculated in blue. 

Type of presentation: Poster

IT-11-P-1720 Operating principles and practical applications of hole-free phase plate imaging

Malac M.1,2, Kawasaki M.3, Beleggia M.4, Pollard S.5, Zhu Y.5, Egerton R.2,1, Okura Y.6
1National Institute of Nanotechnology, Edmonton, Canada., 2Department of Physics, University of Alberta, Edmonton, Canada., 3JEOL USA Inc., MA 01960, USA., 4Center for Electron Nanoscopy, Technical University of Denmark, Lyngby, Denmark., 5Brookhaven National Laboratory, Upton, New York, USA., 6JEOL Ltd., Akishima, Tokyo 198-8558, Japan.
marek.malac@gmail.com

Ideal Zernike phase plate (PP) imaging in a TEM could, in principle, provide a quantitative measure of the phase shift induced by the sample directly from the measured image intensity for weak phase objects. In practice, the contrast transfer in PP imaging is far too complicated to allow for reliable quantification of image intensity. Furthermore, most samples are not weak phase objects. On the other hand, PP imaging, even at its current stage of development, allows to decrease the irradiation dose needed for a desired signal to noise ratio (SNR) [1], and to obtain qualitative information about samples that would otherwise require more complicated methods, such as electron holography, or complicated sample preparation. Here we present novel examples and discuss operating principles of the hole-free phase plate (HFPP) implementation [2] of PP imaging.

The HFPP implementation of PP imaging uses a uniform thin film placed in the back focal plane of the objective lens that charges due to primary beam-induced secondary electron emission. The steady-state electrostatic potential resulting from the charges phase shifts the diffracted beams relative to the direct beam resulting in strong phase contrast.

Figure 1 a) shows an example of a mouse kidney sample about 70 nm thick. Generally, biological sample of this kind are stained to obtain sufficient contrast. Instead, HFPP imaging allows to obtain sufficient contrast to study and measure lateral dimensions of the object without the need for staining. Compared to the standard bright field TEM (BFTEM) in Fig 1. b) the HFPP contrast is significantly higher. The good transfer of low spatial frequencies by the HFPP is seen in the power spectra (insets) where the HFPP in a) shows a bright region at low frequencies that is not present in BFTEM shown in b). Figure 2 a) shows an example of hard magnetic material (PrFeB) imaged using a HFPP. When compared to Fresnel imaging in Fig 2 b), new information can be obtained: the edge of the sample is clearly visible in a) while Fresnel fringes make it difficult to detect the sample edge in b). The HFPP image in a) also allow the fringing magnetic field extending into vacuum to be observed [3]. We have shown that phase plate imaging using the hole-free phase plate set up allows to establish low-dose phase contrast from samples that, when observed in Fresnel mode, would require staining. We have also shown that HFPP imaging on magnetic samples provides information that is not possible to obtain in Fresnel mode.

[1] M. Malac et. al. Ultramicroscopy 108 (2008), p. 126.

[2] M. Malac et. al., Ultramicroscopy 118 (2012), p. 77.

[3] S. Pollard et. al. Appl. Phys. Lett. 102 (2013), p.192401.


JEOL Ltd. NSERC and NRC supported this work. The samples used for Fig.1 were provided by N. Hosogi and H. Nishioka, both JEOL Ltd.

Fig. 1: .

Type of presentation: Poster

IT-11-P-1723 Direct Observation of Magnetic Domain Walls by Lens-Less Foucault Imaging (LLFI)

Taniguchi Y.1, Matsumoto H.2, Harada K.3
1Advanced Microscope Systems Design Dept., Hitachi High-Technologies Corp., Hitachinaka, JAPAN, 2Application Development Dept., Hitachi High-Technologies Corp., Hitachinaka, JAPAN, 3Central Research Laboratory, Hitachi Ltd., Hatoyama, JAPAN
taniguchi-yoshifumi@naka.hitachi-hitec.com

       Lorentz microscopy, categorized as having Fresnel and Foucault modes, is as a practical technique for observing magnetic properties by using a transmission electron microscope (TEM). The Fresnel mode is more popular because it does not require any special equipment in its electron optical system. The Foucault mode, however, requires a magnetic-field shielding lens and an off-center objective aperture. Although this mode can visualize magnetic domains under in-focus conditions, it has been considered impractical. Recently, however, a novel Foucault imaging method, named lens-less Foucault imaging (LLFI) [1], was developed for conventional TEM without any special equipment.

       Figures 1(a), (b) and (c) show the optical systems of different modes of LLFI in a 300-kV field emission TEM (HF-3300; Hitachi High-Technologies Corp.). Figures 1(a) and (b) are for observing Foucault images and (c) is for small angle electron diffraction (SAED) pattern. The objective lens was switched off and the electron beam was focused with a condenser lens to the crossover on the selected area (SA) aperture plane. The SA aperture was used as an angular aperture selecting for appropriate Foucault images, and the focal length of the magnifying lens was changed in order to observe the Foucault images and diffractions. The irradiated area on the specimen was set by selecting an appropriate diameter for the condenser aperture. Figure 1(d) is an example of SAED of a 90° ferromagnetic domain structure of La0.75Sr0.25MnO3 (LSMO). The observation was done with a camera length of 150 m. White circles in the lower panel of Fig. 1(d) indicate the diameters and positions of the SA aperture for the LLFI observations and lowercase letters of individual circles correspond to Foucault images in Figs. 1(e) – (l).

       Figures 1(e) – (h) show Foucault images of each dispersed deflection spot (see in Fig. 1(d)). The magnetization structure among the 90° domains is directly visualized. Figures 1(i) – (l) show Foucault images of the magnetic domain walls in the same area as Figures 1(e) – (h). Figure 1(i) shows 90°/180° domain walls imaged with four streaks including the optical axis by using an SA aperture 5 μm in diameter. Figure 1(j) is an image of 180° domain walls made using an SA aperture 3 μm in diameter. Figure 1(k) and (l) are 90° domain walls imaged with individual streaks from the upper and lower positions on the optical axis.

       The LLFI method is advantageous for observing not only magnetic domains and domain walls but also SAEDs less than 2×10-5 rad. In combination with high-angular-resolution imaging, it will open the way to developing new applications in Lorentz microscopy.

 

References:
 [1] Y. Taniguchi et al., Appl. Phys. Lett., 101, 093101 (2012).

 


The authors would like to thank Prof. S. Mori of Osaka Prefecture University for supplying the LSMO specimens and valuable discussion about the materials.

Fig. 1: (a) Optical system for domain observation, (b) for domain-wall, (c) for small-angle electron diffraction, (d) SAED from 90°/180° domain structure of LSMO, (e)–(h) Foucault images of 90°/180° domains from each single deflection spot, (i) –(l) Foucault images of domain walls from streaks, (j) 180° domain walls, (k) and (l) 90° domain walls.

Type of presentation: Poster

IT-11-P-1725 Twin-Foucault Imaging for Observing 180° Domains in Magnetic Materials

Harada K.1
1Central Research Laboratory, Hitachi Ltd., Hatoyama, Saitama 350-0395, Japan
ken.harada.fz@hitachi.com

     The conventional Foucault method can visualize magnetic domains under in-focus conditions, but it cannot observe the whole region irradiated by the incident electron beam at once. This is because an off-center objective aperture filters out one of the two deflected beams from materials that have 180° magnetic domains whose magnetizations are in opposite directions to one another. To solve this problem, the twin-Foucault imaging (TFI) method uses an electron biprism instead of an objective aperture to obtain two Foucault images simultaneously [1].

     Figure 1(a) shows the optical system for the TFI method using a 300-kV field-emission TEM. The electron biprism was installed between the objective and the first magnifying lenses. When a negative potential is applied to the filament electrode of the biprism, the two electron beams are deflected in dispersive directions away from the optical axis and form two individual Foucault images on the image plane simultaneously. Figure 1(b) shows small angle electron diffraction (SAED) spots from 180° magnetic domains and the shadow of the biprism.

     Figure 1(c) shows micrographs of La0.825Sr0.175MnO3 (LSMO) film taken by a CCD camera. The ordinary electron micrograph in the middle panel was divided into two series (upper/lower) of Foucault images with reversed contrast by applying negative/positive potentials to the biprism of ± 50 V and ± 100 V. The winding fringes in the central parts of these images are bend contours, and the vertical black and white stripes in Figs. 1(c) correspond to individual 180° magnetic domains. The Foucault images are switched by the polarity of the applied potential to the biprism.

     The TFI method can visualize magnetic domain structures. Figure 1(d) shows examples of processed images of the domain structures. The left panel is an image made by subtracting the right image from the left, i.e., lr, and the right panel is an image made by subtracting the left from the right, i.e., rl, in the uppermost panel of Fig. 1(c). In Fig. 1(d), the magnetic domain structures are clearly visible in the reversed contrast in images, whereas the bend contours and other contrasts have been eliminated.

     The TFI method can be used to extend the conventional Fresnel method when the applied potential to the biprism is switched off and the specimen is observed in defocused conditions. Figure 1(e) shows over- (left panel) and under-focused (right panel) Fresnel images and in-focus micrograph (center panel).

     The TFI method not only has the advantages described here; it can also be used, for example, to observe the dynamics of imaging domain switching. It will lead to new applications in Lorentz microscopy.

 

References:
[1] Harada, K., Appl. Phys. Lett., 100, 061901 (2012).


The authors would like to thank Prof. S. Mori of Osaka Prefecture University for supplying the LSMO specimens and valuable discussion about the materials.

Fig. 1: (a) Optical system for TFI method, (b) SAED pattern with shadow of a filament electrode of a biprism, (c) twin-Foucault images with different potentials from −100 V to 100 V, at intervals of 50 V, (d) 180° domain structures made by subtraction processing of twin-Foucault images, (e) Fresnel images and in-focus micrograph (center panel).

Type of presentation: Poster

IT-11-P-1727 Observation of Magnetic Field by Combination of Electron Tomography and Holography

Tsuneta R.1, Ikeda M.1, Ono S.1, Yamane M.1, Sugawara A.1, Harada K.1, Koguchi M.1
1Central Research Laboratory, Hitachi Ltd., Hatoyama, Saitama 350-0395, Japan
ken.harada.fz@hitachi.com

     Electron tomography has developed in the last decade with the progress of electron microscopy and the development of algorithms for the Radon transformation and image processing with interpolation [1]. On the other hand, electron holography is one of the standard techniques for observing phase maps of electron waves. The combination of tomography and holography has also led to elucidation of the distributions of the mean inner potential of materials [2]. In the case of the magnetic field, the component By parallel to the rotation axis (see Fig. 1(a)) can be calculated from the phase shift Δφy by using the conventional tomography algorithm [3], but the other two components (Bx, Bz) perpendicular to the rotation axis cannot be calculated, because they are mixed together with a rotation angle θ. In the case of a magnetic field in free space, however, the phase shift Δφθ projected to the optical axis can be described simply [4], for which Bx and Bz can be separated.

     Figure 1(a) is a schematic diagram of the tomography/holography experiment. A thin magnetic pillar made of C/CoFeB/SiO2 was prepared with a focused ion beam machine (see Fig. 1(b)) and mounted on the 360° rotation axis of the specimen holder. Two holograms projected from the front surface and back surface of the material were processed in order to divide the magnetic and electric fields around the specimen. Figures 1(c) and (d) are examples of reconstructed interferograms of the magnetic and electric fields. Holograms made with a double-biprism interferometer [5] for tomography were recorded every 10°. Eighteen phase maps corresponding to the projected magnetic field were reconstructed from thirty-six holograms. The phase maps were processed into a three-dimensional (3-D) magnetic field distribution in free space by using a modified tomography algorithm.

     Figure 2 shows the reconstructed 3-D magnetic field distribution on the plane perpendicular to the pillar-shaped magnet. Figure 3 shows the reconstructed magnetic field in three planes parallel to the pillar.

     To improve the precision and resolution of 3-D reconstructions of the magnetic field, especially inside the material, the two components of the magnetic field should be measured independently. A dual-axis 360° rotation specimen holder has already been developed for this purpose [6]. The results of an experiment using this dual-axis holder will be reported soon.

 

References:
[1] S. Ono et al., Appl. Phys. Express, 4, (2011) 066601.
[2] G. Lai et al., Appl. Opt., 33, (1994) 829.
[3] G. Lai et al., J. Appl. Phys., 75, (1994) 4593.
[4] H. Shinada et al., IEEE Transaction on Magnetics, 28, (1992) 1017.
[5] K. Harada et al., Appl. Phys. Lett., 84, (2004) 3229.
[6] R. Tsuneta et al., Kenbikyou, 48, (2013) 205 (in Japanese).


The authors would like to thank Mr. H. Hasegawa of the Central Research Lab., Hitachi Ltd. for supplying the C/CoFeB/SiO2 specimens and valuable discussion.

Fig. 1: (a) Experimental setup for tomography/holography, (b) scanning ion micrograph of C/CoFeB/SiO2 pillar, (c) reconstructed interferogram of magnetic field, (d) interferogram of electric field.

Fig. 2: Reconstructed 3-D magnetic field distribution in one horizontal plane.

Fig. 3: Reconstructed 3-D magnetic field distributions in three vertical planes (a, b and c).

Type of presentation: Poster

IT-11-P-1784 Accumulation Processing in Reconstruction for Electron Holography

Kasai H.1, Harada K.1
1Central Research Laboratory, Hitachi Ltd., Hatoyama, Saitama 350-0395, JAPAN
hiroto.kasai.qm@hitachi.com

 

     One of the long standing problems affecting electron holography, the lateral coherence limitation, has been solved by using the split illumination method with a specially customized TEM. The customized TEM has one or two biprisms in a condenser optical system [1, 2]. Conventional holography TEM, however, does not have any biprisms in the condenser system. The problem, therefore, remains unsolved in practice. In order to solve the problem by using a conventional TEM, an "accumulation processing" method was developed.

     Figure 1(a) explains the concept. The target sample is in the object region (n) far from the specimen edge, and the distance from the region (n) to the vacuum region (ref) exceeds the coherent length R. The first hologram is recorded with two waves in region (n) and region (n-1); the second hologram is also recorded with these waves but in regions (n-1) and (n-2). The specimen is shifted perpendicular to the hologram through a distance W that is equal to the hologram's width. The last hologram is recorded with the waves in region (1) and the reference region (ref).

     The reconstructed phase distribution Δηn is described as the difference between the phase distributions ηn of the waves: Δηn = ηn-ηn-1. After reconstruction of all of the holograms, the reconstructed phase distributions are summed one by one, i.e., ΣΔηn =Σ(ηn-ηn-1) =ηn-ηref= ηn. The summed result is the same as the phase distribution reconstructed from the hologram recorded in region (n) and the reference region (ref). This means that the problem is solved in principle.

     Figure 1(c) shows a conventional phase map corresponding to the magnetic lines of force from the apex of the magnetic force microscopy (MFM) tip (see Fig. 1(b)). When the hologram width was changed, the density of the magnetic lines of force changed. This means that the magnetic field from the tip leaked into the reference region. Thus, it is important to keep the reference wave far from the magnetic material.

     Figure 1(d) shows the phase distribution reconstructed by the accumulation method. All of the regions were reconstructed by using just one reference wave and put in order from (n) to (1). The reconstructed area was seven times as wide as the possible area reconstructed by conventional holography. Accordingly, magnetic lines of force not only from the apex of the tip, but also from other areas can be visualized.

     Figure 1(e) shows the subtracted phase distribution processed from (d). A comparison of the results in Figs. 1(c), (d) and (e) clarifies that conventional holography visualizes differentiation of the magnetic field around the tip.

 

References:
[1] T. Tanigaki et al., Appl. Phys. Lett., 101, 043101 (2012).
[2] T. Tanigaki et al., Ultramicroscopy, 137, 7 (2014).


Fig. 1: Fig. 1 (a) Principle of "accumulation processing", (b) SEM image of MFM tip, (c) magnetic lines of force (X 4) from the apex of the tip reconstructed by the conventional method, (d) magnetic lines of force (X 1) around the tip reconstructed by the accumulation method, (e) differential distribution from (d).

Type of presentation: Poster

IT-11-P-1790 Phase Reconstruction in Annular Bright Field STEM

Ishida T.1, Kawasaki T.1, Kodama T.2, Ogai K.3, Ikuta T.4, Tanji T.1
1Nagoya University, Nagoya, Japan, 2Meijo University, Nagoya, Japan, 3APCO Ltd., Hachioji, Japan, 4Osaka Electro-Communication University, Neyagawa, Japan
takafu_i@echo.nuee.nagoya-u.ac.jp

In scanning transmission electron microscopy (STEM), an annular bright field (ABF) imaging [1] enabled simultaneous imaging of light and heavy elements. The ABF image contrast does not change in a thick specimen. On the other hand, in a thin specimen, the ABF image contrast oscillates by thickness and defocus [2,3]. This effect seems to same as the conventional bright field phase contrast imaging. Exact interpretation of the ABF image needs to be reconstructed to phase and amplitude information on specimens. In STEM, a new phase reconstruction technique revealed phase and amplitude images using a multi-channel detector [4]. We apply this technique to the ABF. This is called ABF phase imaging. The ABF phase imaging requires an annularly arrayed 24 detectors. We will show that this new method works effectively for reconstructing the phase of electron wave.

In the present study, the STEM (Hitachi HD-2300S) equipped with a spherical aberration corrector and an annular condenser aperture. The STEM was operated at the acceleration voltage of 200 kV with a convergence semi-angle of 20 - 25 mrad. The annular array detector is placed under a post-specimen projection lens. The inner- and outer-side angles of the annular array detector were set to 20 and 25 mrad by changing the excitation of the post-specimen projection lens. The STEM system takes 24 images simultaneously with a scanning time of 8.3 second.

Figs. 1(a)-(c) show some parts of 24 images of the graphite and the number k in these images corresponds to the position of the detector used for forming these images, as shown in Fig. 1(d). Images obtained with different detectors show different direction lattice fringes of 0.34 nm. Fig. 2 shows the reconstructed amplitude and phase images. The phase image is clearly visible in 0.34 nm lattice fringes. In contrast, the amplitude image mainly shows the thickness of the specimen.

We also carry out Image simulation for atomic resolution ABF phase images using multislice method. The image simulation will be presented an image contrast is proportional to atomic number Z. These results show the capability of this new method for a high resolution electron phase retrieval technique.

 References
[1] S.D. Findlay, N. Shibata, H. Sawada, E. Okunishi, Y. Kondo, T. Yamamoto and Y. Ikuhara, Appl. Phys. Lett. 95 (2009) 191913
[2] R. Ishikawa, E. Okunishi, H. Sawada, Y. Kondo, F. Hosokawa and E. Abe, Nat. Mater. 10 (2011) 278–281
[3] S. Lee, Y. Oshima, E. Hosono, H. Zhou, K. Takayanagi, Ultramicroscopy 125 (2013) 43–48.
[4] M. Taya, T. Matsutani, T. Ikuta, H. Saito, K. Ogai, Y. Harada, T. Tanaka, Y. Takai, Rev. Sci. Instrum. 78 (2007) 083705.


This work was supported by JSPS KAKENHI Grant Numbers18GS0211, 24360020.

Fig. 1: (a)-(c) Observed images obtained by the annular array detector. (d) Schematic of the detector layout.

Fig. 2: Reconstructed (a) amplitude and (b) phase images, respectively.

Type of presentation: Poster

IT-11-P-1940 Experimental evaluation of magnetic phase reconstruction in Lorentz TEM using the ‘transport-of-intensity’ equation

Kohn A.1, Habibi A.1
1Department of Materials Engineering and Ilse Katz Institute for Nanoscale Science and Technology, Ben-Gurion University of the Negev
akohn@bgu.ac.il

Imaging the micromagnetic structure of materials at the nanometer scale is motivated by the scientific study of new magnetic phenomena, and the technological drive for new information storage devices, which increase the storage density.
The micromagnetic structure can be imaged using transmission electron microscopy (TEM) in a variety of contrast modes, termed ‘Lorentz TEM’; for example, Fresnel-contrast (defocused images). Magnetic imaging in the TEM is possible because in the presence of a magnetic (and electric) potential, the electron wave-function undergoes a phase shift. Therefore, for quantitative mapping of the magnetic induction, the phase shift of the electron-wave needs to be reconstructed.

The ‘transport-of-intensity’ equation (TIE) is a general phase reconstruction methodology that can be applied to Lorentz TEM through the use of Fresnel-contrast images. We present an experimental study of sub-micrometer sized Permalloy elements in order to test the application of the TIE for quantitative magnetic mapping. We find that quantitative phase reconstructions (e.g. Fig. 1) are possible for defoci distances ranging approximately between 200 and 800 μm. The lower defocus limit is attributed to competing sources of image intensity variations in the Fresnel-contrast images such as structural defects and diffraction contrast. The upper defocus limit is shown to originate from a numerical error in the estimation of the intensity derivative.

Three sources of magnetic phase information are compared: domain walls, element edges and vortex cores. The vortex cores are shown to enable quantitative phase reconstructions while the domain walls and element edges enable only qualitative phase reconstructions. Considering the above limitations, we show quantitative reconstructions of elements sized down to approximately 100 nm and 5 nm thick. Thus, the minimal detection of the product of the magnetic induction and thickness is 5 Tesla·nanometer and magnetic structures are spatially resolved down to a size of 12 nanometers.


Fig. 1: Calculated (a, b, c, d) and experimental (e, f, g, h) eqi-phase contour maps (spaced at 1 radian) for Permalloy elements: triangular, 1µm diagonal, 10nm thick (a, e), circular, 250nm in diameter, 20nm thick (b, f), square, 130nm edge, 10nm thick (c, g) and circular, 1µm in diameter, 5nm thick (d, h).

Type of presentation: Poster

IT-11-P-1945 Dopant mapping in thin FIB prepared silicon samples by off-axis electron holography

Kohn A.1, Pantzer A.1,2, Vakahy A.2, Eliyahou Z.1, Levi G.2, Horvitz D.2
1Department of Materials Engineering and Ilse Katz Institute for Nanoscale Science and Technology, Ben-Gurion University of the Negev , 2Micron Semiconductors Israel Ltd.
akohn@bgu.ac.il

Modern semiconductor devices function due to accurate dopant distribution. Off-axis electron holography (OAEH) in the transmission electron microscope (TEM) can map quantitatively the electrostatic potential in semiconductors with high spatial resolution. For the microelectronics industry, ongoing reduction of device dimensions, 3D device geometry, and failure analysis of specific devices require preparation of thin TEM samples, under 70 nm thick, by focused ion beam (FIB). Such thicknesses, which are considerably thinner than the values reported to date in the literature, are challenging due to FIB induced damage and surface depletion effects.
We report1 on preparation of TEM samples of silicon PN junctions in the FIB completed by low-energy (5 keV) ion milling, which reduced amorphization of the silicon to 10 nm thick. Additional perpendicular FIB sectioning (e.g. Fig. 1) enabled a direct measurement of the TEM sample thickness in order to determine accurately the crystalline thickness of the sample. Consequently, we find that the low-energy milling also resulted in a negligible thickness of electrically inactive regions, approximately 4 nm thick. The influence of TEM sample thickness, FIB induced damage and doping concentrations on the accuracy of the OAEH measurements were examined by comparison to secondary ion mass spectrometry measurements as well as to 1D and 3D simulations of the electrostatic potentials. We conclude that for TEM samples down to 100nm thick, OAEH measurements of Si-based PN junctions, for the doping levels examined here, resulted in quantitative mapping of potential variations, within ~0.1V. For thinner TEM samples, down to 20nm thick, mapping of potential variations is qualitative, due to a reduced accuracy of ~0.3V (Fig. 2).

1. A. Pantzer et al., Ultramicroscopy 138 (2014) 36–45


We thank A. Ripp for advising and implanting PN junction samples; M. Sokolovsky for SIMS analysis; I. Amit and Y. Rosenwaks for assistance with ‘Sentaurus’ 3D simulator.

Fig. 1: Example (50 nm thick sample) of additional perpendicular FIB sectioning to measure layer thicknesses directly: (a) bright-field TEM overview image of the sample; (b) High-resolution TEM image of the region denoted schematically by the black rectangle in (a). A 6 nm thick amorphous Si surface layer is measured and 154 nm thick crystalline Si layer.

Fig. 2: Comparison of potential profiles for sample thicknesses varying between 20 and 100 nm as derived from OAEH to the band potential as calculated by a 1D Poisson simulation using data from SIMS measurements.

Type of presentation: Poster

IT-11-P-1954 Specimen Charging Measured by Off-axis Electron Holography

McLeod R. A.1, Beleggia M.2,3, Malac M.4
1Fondation Nanosciences, Grenoble, France, 2Denmark Technical University, Lyngby, Denmark, 3Hemholtz-Zentrum-Berlin, Berlin, Germany, 4National Institute for Nanotechnology, Edmonton, Canada
robbmcleod@gmail.com

In off-axis electron holography, the acquisition of hologram series for the purpose of averaging is becoming popular to improve the signal-to-noise ratio. Image series can also be considered as a form of tomography, with the 3rd axis being time. In this case, a hologram series can measure dynamic electric and magnetic characteristics of a specimen.


Hologram series have a total exposure time of minutes. During this period, interaction with the electron beam will cause many secondary electrons to be emitted from the specimen. The generated holes are screened by the mobile charges and the dielectric response of the material, generating a screening Coloumb potential. If a metal is nearby, the potential is further screened by an image charge, resulting in a phase shift of dipole character. Thus the holes left-behind by SE emission constitute a form of radiation damage for phase retrieval techniques. The holes will be refilled at a rate determined by the conductance and morphology of the specimen. The hole half-life relative to the total exposure time governs the ultimate accuracy of the measured phase. The technique is expected to improve understand of charging behavior, especially in insulating specimens.


Here we have conducted experiments with a series of short exposure (0.25 – 0.5 s) off-axis holograms at 300 keV of latex nanoparticles on a lacey carbon support. The beam was initially blanked for ~15 minutes. Based on the Berriman effect, we expect the latex particle to only partially discharge over this time period. The beam was unblanked at the start of the hologram series acquisition. The experiment was repeated at three current density levels: high (1240 A m-2), medium (260 A m-2, shown in Fig. 1), and low (70 A m-2). The phase shift difference inside the particle boundaries between the start and end of the series measured for the high (0.8 rad) and moderate (0.2 rad, shown in Fig. 2) current densities. For the low current density series the phase error (~0.12 rad) was too high to directly estimate the phase shift. To estimate the charge on the latex particle as a function of dose, the vacuum projected potential was found from the rotational average of the vacuum phase around the particle. A least-squares best-fit to a model of a surface-charged sphere plus image charge yielded an estimate of the relative charge, as shown in Fig. 3. The measured phase is relative rather than absolute because the average phase of the entire series was used as a reference to remove fringing fields. The high and medium current densities charged the particle, whereas for the low current density there was actually discharge during illumination. The time-resolved behavior is visualized by videos, which will be shown at the conference.


RAM acknowledges the financial support of Fondation Nanosciences and CEA and MM the financial support of NINT. RF Egerton provided valuable discussion.

Fig. 1: Phase contours (0.1 rad each) for the averaged phase shift over the medium dose-rate series.

Fig. 2: Phase difference between first seven and last seven holograms for medium dose-rate series.

Fig. 3: Accumulated electron charge as a function of cumulative dose for the three current density levels. The total series exposure time was 119 s for the low and med cases and 141 s for the high case.

Type of presentation: Poster

IT-11-P-2006 Solving the transport of intensity equation using flux-preserving boundary conditions

Parvizi A.1, Koch C. T.1
1Institute for Experimental Physics, Ulm University, Albert-Einstein-Allee 11, 89081 Ulm, Germany
amin.parvizi@uni-ulm.de

The Transport of Intensity Equation (TIE) is a non-interferometric phase reconstruction method which overcomes disadvantages of interferometric methods such as coherent illumination and interferometer stability [1]. The TIE is a Poisson type equation which relates a modified Laplacian of the phase of the wave to the intensity variation along the optical axis (see eqn.(a.1) in Fig 1). In the absence of singularities in the principle (central) plane of focus, many approaches of solving the TIE exist. The problem common to all these approaches is that the necessary boundary conditions are not known. The most popular approach is based on the Fast Fourier Transform (FFT) and solves the TIE non-iteratively in the frequency domain. The implicitly assumed periodic boundary conditions are the main drawback of this method [1]. Another method is based on the multigrid approach for solving partial differential equations, yielding an exact solution of the TIE in the spatial domain. This approach allows us to define zero-flux boundary conditions (Fig1,eq. (a.2), where n ⃑ is the normal to the boundary) and the a vanishing phase shift within regions of vacuum. Since the phase in vacuum is constant, a circle (the number or the shape is arbitrary) with Dirichlet boundary condition can be placed in an area containing vacuum. Figure 1b shows the graphical representation of above-mentioned boundary conditions. Fig.2a and b show under and over focused images which are used to compute the intensity variation along the optical axis (Fig.2c).
Fig. 3 compares phase maps reconstructed from the simulated data by different methods. Fig. 3c shows the phase recovered by the Fourier method and our finite element multigrid solution obtained by using the software package COMSOL is shown in Fig. 3b. This figure shows clearly that especially the low frequency details of the phase reconstructed by the flux-preserving approach agree better with the phase used for simulating the images, than the reconstruction obtained by the Fourier transform method. In our presentation we will show applications of this method to experimental data acquired in the TEM and also the optical microscope. 
[1] D. Paganin and K.A. Nugent, Non-interferometric phase imaging with partially-coherent light Phys. Rev. Lett., 80, 2586-2589 (1998)


This work was supported financially by the Carl Zeiss Foundation as well as the German Research Foundation (DFG, Grant No. KO 2911/7-1).

Fig. 1: a) (a.1)TIE equation and (a.2) Equation defining the zero-flux boundary condition b) Graphical representation of the combination of boundary conditions used for the reconstruction scheme presented here.

Fig. 2: Under-focused image b) Over-focused image c) dI/dz approximated by the difference of the images shown in a and b.

Fig. 3: a) Original phase, b) Phase reconstructed by the method presented here c) Phase reconstructed by the FFT method for solving the TIE.

Type of presentation: Poster

IT-11-P-2111 Optimisation of spatial and phase resolution of off-axis electron holography for detection of single dopant atoms.

Mayall B.1, McLeod R. A.2,3, Cooper D.1
1CEA-LETI, Minatec, 17 rue des Martyrs, 38054 Grenoble Cedex 9, France, 2Fondation Nanosciences, Grenoble, France, 3CEA-INAC, Minatec, 17 rue des Martyrs, 38054 Grenoble Cedex 9, France
benjamin.mayall@cea.fr

The reduction of the size of semiconductor devices leads to use of higher dopant concentrations. The scale of these devices implies that what was known as a very high dopant concentration will be only a few tens of atoms. The small numbers of these dopants mean that the just one of these atoms out of place will become statistically significant with respect to the electrical properties of the device. Thus, it is becoming increasingly important to be able to characterise the position and activity of these individual dopants.

Assuming a perfect specimen and vacuum, simulations show a phase sensitivity of better than 2π/2000 will be required in order to detect a single ionised dopant atom in silicon. The current state of the art for atomic resolution is around 2π/200, thus an improvement of an order of magnitude is required. The phase sensitivity of a reconstructed phase image is given by the relation,

Δφ≈(2/NV2)0.5

where N is the number of electron counts and V is hologram contrast. In order to achieve the best spatial resolution at atomic resolution it will not be possible to acquire holograms for very long periods due to specimen drift. Thus different approaches must be employed to improve the sensitivity.

Off-axis electron holography has been performed using a double aberration corrected FEI Titan Ultimate TEM equipped with an X-FEG. To improve the number of counts, a large series of holograms can be added together [1]. Fig. 1 shows a phase image of a silicon calibration specimen delta-doped with boron atoms. In Fig. 1(a) a single phase image acquired for 8 seconds is shown whereas (b) shows improvements from summing 25 holograms. In order to improve the contrast of the holograms, a monochromator can be used. Fig. 2 shows the effects of using the monochromator on the fringe contrast measured on reference holograms. An increase from 25 to 35 % will make a significant step towards the target of 2π/2000. Other approaches have been used to improve the reconstruction procedure. The spatial resolution of holograms containing strong phase objects is usually one third of the carrier frequency to avoid cross-talk from the centreband. The suppression of centreband in Fourier space by using phase shifting holography [2] improves the spatial resolution for a given carrier frequency and relaxes the electron biprism bias, resulting in higher V.

In this presentation we will show how combinations of these different methods of optimising phase noise and spatial resolution have been combined and then applied to wedge polished silicon specimens that contain different types of dopant atoms.

[1] R. McLeod et al. In press Ultramicroscopy (2013)

[2] V. Volkov et al. Ultramicroscopy 134 p175 (2013)


This work has been performed on the nanocharacterisation platform (PFNC) at Minatec. DC thanks the European Research Council for the Starting Grant “Holoview”.

Fig. 1: (a) Secondary Ion Mass Spectrometry (SIMS) profile of boron doped delta layers. (b) Reconstructed phase image from single hologram acquired for 8 seconds. (c) Reconstructed phase image from 25 holograms summed together.

Fig. 2: Electron holograms acquired without (a) and with (b) monochromator active. (c) Fringe intensity profiles acquired from hologram and (d) measured contrast as a function of C2 with and without monochromator.

Type of presentation: Poster

IT-11-P-2202 Toward a 3D strain mapping at nanometer scale with Dark-Field Electron Holography

Javon E.1, Gatel C.2, Hÿtch M. J.2, Bals S.1, Van Tendeloo G.1
1EMAT, University of Antwerp, Antwerp, Belgium, 2CEMES-CNRS and Université de Toulouse, Toulouse, France
elsa.javon@uantwerpen.be

Dark-Field Electron Holography (DFEH) is a technique developed to map strain at the nanometre scale with a large field of view. The principle is based on measuring the variation phase of diffracted beam due to small strain which is the so called geometric phase1. Due to the diffraction conditions used for obtaining dark-field holograms, it is not possible to combine classical electron tomography with DFEH as it was done with electron holography for 3D magnetic and electric field mapping. Until now and similarly to GPA, two diffracted beams were selected to map 2D strain field projected on the electron direction. Here, we propose to combine 3 or 4 non collinear diffracted beams to reach the 3D strain map of a multilayer sample constituted by the repeated stacking of SiGe/Si layers grown along the [001] direction.
In order to keep the same dynamical conditions, all the diffracted beams belong to the {220} family. Two of them (0-22) and (02-2) are included in the [100] zone axis and two others (20-2) and (-20-2) in the 90° tilted zone axis namely [010] as shown in the Fig.1. Since DFEH imposes very strict constraints relating to the sample such as flatness and uniformity of the thickness on both reference and strain areas, specific FIB techniques have been developed for 2D strain mapping with DFEH. Here, we created a needle sample with a squared cross section in order to obtain a symmetric configuration and exactly the same thickness in both zone axis (Fig.2). The value of the thickness was chosen such that the crystalline thickness corresponds to a half integer of the extinction distance2 ξ220 and that the field of view was large enough.
The strain maps corresponding respectively to [100] and [010] are presented in the Fig.3. Since sample edges yield relaxation effects which are even stronger in strain regions, the measured profiles does not exactly correspond to the expected bulk values. However, the similarity between the strain profiles from both zone axis proves of the symmetry of the boundary conditions. All the projections that are required to reconstruct the 3D strain map are now known and the information common of both zone axis serves to validate the method.

In this study we present a method yielding for the first time a 3D reconstruction of the strain at nanometer scale. These results are also compared with the new 2-beams theory considering dynamical effects and resulting simple projection rule for the measurement of the geometric phase within the thickness. To carefully develop this method we chose a well characterized strained sample, however, the scope of this new technique is of high importance for entirely controlling the electron mobility of new nano devices.

1.M.Hÿtch et. al Nature453(2008)
2.A.Lubk, E.Javon et. al Ultramic.136(2014)


The authors acknowledge financial support from the European Union under the Seventh Framework Program under a contract for an Integrated Infrastructure Initiative (ESTEEM2) and the French National Agency (ANR) in the frame of its program in Nanosciences and Nanotechnologies.

Fig. 1: Schematic illustration of the needle preparation: a squared cross section yielding the same thickness in both [100] and [010] zone axis.

Fig. 2: (a) SEM image of a FIB of the squared needle. (b) Conventional dark field image of the multilayer SiGe.

Fig. 3: Strain map εzz corresponding respectively to the (a) [100] and (b) [010] zone axis. (c) Comparison of both strain profiles from the [100] and [010] zone axis along the [001] direction with the bulk value.

Type of presentation: Poster

IT-11-P-2409 Low-keV electron microscope based on a single-atom electron source

Chang W. T.1, Lin C. Y.1, Hsu W. H.1, Chen Y. S.1, Hwang I. S.1
1Institute of Physics, Academia Sinica, Taipei, Taiwan, R.O.C.
wtchang@phys.sinica.edu.tw

Imaging of biological and organic molecules is a challenge in current electron microscopes due to insufficient contrast. Low-energy electrons emit from a single-atom tip1 (SAT) provide better contrast in light element materials and low radiation damage, and wave phases scattered from an object can be used to reconstruct images of the object due to its full spatial coherence2. Therefore, development of low-energy electron microscopes based on a reliable single atom emitter is highly desirable in study of light element materials.
We have built a low-energy electron point projection microscope (PPM) to evaluate emission characteristics of single atom emitter, as shown in Fig. 1. The beam divergence angle, measured at the half maximum intensity [Fig. 2(a)], becomes larger and the beam size at extractor becomes smaller with decreasing emitter-extractor separation [Fig. 2(b)]. This measurement may provide useful information for constructing an electron gun based on a single-atom source.
The beam energy used in PPM is always less than 500 eV and is not suitable for imaging molecules of 1 nm or thicker because the inelastic mean free path is less than 1 nm. Here we propose a low-keV (500 eV~ 5 keV) electron microscope based on a single-atom electron gun and a focusing lens, as shown in Fig. 3(a). In this electron energy range, samples with thickness of 3 nm or thinner can be imaged. It allows different imaging modes, including SEM imaging, coherent electron diffractive imaging, and holographic imaging.
A new challenge for the design is alignment. A narrow beam emitted from a single-atom source facilitates focusing, but also makes the alignment of the electron beam become critical, because misalignment by merely 1˚ causes a significant decrease in the beam intensity. Design for fine alignment of the source is essential for the successful extraction of the electron beam. Therefore, a piezo-stage with precision linear and angular adjustments is used to fine position the emitter. The custom-made stage enables building a compact and rigid system and provides the freedom to align and optimize the performance the entire system, as shown in Fig 3(b). Fig 3(c) presents a diffraction pattern of graphene with an imaging area around ~20 μm. We are improving the lens system to get better focusing of the system. This new instrument may allow determination of the atomic structures of individual thin nano-objects, such as graphene, carbon nanotubes, DNA molecules, or protein molecules.

Ref:

1.  Nano Lett. 4 (2004), 2379.

2. Nanotechnology 20 (2009), 115401.


This work was sponsored by the Academia Sinica. We pay our great thanks to their financial support

Fig. 1: Schematic diagram of a point projection microscope based on a single-atom emitter.

Fig. 2: (a) Intensity profile of an extracted beam from a single-atom electron source; (b) Opening angle & beam size v. s. the emitter-extractor separation.

Fig. 3: (a) Schematic diagram of the low-keV microscope. (b) Photograph of the microscope. (c) Diffraction pattern of graphene.

Type of presentation: Poster

IT-11-P-2437 Reconstruction of the projected crystal potential using HRTEM – prospects for materials science investigations

Lentzen M.1, Barthel J.2
1Ernst Ruska Centre, Research Centre Jülich, Jülich, Germany, 2Central Facility for Electron Microscopy, RWTH Aachen University, Aachen, Germany
m.lentzen@fz-juelich.de

Potential reconstruction is the logical continuation of wave function reconstruction [1–3] in high-resolution electron microscopy. It aims at eliminating the problems in the structural interpretation of reconstructed wave functions, chiefly imposed by the effects of dynamical electron diffraction. These effects cause a non-linear relation of atomic scattering power and modulation of the wave function [4], and the local modulation near atomic columns can be further obscured through delocalisation and asymmetries induced by crystal tilt.

A series of investigations using the channelling model of dynamical electron diffraction [4] and a rapid and stable potential reconstruction algorithm revealed that the projected crystal potential can be determined for thick objects [5]. Object thickness, residual defocus aberration of the wave function, and phenomenological absorption, parameters often unknown in experiment, can be fitted self-consistently together with the projected potential [6], as well as crystal tilt [7].

In a materials science investigation of a thin Ba0.5Sr0.5Co0.8Fe0.2O3 (BSCF) crystal a through-focus series of 20 images was recorded with an aberration-corrected TITAN 80-300 microscope operated at 300 kV. After wave function reconstruction and numerical aberration correction up to the information limit of 0.08 nm (Fig. 1, right) the projected potential (Fig. 2, left) was reconstructed with a best fit of 8.4 nm object thickness, 1.6 nm residual defocus, 7.0 nm–1 crystal tilt, and a small residual of S = 4.5%. The potential map is free from non-linear contrast modulation, and the effects of tilt are strongly reduced. Column-by-column measurement of the potential maxima at the oxygen sites reveals through a histogram single oxygen atom precision of 2.6 volt per atom (Fig. 2, right). The three maxima of the distribution indicate a high concentration of oxygen vacancies.

[1] H Lichte, Ultramicroscopy 20 (1986), p. 293.

[2] W Coene et al, Phys. Rev. Lett. 69 (1992), p. 3743.

[3] A Thust et al, Ultramicroscopy 64 (1996), p. 211.

[4] K Kambe, G Lehmpfuhl and F Fujimoto, Z. Naturforsch. A29 (1974), p. 1034.

[5] M Lentzen and K Urban, Acta Cryst. A56 (2000), p. 235.

[6] M Lentzen, Ultramicroscopy 110 (2010), p. 517.

[7] M Lentzen, Proceedings MC2011 Kiel (2011), IM2 P133.


JB gratefully acknowledges funding from the German Federal Ministry of Economics and Technology within the COORETEC initiative.

Fig. 1: (left) High-resolution image of BSCF at bright-atom pass-band conditions, red: Ba/Sr, green: Co/Fe, blue: O. (right) Phase of reconstructed exit wave function of BSCF. Frames are 3.4 nm × 3.4 nm.

Fig. 2: (left) Reconstructed projected potential of BSCF, frame 3.4 nm by 3.4 nm, red: projected unit cell, blue: oxygen columns used for histogram analysis, dashed: Σ3 twin boundaries. (right) black: frequency of oxygen potential maxima versus maximum of oxygen potential (volt), grey: fit of the distribution with three gaussians.

Type of presentation: Poster

IT-11-P-2687 Direct Imaging of Two-Dimensional Electron Gas at Oxide Interfaces using Inline Electron Holography

Song K.1, Ryu S.2, Lee H.2, Choi S.3, Paudel T. R.6, Koch C. T.4, Rzchowski M. S.5, Tsymbal E. Y.6, Eom C.2, Oh S.1
1POSTECH, Pohang, Republic of Korea, 2University of Wisconsin-Madison, Madison, USA, 3Korea Institute of Materials Science, Changwon, Republic of Korea, 4Ulm University, Ulm, Germany, 5University of Wisconsin-Madison, Madison, USA, 6University of Nebraska, Lincoln, USA
ksong@postech.ac.kr

Recently, a variety of new physical properties and phenomena have been discovered to emerge at atomically engineered interfaces of complex oxide systems. One example is the two-dimensional electron gas (2-DEG) forming at the interface between two insulating perovskite oxides, LaAlO3 (LAO) and SrTiO3 (STO). Theoretically, the electron concentration at this atomically-controlled interface can be manipulated by means of the polarity-induced electric field, which is facilitated by simply changing the film thickness of LAO on a STO substrate. The resulting conducting “interface material” is known to be localized within a few nm from the interface. Although the existence of 2-DEG has been proved and utilized in many prototype devices, there are still compelling debates related to the origin, spatial distribution, and electrostatic compensation of this “interface material”. Here, we directly visualize and quantify the 2-DEG forming at the interface of LAO/STO by using inline electron holography. By combining electron energy loss spectroscopy (EELS) and quantitative measurements of the atomic displacement of cations the possible origin of 2-DEG will be discussed.
A wide area 2-D charge density map with sub-nanometer resolution (~0.8 nm) was obtained by applying a Laplacian image filter to the electrostatic potential map as shown in Fig. 1. The electrostatic potential maps were retrieved by carefully calibrating the mean inner potentials and local thicknesses of LAO and STO. Whilst the charge density map obtained from the 3 unit cell (u. c.) sample, which is below the known critical thickness of 4 u. c., seems not to host any significant charge density near the interface (Fig. 1a), the 10 u. c. sample exhibits the negative charges beneath the interface (Fig. 1b). The width of 2-DEG, measured at full width at half maximum, is 0.82 ± 0.34 nm. In order to extract a correct density value of 2-DEG from the total charge density map, one has to take account of a change of the dielectric constant of STO near the interface due to a large intrinsic electric field. The calibration of the dielectric constant using Landau theory yields a 2-DEG density close to the theoretical expected value of ~3.3×1014 e cm-2 corresponding to the transfer of 0.5 e per unit cell. The scanning transmission electron microscopy (STEM) analysis combined with EELS indicates that the origin of this interfacial 2-DEG is most likely related with the oxygen vacancies formed at the LAO surface (Fig. 3), which agrees well with the recent first principles calculations.


This works has been supported by the AFOSR under Grant numbers FA2386-13-1-4136 and FA9550-12-1-0342.

Fig. 1: Charge density maps across the LaAlO3/SrTiO3 interfaces obtained by using inline electron holography for: a. 3 u. c. and b. 10 u. c. LaAlO3 samples.

Fig. 2: Charge density profile extracted from the charge density map of 10 u. c. LaAlO3/SrTiO3 shown in Fig. 1b.

Fig. 3: High-angle annular dark field (HAADF) image of a 10 u. c. LaAlO3/SrTiO3 and EELS spectra of Ti-L2,3 and O-K edges obtained from 2D line scans.

Type of presentation: Poster

IT-11-P-2718 80 kV double biprism electron holography

Genz F.1, Niermann T.1, Lehmann M.1
1Technische Universität Berlin, Institut für Optik und Atomare Physik, Straße des 17. Juni 135, 10623 Berlin, Germany
florian.genz@physik.tu-berlin.de

Off-axis electron holography is a powerful method to retrieve the image phase in high quality[1]. Recently, off-axis electron holography performed with an acceleration voltage of 80 kV has moved into the focus of interest because electrons accelerated with 80 kV produce less knock-on damage and hence can be used to investigate, e.g., carbon-based materials [2]. The performance of a double biprism setup at 80 kV acceleration voltage and possible advantages in comparison to 300 kV were investigated.
A FEI Titan transmission electron microscope equipped with a high-brightness Schottky field-emission gun, an image Cs-corrector and a 2k by 2k camera (Gatan US1000) was used for this work. The camera's modulation transfer function (MTF) was determined with the edge method (fig. 1). The MTF is significantly improved at an acceleration voltage of 80 kV compared to 300 kV, probably due to a smaller electron scattering volume in the scintillator. The improvement is best for a sampling rate of about 10 pixels per fringe, giving an improvement factor of 1.9. In experimental electron hologram series with different fringe spacings, but keeping a sampling rate of 10 pixels per fringe realized by accordingly adjusting the magnification, about the same improvement of the standard deviation of the reconstructed phase was measured (fig. 2). Thereby, other effects like, e.g., electron energy dependent detection quantum efficiency of the camera, seem to be negligible.
To further decrease the standard deviation of the reconstructed phase and hence improve the phase sensitivity, a double biprism setup [3] was used. It allows the minimization of the applied biprism voltages. With the resulting increased stability of the holographic system, longer exposure times are possible, leading to an observed standard deviation of the reconstructed phase of about 2π/740 in a single empty hologram with an exposure time of 20 s.
The 80 kV off-axis electron holography is demonstrated using a thin GaN-foil oriented along [11-20] zone axis. Here, to avoid image resolution due to specimen drift, ten holograms with an exposure time of two seconds were made and averaged. The residual lens aberrations in the reconstructed image wave were corrected to retrieve the object exit wave. For comparison with the experiment, a GaN foil with thickness 1.9 nm, a chromatic aberration of 1.3 mm and an energy spread of 0.8 eV was simulated (fig. 3) [4]. Amplitude and phase of the object exit wave show a good match to the simulation.


  1. M. Lehmann et al., Microscopy and Microanalysis 8 (2002), p. 447-466.
  2. M. Linck et al., Microscopy and Microanalysis 18 (Suppl 2) (2012), p. 478.
  3. K. Harada et al., Applied Physics Letters 84 (2004), p. 3229.
  4. P. A. Stadelmann, Ultramicroscopy 21 (1987), 131.

Fig. 1: Damping by the modulation transfer function (MTF) of the camera as a function of the sampling frequency g at acceleration voltages of 80 kV and 300 kV.

Fig. 2: Standard deviation in the reconstructed phase of electron holograms with a sampling rate of 10 pixels per hologram fringe as a function of the inverse fringe spacing s related to the object plane. The exposure time was 2 s. A smaller standard deviation is equivalent to a higher phase sensitivity.

Fig. 3: Comparison of experimental and simulated object exit-wave of a thin GaN foil in [11-20]-orientation with a thickness of approximately 1.9 nm. Both amplitude and both phase images have the same grey levels.

Type of presentation: Poster

IT-11-P-2738 Nanoscale strain analysis by dark-field electron holography of shallow trench isolation structures for MOSFET technology

Ravaux F.1, Cherkashin N.2, Lolivier J.3, Alexander D. T.1
1Interdisciplinary Centre for Electron Microscopy (CIME), École Polytechnique Fédérale de Lausanne (EPFL), Lausanne, Switzerland, 2CEMES-CNRS, Toulouse, France, 3EM Microelectronic SA, Marin, Switzerland
duncan.alexander@epfl.ch

As microelectronic devices shrink, mechanical deformations induced by the various fabrication process steps play a growing role in their technological properties. These parameters must be understood and controlled because strained silicon can significantly boost transistor performance [1], but can also lead to product failures. Of interest in this study are transistors based on shallow trench isolation (STI) structures (Fig 1.) used in metal-oxide-semiconductor field effect transistor (MOSFET) technology. The main goal is to understand the mechanisms responsible for strain formation inside the active silicon area.
Off-axis dark field electron holography (DFEH) in TEM is used to perform the strain analysis on dedicated test structures that simulate the mechanical deformation on the active silicon areas that are supposed to receive the MOSFET. The choice of removing the MOSFET was made in order to analyze/evaluate purely the influence of the STI structure, within a systematic study of the influence of STI fabrication steps. The DFEH technique [2] is based on the interference of diffracted beams coming from two different regions of the sample with the aid of an electrostatic biprism (Fig. 2). If the two regions present a lattice constant difference, a phase difference will be measured from the holographic fringes and the strain information can be retrieved.
For the experiments, classical focused ion beam (FIB) TEM lamellae were prepared but thickness variations due to the STI geometry and lamella bending due to strain relaxation prevented the extraction of the strain tensor from the entire structure. For this reason, a dedicated sample preparation method has been developed to match the requirements imposed by the DFEH method (i.e. uniform 120 nm sample thickness with no specimen bending). Backside FIB milling [3] combined with an innovative “double-bar” rigid-thin TEM lamella geometry succeeds in avoiding these artefacts (Fig. 3). Measurements made on the HITACHI I2TEM at CEMES-CNRS (Toulouse, France) now permit the extraction of the complete sample-plane strain tensor in the active silicon area sandwiched by two STI structures (Fig. 4). The results are in accordance with values obtained by CBED and TCAD simulations in the literature [4].
[1] M. Cai et al., IEEE Trans. Electron Devices, vol. 57 (2010) 1706–1709
[2] M. Hÿtch et al., Nature, vol. 453 (2008) 1086–1089.
[3] J. Gazda et al., Microsc. Microanal., vol. 16 Supplement S2 (2010) 230–231.
[4] A. Steegen and K. Maex, Mater. Sci. Eng. R Rep., vol. 38 (2002) 1–53.


The authors acknowledge funding by the Swiss Commission for Technology and Innovation (CTI), Franc Fort project number 13496.1 PFFLE-NM, and ESTEEM2 for Transnational Access to the I2TEM at CEMES. Marco Cantoni at CIME is thanked for many useful discussions.

Fig. 1: Schematic figures of CMOS transistor (a) and STI structures (b).

Fig. 2: Principle of the dark-field electron holography (DFEH) technique.

Fig. 3: The double-bar “rigid-thin” TEM lamella geometry (a). SEM images of the thin part (b) and the rigid frame (c). Bright-field TEM image of the region of interest (d).

Fig. 4: Strain measurement of the active silicon area sandwiched by STIs.

Type of presentation: Poster

IT-11-P-2745 Electron Holography and Electron Energy Loss Spectroscopy for Studying Electrochemical Reactions at Electrode/Electrolyte Interfaces in a Lithium-Ion Battery

Hirayama T.1, Yamamoto K.1, Shimoyamada A.1, Yoshida R.1, Iriyama Y.2
1Nanostructures Research Laboratory, Japan Fine Ceramics Center, 2Department of Materials, Physics and Energy Engineering, Nagoya University
t-hirayama@jfcc.or.jp

It is essential that detailed knowledge about the electrochemical reactions near the interfaces between electrodes and electrolyte be obtained to find clues for the development of more efficient batteries [1]. For this purpose, measurement of electric potential distributions and direct observation of ion distributions are of fundamental importance. Here we report our results of observations of such phenomena in all-solid-state lithium-ion batteries.

Electron holography is a powerful technique to map electric potential distributions in working batteries [2,3], while electron energy loss spectroscopy (EELS) is an effective technique to directly detect lithium distributions. The two methods provide a powerful means of revealing the electrochemical reactions that occur at electrode/electrolyte interfaces.

Figure 1(a) shows a schematic of the model battery used in our experiments. Fig. 1(b) shows the electric potential distribution measured at the negative side by electron holography after the first charge-discharge cycle. The negative potential indicates that a negative electrode region was formed in situ during this cycle [2]. Fig. 2(a) and 2(b) respectively show a transmission electron micrograph of a region near the negative electrode and corresponding spectrum obtained from Spatially Resolved Electron Energy Loss Spectroscopy (SR-EELS) measurements for a specimen that had undergone 50 charge-discharge cycles. This spectrum indicates that the in situ negative electrode contains excess lithium.

In summary, we have successfully observed distributions of electric potentials and lithium ions at the negative electrode side of an all-solid-state lithium-ion battery. These microscopy techniques provide new insights into the electrochemical reactions that take place in all-solid-state batteries during cycling, and should aid the design of high-performance devices.

References

[1] M. Armand and J.-M. Tarascon, Nature 451, 652-657 (2008).

[2] K. Yamamoto, et al., Angew. Chem. Int. Ed. 49, 4414-4417 (2010).

[3] K. Yamamoto, et al., Electrochem. Commun. 20, 113-116 (2012).


The authors would like to thank Dr. Y. Sugita and Mr. K. Miyahara of Chubu Electric Power Co., Inc., and Dr C. Fisher of Japan Fine Ceramics Center for valuable discussions. SR-EELS measurements were performed as part of the RISING project of NEDO, Japan. We are grateful to Profs. T. Abe, Y. Uchimoto and Z. Ogumi of Kyoto University for their encouragement and useful suggestions.

Fig. 1: FIG. 1. (a) Schematic of the model battery sample used in the experiments. An area corresponding to that enclosed by the red rectangle was observed by electron holography. (b) Electric potential distribution obtained by electron holography.

Fig. 2: FIG. 2. Spatially Resolved Electron Energy Loss Spectroscopy (SR-EELS) measurements. (a) Transmission electron micrograph of the region at the negative electrode and electrolyte Li1+x+3zAlx(Ti,Ge)2-xSi3zP3-zO12 (LICGC) interface; (b) the corresponding spectrum obtained by SR-EELS.

Type of presentation: Poster

IT-11-P-2987 Distribution of electric field and mechanical vibrations of CdS nanocombs during field emission

Migunov V.1, Duchamp M.1, Liu R.3, Kamran M. A.3, Zuo B.3, Li Z.2, Farle M.2, Dunin-Borkowski R. E.1
1Ernst Ruska-Centre for Microscopy and Spectroscopy with Electrons and Peter Gruenberg Institute, Forschungszentrum Juelich, Juelich, Germany, 2Faculty of Physics and Center of Nanointegration (CeNIDE), University of Duisburg-Essen, Duisburg, Germany, 3Beijing Key Lab of Nanophotonics and Ultrafine Optical Systems, School of Physics, Beijing Institute of Technology, Beijing, China
v.migunov@fz-juelich.de

CdS comb-like nanostructures (Fig. 1a) were recently shown to have remarkable optical properties and possible applications as waveguides [1]. On the other hand, nano-cone-based field-emission guns (FEGs) have shown a strong improvement of the coherence compared to classical W tips [2] that is of particular interest for electron holography for example. As CdS-based nanostructures are also used in optically-induced field emission [3], it combination with wave guiding properties may allow their use in FEGs, in which the light that triggers the emission of electrons is focused onto a different part of the nanostructure to that from which the electrons are emitted. With this an increase of the local temperature at the electron emission point that results in increase of electron energy spread can be avoided. Here, we study the field emission properties of such nanocombs in situ in the transmission electron microscope (TEM) using off-axis electron holography combined with electrical biasing.

The experiments involved mounting a CdS nanocomb onto a W needle and bringing a movable W probe towards CdS comb in situ in the TEM (Fig. 1b). Bias voltages of between 0 and -140V were applied to the nanocomb. Off-axis electron holograms were acquired before and during field emission from the apex of one of the branches of the comb. The phase shift measured using electron holography is proportional to the projected electrostatic potential. The density of the measured equiphase lines shown in Fig. 2 facilitates identification of the region from which electrons are emitted. This information can be correlated directly with the morphology of the apex. We measured threshold bias voltages for field emission for different electrode configurations. Unexpectedly, the nanocomb was observed to vibrate when field emission was initiated, as shown in Fig. 3.

[1] Liu R. et al., Nano Letters, 2013, 13, 2997−3001

[2] Houdellier F. et al., 15th European Microscopy Congress Manchester Central, 2012

[3] Zhang J. et al., Sci. China Phys. Mech. Astron., 2011, 54 (11), 1963–1966


We thank Giulio Pozzi for fruitful discussions.

Fig. 1: a) Scanning electron micrograph of a CdS nanocomb deposited on a holy carbon grid. b) Bright-field TEM image showing the geometry of the field emission experiment.

Fig. 2: Equiphase contours recorded using electron holography for a bias voltage of -23 V between the apex of one tooth of the CdS nanocomb (bottom) and the W probe (top).

Fig. 3: Frames from a video recorded in a TEM, showing the end of a tooth of the CdS comb (bottom) before the onset of field emission at a bias voltage of -32 V (left) and during field emission at a bias voltage of -33 V with a current of 4 nA (right). The contrast of the tooth is blurred due to the onset of mechanical vibrations with MHz range frequency.

Type of presentation: Poster

IT-11-P-3018 Experimental electron holographic tomography of magnetic vector fields in nanoscale materials

Diehle P.1, Caron J.1, Kovacs A.1, Ungermann J.2, Kardynal B.3, Dunin-Borkowski R. E.1
1Ernst Ruska-Centre for Microscopy and Spectroscopy with Electrons and Peter Grünberg Institute 5, Forschungszentrum Jülich, Jülich, Germany, 2Institute of Energy and Climate Research, Forschungszentrum Jülich, Jülich, Germany, 3Institute of Semiconductor Nanoelectronics and Peter Grünberg Institute 9, Forschungszentrum Jülich, Jülich, Germany
p.diehle@fz-juelich.de

It is important to develop a characterization technique that can be used to measure three-dimensional magnetization distributions in nanoscale magnetic materials, which are of interest for a variety of applications that include future information storage technologies [1].

Off-axis electron holography allows the phase shift of the electron wave that has travelled through a thin sample to be measured in the transmission electron microscope (TEM) [2]. The phase shift is, in turn, sensitive to the in-plane component of the magnetic flux density within and around the specimen integrated in the electron beam direction. Although conventional tomographic reconstruction algorithms can in principle be used to reconstruct the magnetic flux density within and around a TEM specimen from two tilt series of magnetic phase images, significant artefacts can result from the difficulty of acquiring two ideal tilt series of images about independent axes over a complete range of specimen tilt angles.

We have therefore chosen to develop a different approach, which is based on the use of a model-based reconstruction algorithm to reconstruct the three-dimensional magnetization distribution in a TEM specimen, rather than the magnetic flux density, from two tilt series of phase images recorded using off-axis electron holography. Our approach involves repeated forward calculation of magnetic phase images until a best-fitting simulated magnetization distribution to the experimental images is obtained.

Practical challenges include the need to subtract the mean inner potential contribution to the phase shift recorded at each specimen tilt angle, the need to take into account the perturbed reference wave in the simulations, the minimization of contributions to the recorded phase from diffraction contrast and the preparation of a TEM specimen that is neither too thick nor obscured by another part of the specimen or the specimen holder when it is tilted to high angles about two axes.

Figure 1 shows simulated magnetic induction maps and magnetic phase images of an elliptical magnetic element for two different specimen tilt angles. Figure 2 illustrates the use of an electron-transparent silicon nitride membrane to acquire two independent tilt series of electron holograms of a lithographically patterned magnetic element. We will compare simulated and experimental results from both lithographically patterned elements and more three-dimensional nanoscale magnetic specimens and discuss the influence of noise and experimental artefacts on the final reconstructed magnetization distribution.

[1] S. S. P. Parkin et al., J. Appl. Phys. 85, 5828 (1999)
[2] D. Gabor, Proc. Roy. Soc. A 197, 454 (1949)


RDB acknowledges the European Commission for an Advanced Grant.

Fig. 1: Simulated magnetic phase shift and magnetic induction map of a uniformly magnetized elliptical element (saturation magnetisation = 3T) with semi-axis lengths of 600nm and 200nm and a thickness of 50nm for two different specimen tilt angles. The colors represent the direction and magnitude of the phase gradient, according to the color wheel shown.

Fig. 2: Illustration of the design of lithographically patterned magnetic elements that can be tilted to high angles about two independent axes.

Type of presentation: Poster

IT-11-P-3140 Electron Beam induced Currents in GaN p-n Junctions

Park J. B.1, Niermann T.1, Lehmann M.1
1Institut für Optik und Atomare Physik, Technische Universität Berlin, Straße des 17. Juni 135, D-10623 Berlin, Germany
park@physik.tu-berlin.de

Off-axis electron holography (EH) is a unique approach for analyzing potential variation down to nanometer scale. In absence of magnetic fields and in kinematical diffraction condition, the phase modulation is proportional to the potential (sum of mean inner potential and built-in potential Vpn) and the specimen thickness. However, quantitative measurement of Vpn in GaN shows a huge deviation with the theoretical value. Additionally, it seems that the hypothetical “dead layer” on both surfaces affected by focused ion beam (FIB) preparation is not the only reason for the discrepancy since the measured Vpn from the bulk (n-GaN)/shell (p-GaN) geometry of the GaN p-n junction is not affected by the FIB preparation [1].
In this study, we assess the influence of the electron beam on the measured Vpn at GaN p-n junctions. Two GaN p-n junctions of comparable dopant concentration were prepared by FIB. Thereby, a needle-like shape geometry is deliberately chosen to ensure that the whole specimen is illuminated during the experiment [2]. Moreover, a defined path of the induced current arising from electron-hole pair generation and secondary electron emission is provided. Low voltage (5 kV) is applied for the final polishing step to mitigate the surface damaging by FIB.
Fig. 1 shows the FIB milled needle with a sketch of the layer structure. In order to enhance the signal to noise ratio of the phase, series of holograms separately for each illumination intensity were recorded and subsequently averaged. The reconstructed phase depicts a clear phase jump at the p-n junction (Fig. 1 (c)).
Two electron beam induced effects are observed. Firstly, in Fig. 2 the measured Vpn increases with reducing the electron dose rate. This illumination dependency of Vpn can be quantitatively explained by a solar-cell model. Hence, Vpn = Vpn, expectedVbias is fitted on the data set, where Vpn, expected is determined by the doping concentration (sample A: 3.41 V, sample B: 3.43 V), Vbias is the voltage drop across the p-n junction, which is affected by illumination induced currents owing to secondary electron emission and e-h pair generation [3]. Additionally, fluctuation of Vpn is depicted by two needles, which are prepared from a same wafer and with the same FIB preparation steps. However, both needles (I and II) show the same behaviour of Vpn over the electron dose rate. Secondly, beam damage of the specimen is observed in Fig. 3. The Vpn diminishes over a period of time and might converge to a constant value, which is about the half of the initially measured Vpn.

1. S. Yazdi et al., journal of physics: Conf. Ser. 471 (2013), p.012041.

2. A. Lenk, Dissertation, Dresden (2008).

3. S.M. Sze in “Physics of Semiconductor Devices”, 2. Edition, John Wiley & Sons (1981).


This work is carried out within the framework of the DFG collaborative research center SFB787 semiconductor nanophotonics.

Fig. 1: (a) SEM image of the needle-like shape specimen after FIB preparation. (b) The layer structure of the specimen. (c) The reconstructed phase of (a). The phase jump at the p-n junction due to different doping type is clearly visible. The large-area undulation of the phase change along the specimen is caused by thickness variation.

Fig. 2: The measured built-in potential Vpn over the electron dose rate for two needle specimens of sample A (FIB needle I and II). The Vpn is enhanced over roughly 30% by reducing the electron dose rate by a magnitude of 2.

Fig. 3: The Vpn is measured directly after FIB preparation over a period of time (about 1 hour) at sample B. We observe a decrease of Vpn by a factor 2 over time.

Type of presentation: Poster

IT-11-P-3191 Hybridization of Off-Axis and Inline High-Resolution Electron Holography

Ozsoy Keskinbora C.1, Boothroyd C. B.2, Dunin-Borkowski R. E.2, van Aken P. A.1, Koch C. T.3
1Max Planck Institute for Intelligent Systems, Stuttgart, Germany, 2Ernst Ruska-Centre for Microscopy and Spectroscopy with Electrons and Peter Grünberg Institute, Forschungszentrum Jülich, Jülich, Germany, 3Institute for Experimental Physics, Ulm University, Ulm, Germany
ozsoy@is.mpg.de

In conventional TEM experiments, only the intensity (i.e., the square of the amplitude) of the wave function can be measured. The phase information gets lost when the electron is detected by the CCD camera. Denis Gabor introduced an approach that could be used to solve this problem 66 years ago [1]. In Gabor’s original setup, which is the pioneering scheme for inline holography, the wave that has been scattered by the specimen (the object wave) interferes with a reference wave propagated along the same axis. Using laser light, Leith and Upatnieks [2] showed that separation of the axes of propagation of the reference and object waves could be used to solve the twin-image problem. Möllenstedt later translated this idea back to electron microscopy, creating the field of off-axis electron holography [3,4]. Inline electron holography, or focal series reconstruction, is now a common method in high-resolution TEM. Although it is very efficient for recovering high spatial frequency variations in phase, it is inefficient for recovering phase information at low spatial frequencies. In contrast, high-resolution studies are very challenging for off-axis holography, because the interference fringes must be at least twice as fine as the finest feature of interest in the object to be resolved.
In this study, we present a new approach that combines off-axis and inline holography and allows reliable phase information to be recovered for all spatial frequencies. For a desired signal-to-noise ratio, the required total exposure time is lower than that for traditional high-resolution off-axis electron holography.
All holographic data were acquired using round illumination with a FEI Titan TEM operated at 300 kV and using a bi-prism voltage of 97.4 V for off-axis electron holography. Figure 1 shows phase and amplitude images of a gold particle obtained using inline and off-axis electron holography and the hybrid method, respectively, for a total exposure time of 7s. Although the noise level in the vacuum region is slightly higher than for inline electron holography (0.055π vs. 0.046π), the recovery of low spatial frequencies is far better than for inline holography alone. Such noise levels are difficult to achieve using off-axis holography for the exposure time utilized here.
[1] D. Gabor, Nature vol. 161 (1948), p. 777–778.
[2] E.N. Leith, J. Upatnieks, J. Opt. Soc. Am. vol 52 (1962), p. 1123–1130.
[3] G. Möllenstedt and H. Düker, Naturwissenschaften vol. 42 (1955), p. 41–41
[4] G. Möllenstedt and H. Wahl, Naturwissenschaften vol. 55 (1968), p. 340–341


The authors thank to; John Bonevich for offering free public use of HolograFREE reconstruction software. Wilfried Sigle, Luis M. Liz-Marzan and Cristina Fernandez-Lopez for samples. The research leading to these results received funding from the European Union Seventh Framework Programme [FP7/2007-2013] under grant agreement no312483 (ESTEEM2) and the Carl Zeiss Foundation.

Fig. 1: a)-c) Reconstructed phase, d-f) reconstructed amplitude images from inline, off-axis and hybrid methods, respectively.

Type of presentation: Poster

IT-11-P-3505 Electron holography for magnetic and electric in situ imaging

Ponce A.1, Cantu-Valle J.1, Diaz Barriga E.2, Luna C.2, Mendoza-Santoyo F.1, Jose Yacaman M.1, Eder J.1
1University of Texas at San Antonio, 2Universidad Autónoma de Nuevo León
arturo.ponce@utsa.edu

In this work we report the local magnetic behavior of multi-segmented Cox-Ni1-x nanowires by off-axis electron holography as well as electric contribution in ZnO nanostructures. The nanowires were grown by electrode deposition, by alternating cycles that produce the multi-segmented structure. The crystalline phase of each segment and magneto-crystalline anisotropy will be resolved by the phase maps obtained by electron holography.
Samples were studied using JEOL JEM ARM-200F. Holograms were obtained under two different conditions; first following the dual-lens imaging system, using a voltage in the objective lens of 1V, and secondly under Lorentz mode, with the objective lens turned off.
The ZnO nanorods were studied under external bias applied to the sample in situ TEM. The experimental setup consists in a connection from the holder (Nanofactory electrical holder) to the external source. The holograms have been live recorded and the electric variation is observed in the reconstructed phase.
In order to select the optimum parameters for the holograms reconstruction, the fringe spacing, interference width and fringe contrast were measured for different biprism voltages. The holograms were recorded in-focus with a biprism voltage of about 20-23V and interference fringe spacing of 5nm and a fringe contrast of 22%. The quality of the holograms will depend in a high fringe contrast and number of electrons on the holograms, which imply better signal to noise ratio, this will be reflected in the reconstructed phase and amplitude images. The holograms acquisition was obtained using specialized software from Gatan, Digital Micrograph (DM) and processed by the latest version of HoloWorks, which includes a new feature to extract the magnetic induction and magnetic contours from a phase image.


This project was supported by grants from the National Center for Research Resources (5 G12RR013646-12) and the National Institute on Minority Health and Health Disparities (G12MD007591) from the National Institutes of Health. The authors would also like to acknowledge the NSF PREM # DMR 0934218. A special recognition to Holowerks LLC., and Dr. Edgar Voelkl for their guidance and support.

IT-12. Surface microscopy and spectroscopy

Type of presentation: Invited

IT-12-IN-2866 Addressing Fundamental Problems in Information technology: Opportunities for X-Ray Photoelectron Spectromicroscopy

Schneider C. M.1,2
1Peter Gruenberg Institute PGI-6, Research Centre Juelich, D-52425 Juelich, Germany, 2Faculty of Physics and CENIDE, University Duisburg-Essen, D-47057 Duisburg, Germany
c.m.schneider@fz-juelich.de

Modern information technology must exploit the full potential of complex material systems for the meticulous control of state variables. These state variables are used to encode an information bit and may be electron charges in semiconductor nanoelectronics, electron spins in the case spintronics, or local redox configurations in resistive switching elements. Consequently, the materials encompass intermetallic compounds, oxides or chalcogenides, elementary and compound semiconductors or even molecular components. In addition, the functional elements, for example, individual memory cells or transistor structures often involve nanometric dimensions and operate on nanosecond timescales or even below. This imposes considerable challenges on the characterization of electronic, chemical and magnetic states in the steady state or during operation.

Immersion lens microscopy with synchrotron radiation has matured into a versatile and powerful tool to investigate a broad range of issues in condensed matter physics and materials science. It combines high-resolution imaging with spectroscopic capabilities in a unique fashion. The excitation with photons from the soft to the hard x-ray regime ensures element selectivity and variable information depth. The polarization state of the synchrotron radiation enables a distinction of different magnetic orderings (Fig. 1), whereas the intrinsic time structure of the synchrotron radiation permits the study of processes with picosecond time-resolution.

In this contribution we will review the present status of x-ray photoemission spectromicroscopy with emphasis on applications in information technology. In particular, we will cover model systems in spintronics and in resistive switching (Fig. 2). The results will cover both static properties and dynamic processes. We will also discuss new developments, such as photoemission microscopy with hard x-rays and imaging spin polarimetry.


I would like to thank N. Barrett, S. Cramm, R. Dittmann, W. Drube, M. Escher, V. Feyer, A. Gloskovskii, A. Kaiser, J. Kirschner, A. Koehl, I. Krug, Ch. Lenser, M. Merkel, M. Patt, L. Plucinski, J. Rault, O. Renault, Ch. Tusche, N. Weber, R. Waser, and C. Wiemann for their cooperation. Financial support through the Deutsche Forschungsgemeinschaft (SFB 917) is gratefully acknowledged.

Fig. 1: Ferro- (left) and antiferromagnetic domain pat- terns in the system NiO/Fe3O4(110) exploiting XMCD (Fe) and XMLD contrast (Ni) at the indicated photon energies. Wide arrows indicate the local spin alignment axis.

Fig. 2: Spatially resolved hard x-ray photoemission from a Fe:SrTiO3. The dark squares result from 7nm thick Au electrode pads deposited for the resistive switching experiments.

Type of presentation: Invited

IT-12-IN-3176 Towards 1.5λ resolution with low energy electrons

Tromp R. M.1,2
1IBM T.J. Watson Research Center, Yorktown Heights, NY, 2Kamerlingh Onnes Laboratory, Leiden University, The Netherlands
rtromp@us.ibm.com

The highest resolution aberration-corrected electron microscopes today, operating at 300 keV, achieve a spatial resolution of 50 pm, or about 25 times the wavelength of the electron, l. With the third-order spherical aberration of the objective lens compensated, this resolution is limited by chromatic aberration, fifth-order spherical aberration, parasitic aberrations, and various microscope instabilities. On the other end of the spectrum, Low Energy Electron Microscopy (LEEM) without aberration correction has achieved a spatial resolution of 4 nm at 3.5 eV, or about 6l. The resolution in such instruments is primarily limited by the spherical and chromatic aberrations of the uniform electrostatic field between sample and cathode objective lens. This uniform field is the first (virtual) image-forming element of the microscope. When used as a Photo Electron Emission Microscope (PEEM) resolution is usually limited to the range of 10-20 nm, depending on the details of the imaging conditions. In LEEM/PEEM aberration coefficients are strongly energy-dependent, and must be readily adjustable even in a single experiment, so as to track the aberrations as they change with electron energy.

Over the last several years we have developed an aberration-corrected LEEM/PEEM instrument[1], using a relatively simple catadioptric (i.e. electrostatic lens + mirror) correction system which provides independent control over the lowest order spherical and chromatic aberration coefficients, and the focal length of the correction optics. We have demonstrated the practical feasibility of aberration correction in LEEM/PEEM, achieving spatial resolution below 2 nm for the first time [2]. Detailed studies of the wave-optical image formation process show that resolution well below 1 nm is possible in principle [3].
In this talk I will review challenges and recent progress towards reaching the goal of a spatial resolution of just 1.5 times the wavelength of the electron in LEEM.

1. A new aberration-corrected, energy-filtered LEEM/PEEM instrument. I. Principles and design, R.M. Tromp, J.B. Hannon, A.W. Ellis, W. Wan, A. Berghaus, O. Schaff; Ultramicroscopy 110 (2010) 852

2. A new aberration-corrected, energy-filtered LEEM/PEEM instrument. II. Operation and results; R.M. Tromp, J.B. Hannon, W. Wan, A. Berghaus, O. Schaff; Ultramicroscopy 127 (2013) 25-39

3. A Contrast Transfer Function approach for image calculations in standard and aberration-corrected LEEM and PEEM, S.M. Schramm, A.B. Pang, M.S. Altman, R.M. Tromp; Ultramicroscopy 115 (2012) 88-108


Type of presentation: Oral

IT-12-O-1893 Secondary electron quasi-simultaneous observation of energy selective imaging and diffraction in DualEEM

Grzelakowski K. P.1
1OPTICON Nanotechnology, Wroclaw, Poland
k.grzelakowski@opticon-nanotechnology.com

We present the first results of the AES application in the novel technique based on PEEM [1] and LEEM [2] concepts: DualEEM [3].It utilizes the idea of the imaging α- Spherical Deflector Analyzer (α-SDA) [4] with the total deflection 2π.The image returns exactly to its origin on the optical axis independently of the starting angle and energy.As a consequence,the object of filtration and its image are invariant in the 2π deflection process. Additionally,the final angles of incidence at this plane change the sign after the full angle deflection,which indicates mirroring-like effect.This mathematical analogy to the classical mirror operator is further enriched by the unique property of the α-SDA analyzer: the direction of electron propagation before “reflection” at the symmetry plane is preserved after the 2π deflection process is completed.Therefore, contrary to the classical electrostatic mirror,the propagation direction on both sides of the mirror plane is preserved.This could be referred to as a unique “through the looking-glass” electron optical effect.Thus, the α-SDA imaging analyzer exhibits all the advantages of the electrostatic mirror without the loss of beneficial linear geometry.The α-SDA assures not only the selection of characteristic energy for imaging, but also a beam separation into two imaging channels: energy-selective real image and reciprocal (diffraction) image and their quasi-simultaneous acquisition.The microscope is equipped with an Auger electron gun located inside the immersion objective lens that allows for an unique electron beam sample illumination and thus,opens a new application field for electron spectromicroscopy under laboratory conditions.For the first time that unique kind of the sample illumination is used for the energy selective Auger electron imaging and diffraction. Both are visualized at two independent imaging channels:one for the real and the other for the reciprocal image.These images are acquired quasi-simultaneousely through software based switching of on and off potentials of the one of hemispheres of the α- spherical deflector analyzer.The first results are reported and discussed.                                                                       

1 E. Brueche, Z.Phys. 86  (1933) 448,

2 E. Bauer, in Proc.of the 5th Int. Congr.for El.Micr.,(Academic, N.Y., 1962, p.D-11)

3 K.P. Grzelakowski, Ultramicroscopy, 130 (2013) 29; 4 Ultramicroscopy 116 (2012) 95


The author acknowledges the financial support by the NCBR in Warsaw. My thanks are also due to Krzysztof Wojcik and his team at “Metob” for the excellent machining and advice. I am very grateful to Prof.Ernst Bauer for his valuable suggestions and discussions. I would like also to express my gratitude to Janusz Krajniak for his dedication to this project and Dariusz Mirecki for his support.

Fig. 1: Black and blue areas indicate α-rays and γ-rays, respectively, p1 and p2 denote the symmetry and diffraction planes of the α-SDA, respectively: (a) energy selective k-projection, upper hemisphere switched off, (b) energy selective real image mode, lower and upper hemisphere switched on. PEEM mode: α-SDA switched off (right hand part of Fig.b).

Type of presentation: Oral

IT-12-O-2156 Novel development of very high brightness and highly spin-polarized electron gun with a compact 3D spin manipulator for SPLEEM

Koshikawa T.1, Yasue T.1, Suzuki M.1, Tsuno K.1, Goto S.2, Jin X.3, Takeda Y.4
1Osaka Electro-Communication Univ. , Osaka, Japan, 2San-yu Electronic Corp. , Tokyo, Japan, 3School of Engineering, Nagoya Univ., Nagoya, Japan , 4Aichi Synchrotron Radiation Center, Seto, Japan
kosikawa@isc.osakac.ac.jp

We have already developed a novel very high brightness and high spin-polarized low energy electron microscope (SPLEEM) and applied it to clarify the magnetic property of [CoNix]y/W(110) and Au/CoNi2/W(110) during growth of ultra thin films[1-5]. Such thin multi layered films are important for current-driven domain-wall-motion devices. Our developed SPLEEM can make us the dynamic observation of the magnetic domain images possible. However the size of the spin-polarized electron gun is large and we have started to develop a new compact spin-polarized electron gun with a novel idea. In principle two devices are necessary to operate 3-dimensional spin direction; one is a spin manipulator which changes the out-of-plain spin direction and another one is a spin rotator which can change the in-plain spin direction. We have proposed a multi-pole Wien filter which enables 3-dimensional spin operation with one device. The developed 3D multi-pole spin manipulator which has 8 poles in the present development and the magnetic and electric field in the multi-poles Wien filter as shown in Fig.1. Uniform field can be obtained at the center part of the Wien filter with 8 poles and 12 poles, however 4 poles filter gives non-uniform field even at the center. In the present development 8 poles Wien filter has been adopted. The results of magnetic images and asymmetries of Co(4ML)/W(110) vs. polar and azimuthal angles are shown in Fig.2. The results clearly show that spin direction can be operated three dimensionally with one device.
1)  X.G. Jin et al., Appl. Phys. Express 1, 045002 (2008).
2)  N. Yamamoto et al., J. Appl. Phys.103, 064905 (2008).
3)  M.Suzuki et al., Appl. Phys.Express 3, 026601 (2010).
4)  M.Suzuki et. al., J.Phys. Cond. Matt. 25, 406001-1-8 (2013) (Short News on IOP web and IOP select).
5)  K. Kudo et al., J.Phys. Cond.Matte. 25, 395005-1-6 (2013).


This work was supported by a Grand-in-Aid for Scientific Research (A) (Grand No. 23246015) from the Japan Scoiety for the Promotion of Science (JSPS) and the System Development Program for Advanced Measurement ans Analysis from Japan Science and Technology (JST).

Fig. 1: 3D multi-pole spin manipulator and uniformity of magnetic and electricField.

Fig. 2: The magnetic images and the asymmetries vs. the polar and azimuthal angles for Co(4ML)/W(110).

Type of presentation: Oral

IT-12-O-2690 Time-of-Flight Momentum Microscopy with Imaging Spinfilter: Dirac-Type States on Clean and Oxygen-Covered W(110) and Mo(110)

Medjanik K.1, Schönhense G.1, Chernov S.1, Schertz F.1, Nepijko S. A.1, Elmers H. J.1, Oelsner A.2, Tusche C.3, Kirschner J.3
1Institute of Physics, Johannes Gutenberg-University Mainz, Germany , 2Surface Concept GmbH Mainz, Germany, 3Max Planck Institute for Microstructure Physics, Halle, Germany
medyanyk@uni-mainz.de

We present a novel method for k-space mapping of electronic bands with utmost efficiency. The instrument combines the k-imaging properties of a cathode-lens microscope with the superior resolution of ToF spectroscopy and the parallel acquisition capability of ToF-PEEM [1]. For the first experiments a frequency-doubled Ti-sapphire laser was used for excitation by two-photon photoemission (2hv = 5.8 – 6.6 eV). A delay-line detector serves for rapid single-event counting with 150 ps time resolution and 10 Mcps maximum count rate. The dispersion behavior of the bands is observed in a 3D (kx,ky,E) matrix as schematically depicted in Fig.1, which is confined by the photoemission horizon (condition k_I_=0) in the shape of a E-kII paraboloid. The kII- and energy-range are presently limited by the low excitation energy. Mapping a complete data set with good statistics requires only few minutes of acquisition time. An integral Ir-based imaging spinfilter [2] yields spin resolved 3D-maps.

The low excitation energy is well suited to study surface states close to the centre of the SBZ. As first system we chose the highly anisotropic Dirac-type surface state recently discovered on W(110) [3] and the analogous state on Mo(110). These states arise in a pocket-shaped partial bandgap region, nevertheless the existence of Dirac states on metals was very surprising [3,4]. Fig. 2 shows kx-ky sections at EF (a,b) and E-kII sections displaying the crossover points for clean W(110) at about EB=1.25eV (c) and at 0.6eV for oxygen-covered Mo(110) (d). Hole doping by oxygen shifts the Dirac state to lower binding energy; in addition we found a pronounced pattern of very similar surface states on W(110)-O(1x1) and Mo(110)-O. No bulk bands are visible in this region of k-space. In all cases the Dirac state is highly anisotropic (2mm symmetry), revealing massless behavior (i.e. linear band dispersion) along one mirror plane and a high effective mass (flat band region) along the second mirror plane (g,h), similar as measured and calculated for clean W(110) [3,4]. For the oxidic surfaces we observe a complex 3D k-space behavior that is hard to elucidate in conventional ARPES. Using s-, p- and circular polarization we probe orbital symmetries and hybridization effects of the band states. The dichroism pattern in Fig. 2(e) is antisymmetric with respect to all mirror planes and can be understood in terms of a simple dz2-orbital model.

[1] G. Schönhense et al., Surf. Science 480 (2001) 180;

[2] C. Tusche et al., Appl. Phys. Lett. 99 (2011) 032505; D. Kutnyakhov et al., Ultramicroscopy 130 (2013) 63

[3] K. Miyamoto et al., Phys. Rev. Lett. 108 (2012) 066808;

[4] H. Mirhosseini et al., New J. of Phys. 15 (2013) 033019.


Project funded by BMBF (05K12UM2 and 05K12EF1).

Fig. 1: Scheme of the experiment; the 3D (kx,ky,E) time-resolving single-electron counting detector registers each electron within the E-kII paraboloid with a maximum count rate of 107 counts per second.

Fig. 2: Sections through the (kx,ky,E) matrix (at 2hv=6.6 eV). a,b: Fermi surfaces of W(110) and Mo(110)-O, respectively; c,d corresponding E-kII sections. e: circular dichroism in the region above the Dirac point of clean W(110), g: anisotropic shape of the Dirac point; f,h: calculated patterns for EB=0.9 and 1.2 eV [4].

Type of presentation: Oral

IT-12-O-2910 Shadow Dark-Field LEEM and Scanning Micro-LEED of Epitaxial Graphene on Ru(0001) and Ir(111) Surfaces

Yu K. M.1, Man K. L.1, Altman M. S.1
1Hong Kong University of Science and Technology, Hong Kong, China
phaltman@ust.hk

Spatially resolved measurements using cathode lens microscopies have made notable contributions to the understanding of graphene layers that are customarily spatially inhomogeneous [1]. We have applied low energy electron microscopy (LEEM) and complementary micro-low energy electron diffraction (μLEED) to study the structure and morphology of single layer graphene (g) on Ru(0001) and Ir(111) surfaces, examples of strongly and weakly interacting substrates, respectively. Our investigations of g/Ru(0001) reveal rich structural non-uniformity that depends strongly on preparation conditions. When the g/Ru(0001) layer is prepared using chemical vapor deposition (CVD) by exposure to ethylene at high temperature, we observe strong streaking of superstructure diffraction spots (Fig. 1(a),(c)) for large area (3μm) illumination. This indicates the proliferation of small angle (<0.25°) lattice rotations in the graphene layer. Corresponding small-angle lattice rotational domains are visualized in “shadow” dark field LEEM images (Fig. 1(b),(d)) that are formed by selecting the rotated edge of the streaked superstructure spots using the contrast aperture. The presence of rotation domains also causes the sharp sets of μLEED superstructure diffraction spots around each integer order spot to rotate to-and-fro as a group about their respective stationary foci when the small μLEED illumination beam (250nm) is scanned across the surface. These scanning μLEED measurements provide detailed information about the rotation angle distribution and even indicate a net deviation from perfect alignment between graphene and substrate.


Although the length scale of the rotational domains in g/Ru(0001) can be pushed up to the sub-micron length scale by increasing the growth temperature (Fig. 1(d)), further improvements are limited by the diminishing growth rate at increasingly higher temperature for accessible ethylene pressure. On the other hand, massive single rotation domains can be fabricated by CVD if the crystal is first pre-loaded with carbon by dissolution of a single graphene layer. Although near perfect rotational order was observed (Fig. 1(e),(f)), scanning μLEED revealed substantial spatial variation of lateral periodicity that was not seen when small-angle rotational domains were present. Hence, fabrication of optimal uniform g/Ru(0001) is still elusive. Experiments on g/Ir(111) reveal that small angle rotational microstructure is similarly prevalent when the graphene lattice is nominally aligned with the substrate, but it is substantially suppressed in macro-domains with larger misalignments.

Reference
[1] K.L. Man and M.S. Altman, J. Phys.: Condens. Matter 24, 314209 (2012).


Financial support from the Hong Kong Research Grants Council under Grant No. HKUST600113 is gratefully acknowledged.

Fig. 1: (a),(c),(e) LEED patterns obtained from the areas indicated in (b),(d),(f) shadow dark-field LEEM images of g/Ru(0001). Graphene was prepared by CVD at (a),(b) 1100K; (c),(d) 1270K; (e),(f) 1300K on a preloaded substrate. Contrast fine structure in (b),(d) is due to small angle rotation domains. A uniform rotation domain is seen in (f).

Type of presentation: Poster

IT-12-P-3198 Detecting the topographic, chemical and magnetic contrasts with nanometer spatial resolution

Zanin D. A.1, Erbudak M.1, De Pietro L. G.1, Cabrera H.1, Kostanyan A.1, Vindigni A.1, Pescia D.1, Ramsperger U.1
1Laboratory for Solid State Physics, ETH Zürich, Switzerland
dzanin@phys.ethz.ch

During the last decades, magneto-imaging techniques based on the analysis of secondary electrons helped the discovery of many interesting phenomena related to magnetic-domain patterns, such as re-entrant topological transitions. For those studies, a typical spatial resolution of some tens of nm, achieved e.g. in Scanning-Electron-Microscopy with Polarization Analysis (SEMPA), was more than enough. Nowadays, the quest to resolve magnetic textures in direct space at atomic scale is triggered by novel fundamental and applicative issues. Domain walls, in relation to their potential use in spintronic devices, represent one example. Inspired by the Russel Young topografiner we redesigned the SEMPA setup by replacing the primary electron beam source and the probing method. We dubbed this new technique Near Field-Emission Scanning Electron Microscopy (NFESEM). In NFESEM the sample surface is typically investigated by scanning at constant height with a primary electron beam energy in the range between 20eV and 100eV. A suitable detector analyzes secondary electrons scattered by the surface. We present the resolution improvement on topographic mapping of Fe-patches evaporated on W(110) substrate (Figure 1) and advances in energy analysis of secondary electrons (Figure 2). Moreover, we report on recent efforts to endow NFESEM with the polarization analysis of the detected secondary electrons that emphasize the true potential of this new technique. In particular, the characteristic spatial resolution and the sizeable secondary electrons yield (see Figures 1 and 2) support the technical feasibility of electron spectroscopy and magnetic-domain mapping at nanometer scale with NFESEM.


We thank Andreas Fognini, Thomas Michlmayr and Yves Acreman for the scientific support, Thomas Bähler for technical assistance and the Swiss National Science Foundation and ETH Zurich for financial support.

Fig. 1: (Left) STM Map of 0.4 atomic layers of FE on stepped W(110), showing atomic Fe-patches (bright) residing on the terraces and decorating the steps. (Right) The same surface spot recorded in NFESEM mode. Although the Fe-patches are on top of the W-substrate they appear darker - both the patches on the terraces and along the steps.

Fig. 2: (Left) Energy spectrum of a GaAs(110) surface for a tip-sample distance of 100 nm, both the secondary electron cascade an the elastic peak are clearly distinguishable. (Top right) Map of a GaAs(110) decorated surface produced by secondary electrons with 13 eV energy for a tip-sample distance of 12 nm. (Bottom right) STM reference image.

IT-13. Focused ion beam microscopy and techniques

Type of presentation: Invited

IT-13-IN-2711 Advances in FIB Nanotomography

Cantoni M.1, Knott G. W.1, Burdet P.2
1EPFL-CIME,Lausanne, Switzerland, 2Department of Materials Science and Metallurgy, EM Group University of Cambridge, United Kingdom
marco.cantoni@epfl.ch

FIB-tomography is used in materials science for 3D-analysis of nanostructured materials [1]and in life science for the analysis of complex structures like brain tissue [2]. This presentation summarizes recent technological improvements, which include advancements in detector technology for electron imaging and elemental analysis, scan generator technology for high throughput imaging, and automated drift correction for reliable 3D reconstruction. New in-column detectors have a higher sensitivity for low energy electrons, which is the basis for a very high resolution down to a few nm voxel size. The low kV imaging can be combined with energy filtering in order to detect a pure signal of backscattered electrons (BSE), which improves the reliability of phase segmentation and quantitative analysis. The quality of the 3D reconstructions can also be improved with refined procedures for drift correction based on reference marks. In addition, with the new scan generators image acquisition and ion milling can be performed synchronously. In this way the acquisition speed increases further. Finally, spectral and elemental mapping (XEDS) based on Silicon Drift Detectors (SDD) provides higher X-ray count rates. Increased acquisition rates open new possibilities in chemical analysis that provide larger data cubes with higher representativeness. The new possibilities of FIB-tomography are illustrated with the following examples: a) Reliable phase segmentation is discussed for a superconducting material with trapped pores that cannot be filled with resin. b) Combined analysis of SE- and BSE stacks reveals the complex microstructure of a Sn-solder with different nano-sized precipitates [3] and c) High throughput elemental analysis is performed of a NiTi stainless steel with a complicated multi-phase microstructure [4]. The examples document the recent advancements in resolution, contrast, stability and throughput, which are necessary for reliable and representative 3D-analysis.

References
1. L. Holzer, M. Cantoni, in Nanofabrication Using Focused Ion and Electron Beams—Principles and Applications, I. Utke, S. Moshkalev, P. Russell, Eds. (Oxford University Press, New York, 2012), pp. 410–435.
2. M. Cantoni, C. Genoud, C. Hébert and Graham Knott, Microsc. & Anal. 24(4): 13-16 (2010)2010.
3. M. Maleki, J. Cugnoni, J. Botsis, Acta Mater. 61 (1), (2013).
4. P. Burdet, J. Vannod, A. Hessler-Wyser, M. Rappaz, M. Cantoni, Acta Mater. 61 (8), 3090 (2013).


Fig. 1: Fig 1. Three-dimensional representation of the two different intermetallic phases in the SnCuAg-type solder, segmented based on simultaneously acquired secondary electron and backscattered electron image stacks.

Fig. 2: Fig 2. Complex chemical microstructure of a NiTi—stainless-steel weld with different phases. The yellow and red phases are chemically very close and required segmentation based on the secondary electron image contrast.

Type of presentation: Invited

IT-13-IN-5748 Development of cryogenic FIB-SEM based processes for organic and inorganic samples

Antoniou N.1
1ReVera Inc., Santa Clara, USA
nicholas@cns.fas.harvard.edu

Development of cryogenic FIB-SEM based processes Cryogenic EM is one of the best techniques we have to fix in place samples for EM that would otherwise be destroyed in the vacuum system [1]. Even though FIB-SEM has enabled a rapid, site specific and relatively easy way for TEM sample preparation it has not been easily adopted for cryogenically prepared samples. The impetus to develop a cryogenic sample preparation process with all the advantages of the room temperature one was high. However, the liftout and attach steps of the processes do not work in a cryogenic environment so either a more limited process had to be developed [2] or new equipment and new processes developed. Both have been achieved and we present here on the latter. This technique replicates the room temperature process but in cryo so that all of the developments surrounding FIB-SEM sample prep are available at cryogenic temperature. This not only requires a cold stage in the instrument but also sample transfer capability within and in-between each instrument involved. We describe this process, the equipment and modifications to it as well as applications. It was also discovered that certain inorganic material could benefit from cryogenic processing in FIB-SEM but for completely different reasons. Most FIB systems use gallium liquid metal as the ion sources (LMIS) and this species (Ga) can react chemically with certain compounds such as semiconductors to form undesirable side effects such as spheres and dots [3]. Milling under cryogenic conditions reduces these undesired reactions and in some cases this was the only way we could prepare the sample for TEM imaging (InN for example). We further investigated the effect of warm up after cryo milling for inorganic material and found that in most cases the undesirable side effects were minimized enough so that they did not interfere with our ability to image the sample. References [1] Adrian M., et al., “Cryo-Electron Microscopy of Viruses,” Nature 308, (01 March 1984), pp 32 - 36 [2] Rigort, A., et al., “Focused Ion Beam Micromachining of Eukaryotic Cells for Cryoelectron Tomography,” PNAS, (March 20, 2012), Vol. 109, No. 12, pp. 4449-4454. [3] Grossklaus, K. A., “Mechanisms of Nanodot Formation Under Focused Ion Beam Irradiation in Compound Semiconductors,” Journal of Applied Physics, Vol. 109 , Issue: 1, 2011, pp. 014319 - 014319-11.


The author acknowledges funding from the Center for Nanoscale Systems (CNS), a member of the National Nanotechnology Infrastructure Network (NNIN), which is supported by the NSF (under award no. ECS-0335765). CNS is part of Harvard University. I would also like to thank Dr. Ilan Shalish and Cheryl Hartfileld for their numerous contributions to this work.

Fig. 1: InN nanoparticles grown in a forrest and felled for easier pickup.

Fig. 2: A tungsten tip was sharpened to about 10 nm and used to pick up a single InN nanoarticle using only natural forces (Van der Waals). After pickup, the particle was placed on a TEM grid for milling in cryogenic conditions.

Fig. 3: GaN sample milled at room temperature using 1.5 nA FIB probe. The formation of dropplets is observed at room temperature and impedes clean milling of GaN.

Fig. 4: GaN sample milled at -145 C using 1.5 nA FIB probe as in Fig. 3. The dropplets did not form as in the room temperature milling.

Type of presentation: Oral

IT-13-O-2039 FIB-SEM Tomography of biological samples: Strategy of preparation to resolve high-resolution 3-D volumes

Kizilyaprak C.1, Daraspe J.1, Longo G.2, Humbel B. M.1
1University of Lausanne, Electron Microscopy Facility, Lausanne, Switzerland. , 2EPFL, Laboratory of Physics of Living Matter, Lausanne, Switzerland.
caroline.kizilyaprak@unil.ch

Elucidating the three-dimensional (3-D) spatial distribution of organelles within cells is essential for investigating numerous cellular processes. Tomography in the transmission electron microscope (TEM) is the method of choice for 3-D imaging of cellular structures down to 3nm resolution [1]. However, TEM tomography is typically limited to 500 nm thick sections making the reconstruction of an entire eukaryotic cell very challenging [2]. There is a need for a technology that can be used for rapid 3-D imaging of large mammalian cells to provide information at nanometre resolution. The most promising technology at the moment is the FIB-SEM tomography of fixed biological samples embedded in resin [3-7]. A FIB-SEM microscope is a scanning electron microscope combined with a focused ion beam (FIB) such that both beams coincide at their focal points. This combination enables bulk resin samples to be locally sectioned by ion milling, producing new block face imaged with the electron beam. This process can be repeated allowing 3-D analysis of relatively large volumes with a field of view of several micrometres.

Any fixed and embedded resin samples prepared for TEM examination can be used for FIB-SEM tomography. However, considerations have to be given to artefacts and surface damages induced by FIB milling and imaging [8]. In this study, different protocols of sample fixation and staining were explored in order to improve the signal/noise ratio, preserve the ultra-structure and reduce charging effects of biological samples. In addition, the behaviour of specific resin formulations [9] was investigated in the FIB-SEM microscope. The milling rate was measured and the damages caused by the ion impact on the resin were analysed. The most stable resin was used to improve the milling conditions. Finally, the geometry of the sample was optimized to improve the imaging conditions using detection of the backscattered electrons with the through-the-lens detector (BSE-TLD).

In conclusion, we propose a sample preparation and imaging strategy for high-resolution FIB-SEM tomography (Figure 1).

References:

1. Baumeister,et al.,Trends in cell biology, 1999.9(2): p.81-5.

2. Noske, A.B., et al.,Journal of structural biology, 2008.61(3): p.298-313.

3. Heymann, J.A., et al.,Journal of structural biology, 2009.166(1): p.1-7.

4. Knott, G., et al.,Journal of visualized experiments : JoVE, 2011(53): p.e2588.

5. Bushby, A.J., et al.,Nature Protocols, 2011. 6(6): p.845-58.

6. Villinger, C., et al.,Histochemistry and cell biology, 2012.138(4): p.549-56.

7. Wei, D., et al.,BioTechniques, 2012. 53(1): p.41-8.

8. Drobne, D., et al.,Microscopy research and technique, 2007.70(10): p.895-903.

9. Luft, J.H.,The Journal of biophysical and biochemical cytology, 1961.9: p.409-14.


Fig. 1: FIB-SEM cross section of liver cell imaged at 2kV in backscatter electron mode. This image (4096 x 3536 pixels) comes from a series of 430 images with a voxel resolution of 3 x 3 x 10 nm3.

Type of presentation: Oral

IT-13-O-2051 Modeling of Ion Generated Secondary Electrons

Huh U.1, Ramachandra R.2, Joy D. C.3
1University of Tennessee, Knoxville, TN USA, 2University of California,San Diego, CA, , 3Oak Ridge National Laboratory, Oak Ridge, TN, USA
djoy@utk.edu

The arrival of high performance ion beam scanning microscopes has made it essential to have a quantitative model of the ion beam interactions with specimens and their contribution to the generation of the ion induced secondary electron signal (iSE). We have developed an enhanced Monte Carlo simulation, based on our earlier IONiSE program (1) , which is designed to better understand the physics of ion-solid interactions and to perform quantitative simulations. Two key pieces of data are required for this model. The first is the stopping power of the incident ion in the chosen target. Here we use recent data from Berger et al, (2) whose ASTAR program provides stopping power and other data for the He+ ion . ASTAR stopping power profiles were computed for He+ energies from 10keV to 105keV and for elements with atomic number of 90 as seen in figure 1. The second step is to be able to compute the generation rate, the range, and the subsequent transport of the iSE deposited in the sample which has been done by a Monte Carlo method. The Bethe (3) model of secondary electron production requires two parameters, e which represents the generation rate of iSE in the target material, and l which determines the probability of the generated iSE signals reaching the sample surface and ultimately escaping from the specimen surface. So far there is only limited experimental data for iSE yields as a function of their landing energy but good agreement has been found with what little data is available.

References

1, Ramachandra R, Griffin B, Joy DC, (2009), ‘A model of secondary electron imaging in the

Helium Ion scanning microscope’, Ultramicroscopy 109, 748-757

2. Berger M J, Coursey, J S, Zucker M.A,and J. Chang, J, (2011). ‘Stopping-Power and Range

Tables for Electrons, Protons, and Helium Ion’. This is freely available from:

http://www.nist.gov/pml/data/star/index.cfm

3. Bethe H, (1941), The generation of Secondary Electrons, Phys Rev. 59, 940-942


This work was partially supported by the Center for Materials Processing, University of Tennessee

Fig. 1: Composite plot of the variation in stopping power (eV/cm2/1015) predicted by ASTAR for a helium ion source as a function of its beam energy (keV) and of the target material. The black line shows the averaged stopping power for all materials tested as a function of ion energy

Type of presentation: Oral

IT-13-O-2227 Temperature Evolution during FIB Processing of Soft Matter: From Fundamentals towards TEM Lamella Preparation

Schmied R.1, Froech J. E.1, Orthacker A.1, Kraxner J.1, Hobisch J.2, Trimmel G.2, Plank H.1,3
1Center for Electron Microscopy, Graz, Austria, 2Institute for Chemistry and Technology of Materials, University of Technology, Graz, Austria, 3Institute for Electron Microscopy and Fine Structure Research, University of Technology, Graz, Austria
roland.schmied@felmi-zfe.at

During the last decade focused ion beam (FIB) processing became a well-established technique for the site-specific preparation of ultrathin lamellas for transmission electron microscopy (TEM) but also for sub-surface 3D metrology and 3D prototyping on the nanoscale. Beside the undoubted advantages of straightforward implementation, FIB processing entails unwanted side effects, such as ion implantation, amorphization and partial high thermal stress [1]. While the former two are intrinsic properties and therefore invariable, local heating effects have been shown to depend strongly on the patterning strategy. The minimization of this technically induced heating is essential for low melting materials but requires deeper understanding of thermal effects during scanning.
Therefore, accessing local temperatures, its spatial and temporal evolution together with their consequences is essential for FIB processing of sensitive materials with respect to chemical damage and morphological instabilities. In the first part we present an approach, which uses ion trajectory simulations as input data for a thermal spike model which allows the prediction of local temperatures and its lateral distribution during FIB processing (see Figure 1a). Taking into account the thermal behavior of polymers, combined simulations and calculations reveal very good agreement with FIB experiments on polymers (see Figure 1b) confirming the suitability of this combined approach to predict local temperatures and its spatial and timely evolution.
In second step we apply the gained knowledge together with an alternative patterning strategy for the preparation of TEM lamellas, which minimizes technically induced temperature effects [2]. By this careful adaption of the patterning strategy we will show morphological stabilization, characterized via scanning electron microscopy (SEM) and atomic force microscopy (AFM), and demonstrate the reduced chemical damage via IR-Raman spectroscopy. Based on these results, TEM investigations of polymeric layer systems and organic transistors will be shown which confirms stabilized morphologies and minimized chemical damage.
The study demonstrates the massive thermal stress a polymer is exposed during FIB processing and the capabilities of adapted FIB processing for low melting materials which can be easily implemented in most FIB systems. By that, new possibility for FIB processing capabilities for low melting materials open up which have been considered as very complicated or even impossible in the past.
1. J. Mayer, et al., MRS Bulletin 2007, 32, 5, 400 – 407
2. R.Schmied et al. RSC Adv., 2012, 2 (17), 6932 – 6938


The authors gratefully acknowledge the valuable support provided by Prof. Ferdinand Hofer, Prof. Gerald Kothleitner, Dr. Boril Chernev, and Martina Dienstleder. The authors also thank FFG Austria and the Federal Ministry of Economy, Family and Youth of Austria for their financial support.

Fig. 1: (a) simulated 2D temperature distribution in PMMA showing different degrees of material modification (pristine (green), modified (yellow) and volatized (red)); (b) experimental data of minimum line widths for varying pixel dwell times (squares) during standard FIB processing on PMMA compared to the simulated minimum line width(grey band).

Type of presentation: Oral

IT-13-O-2698 A focused Xe+-ion column for fast materials sputtering at high spatial resolution to carry out time-of-flight mass spectrometry with nanoscale precision within a scanning electron microscope

Sedláček L.1, Hrnčíř T.1, Latzel M.2,3, Hoffmann B.2, Jiruše J.1, Christiansen S.2,4
1TESCAN Brno, s.r.o., Brno, Czech Republic, 2TDSU Photonic Nanostructures, Max Planck Institute for the Science of Light, Erlangen, Germany, 3Institute of Optics, Information and Photonics, University of Erlangen-Nuremberg, Erlangen, Germany, 4Helmholtz Center for Materials and Energy, Berlin, Germany
libor.sedlacek@tescan.cz

A unique combination of a high resolution scanning electron microscope (SEM) and a high current focused ion beam (FIB) using a plasma Xe ion source (FERA from TESCAN company) permits extremely high milling and material removal rates [1,2] while simultaneously being able to watch the process so that a precise end-point detection is at hand. Additional analytical add-ons such as energy dispersive x-ray detection (EDX), electron back-scatter diffraction (EBSD) and orthogonal Time-of-Flight Secondary Ion Mass Spectrometry (TOF-SIMS) (TOFWERK company) [3] permit a novel quality of correlated microscopies/spectroscopies.

For TOF-SIMS the high performance focused Xe-ion beam is used to remove the analyte material with the spatial resolution of a FIB. TOF-SIMS provides secondary ion imaging as well as depth profiling, so that a full three-dimensional isotopic images with better than 100 nm lateral resolution are possible.

Compared to a FIB based on Gallium primary ions, the Xenon ion source provides a better detection limit for most of the elements. A quantitative analysis has been demonstrated using a Xe plasma source for material sputtering and alkali elements, such as Li, Na, K constituting the analyte [4]. Detection limits below 2 ppm have been achieved for these species. Moreover, there is no interference when using Xe-FIB instead of Ga-FIB between the analyte and source Gallium ions for material that contain elements such as e.g. Ce, Ge, Ga. Therefore the Xe-FIB is more suitable e.g. for the analysis of important and widely used semiconductor materials and compounds such as SiGe, (In)GaAs and (In, Al)GaN.

Performance of the TOF-SIMS instrument relying on Xe-FIB materials removal has been demonstrated on samples with light-emitting-diode (LED) structures composed of GaN with InGaN Quantum Wells (QWs). A stack of five QWs, each with a thickness of 2.4 nm has successfully been detected (Fig. 1). The focused e-beam of the SEM has been used during TOF-SIMS measurements to account for charge compensation. An analysis of a different LED layer stack showed that a comparably rough interlayer structure is present inside the multi-QWs as demonstrated by the monitoring of TOF-SIMS 27Al+ intensity which is related to a covering AlGaN layer. A 3D reconstruction of an Al rich layer that covers the QWs is shown in Fig. 2.

References:

[1] T Hrnčíř et al, 38th ISTFA Proceedings (2012), p. 26.

[2] J Jiruše et al, Microscopy and Microanalysis 18 (Suppl. 2) (2012), pp. 652-653.

[3] J A Whitby et al, Adv. Mat. Sci. Eng. (2012), 180437.

[4] F A Stevie et al, Surf. Int. Anal. (in press).


The research leading to these results has received funding from the European Union Seventh Framework Program [FP7/2007-2013] under grant agreement No. 280566, project UnivSEM.

Fig. 1: TOF-SIMS depth profile of a GaN/InGaN multi-QW LED sample. Xe primary ion beam current of 550 pA at 30 kV was used for materials sputtering and SEM e-beam current of 1,2 nA at 10 kV was applied to account for charge compensation. Normalized depth profiles of 115In+ and 69Ga+ show well resolved 2.4 nm thick In rich quantum wells.

Fig. 2: 3D reconstruction of TOF-SIMS 27Al+ signal that shows an Al rich layer covering InGaN/GaN QWs in LED structures. View from top (left) and bottom (right) are shown. The z-axis has been expanded 25 times to highlight the interfacial roughness. Field of view is 60 µm x 36 µm.

Type of presentation: Poster

IT-13-P-1477 Development of Cryo-FIB technique for the structural characterization of liquid samples

Tsuchiya M.1, Iwahori T.2, Morikawa A.3, Nagakubo Y.4
1Hitachi High-Technologies Corporation
tsuchiya-miki@naka.hitachi-hitec.com

It is important to observe and characterize the three-dimensional distribution of materials in liquid samples such as; cosmetics, functional paint, catalysts and similar products. Furthermore the requirement for investigating the structure inside of liquids at a microscopic level such as the interface of a dispersoid and dispersant has increased. In order to meet these requirements, we have developed a fully compatible cryo transfer holder for FIB and (S)TEM systems. Another area of development for this holder centers on controlling the temperature of the specimen during either the fabrication or observation and which makes it possible to transport a specimen in the frozen condition. This holder can be cooled to 100K by liquid N2 for observation or fabrication of the sample in a controlled thermal state. To demonstrate the capabilities of this holder a liquid foundation sample was investigated. The liquid foundation was first plunge frozen on a specimen stub and transferred to the holder. The holder containing the frozen foundation sample was then placed into a Hitachi NB5000 FIB-SEM for FIB fabrication.

As a result 100nm thin lamella was produced in the FIB and we were able to observe the structure of the microscopic dispersoid, and its distribution as first viewed in the NB5000 and then a HD2700 shown in Figure1.

Figure 2 shows the EDX maps of the thin foil specimen (Thickness: approximately 100nm) of the frozen foundation. In this result it is possible to clearly see each dispersoid, such as the Fe needle crystals and the granular Ti containing crystals. The results confirm the low temperature and high stability performance of this cryo-transfer holder.

In conjunction to this cryo holder a modified method for the micro-sampling technique which allows a micrometer sized sample to be taken from a millimetre sized specimen is required. For this a new mechanical cryo-probe was developed for sample extraction. This mechanical cryo probe is able to be cooled to 120K.

Figure 3a show the cross-section SEM images of a frozen facial foundation liquid which was lifted out by using a room temperature probe needle. Some hollows appeared inside the micro sample due to the water sublimating from the sample when contacted by the room temperature needle probe.

Figure 3b shows the cross-section SEM image of the same sample as figure 3a but lifted out by the mechanical cryo probe at a temperature of 120K. By using the mechanical cryo probe it is possible to pick up the frozen micro sample while maintaining its shape and we can observe clearly the pigments and dispersant in the frozen liquid foundation.


Fig. 1: BF-STEM images of liquid foundation at an accelerating voltage of 30kV(a) and 200kV(b). (a) Instrument: NB5000 FIB-SEM, Cooling temperature: 100K, Magnification: x3,500. (b)Instrument: HD-2700 STEM, Cooling temperature: 100K, Magnification: x50,000.

Fig. 2: EDX maps of liquid foundation using the 200kV STEM with the cryo transfer specimen holder. dispersoid dispersantInstrument: HD-2700, Acquisition Pixel: 256×200, Acquisition time: 30 min, Cooling temperature: 100K.

Fig. 3: Cross-section SEM images of liquid foundation. With the mechanical probe at room temperature(a), and with the cryo mechanical probe at 118K(b). Instrument: NB5000 FIB-SEM, Acc. Volt.: 1.5 kV, Magnification: x10,000

Type of presentation: Poster

IT-13-P-1484 Finding the needle in the haystack: FIB-SEM combined with array tomography to achieve higher Z-resolution in selected areas

Wacker I. U.1, Bartels C.1, Grabher C.2, Schertel A.3, Schröder R. R.4
1Center for Advanced Materials, Universität Heidelberg, Heidelberg, Germany, 2Institute of Toxicology and Genetics, Karlsruhe Institute of Technology, Karlsruhe, Germany, 3Carl Zeiss Microscopy, Oberkochen, Germany, 4Cryo-EM, CellNetworks, BioQuant, Universitätsklinikum Heidelberg, Heidelberg, Germany
irene.wacker@bioquant.uni-heidelberg.de

Problems in cell or developmental biology often ask for ultrastructural characterisation of a small volume such as a rare event or a specialized substructure inside a large bulk specimen. We propose an intelligent workflow consisting of hierarchical imaging cascades, potentially also relying on different imaging modalities for different resolution ranges. Based on array tomography (AT) [1,2] this allows a stepwise zooming in to a structure of interest from light microscopy via conventional SEM to FIB-SEM.
As a first example we studied a mixed population of cells, a coculture of human tumor cells with immune cells isolated from Zebrafish. Ribbons of serial sections from chemically fixed, epon-embedded cell pellets were placed on silicon wafers and inspected in a reflected light microscope (Fig. 1a). Cell pairs consisting of a large tumor cell and a small fish cell (circle in Fig. 1a) were then imaged in a FEG-SEM (Fig. 1b) revealing immunological synapses between fish immune cells and human target cells. To further characterize their contact region we applied FIB-milling to selected sections to analyze at higher z-resolution only those regions of interest that enclosed centrosomes, Golgi complex, and other membrane-bound organelles (Fig. 1c).
Next we used our approach to identify a rare structure – the neuromuscular junction (NMJ) – within a large tissue block. Tibialis muscle from mouse was chemically fixed, embedded, and serially sectioned. In a single cross section containing hundreds of muscle cells usually only a few cells exhibit part of an NMJ (circle in Fig. 2a). Once an NMJ was found it was imaged in xy on the surface of the section, which in this case was nominally 1µm thick (Fig. 2b). Then FIB-stacks were produced from a 10µm x 10µm area with 10nm step size. Figure 2c shows several images of such a stack with one postsynaptic fold on the left and actomyosin filaments on the right. After alignment the 3D volume can be resliced in xy (Fig. 3a) or volume rendered (Fig. 3b).
Currently we are recording more stacks from corresponding regions of interest in consecutive sections. Fusion of individual stacks into a larger 3D volume allows observing the convoluted network of the postsynaptic folds at a resolution that allows unambiguous tracking of the membranes.
A combination of AT with FIB-SEM is a good approach whenever it is not necessary for a given problem to create a quasi-native molecular atlas of a cell or a total wiring diagram as needed in brain connectomics approaches. In many cases the region of interest is small enough to be amenable to analysis by FIB-SEM.

[1] Micheva and Smith (2007), Neuron 55, 25
[2] Wacker and Schröder (2013), J Microscopy 252, 93


We thank the German Federal Ministry for Education and Research, project NanoCombine, grants FKZ: 13N11401 and FKZ: 13N11403 for financial support.

Fig. 1: (a) Immunological synapse between Zebrafish immune cell and human cancer cell preselected in reflected light microscope; (b) imaged in SEM (Zeiss Ultra); (c) volume rendering of Golgi complex and centrosome in Amira, scale bars: 10µm (a), 1µm (b)

Fig. 2: Imaging of NMJ (Zeiss Auriga): (a) overview of a cross section from mouse leg muscle, circle shows muscle cell containing part of an identified NMJ; (b) postsynaptic folds (orange overlay) imaged on surface of 1µm thick section; (c) postsynaptic folds (orange) in FIB-stack; scale bars: 100µm in (a), 1µm in (b), (c)

Fig. 3: 3D reconstruction from FIB-stack: (a) stack resliced in xy; (b) volume rendering in Chimera

Type of presentation: Poster

IT-13-P-1517 Cross sectional sample preparation of nanowires for TEM analysis using FIB

Lenrick F.1, Ek M.1, Jacobsson D.2, Wallenberg L. R.1
1nCHREM / Center for Analysis and Synthesis, Lund University, Box 124, SE-221 00 Lund, Sweden, 2Division of Solid State Physics, Lund University, Box 118, SE-22100 Lund, Sweden
filip.lenrick@polymat.lth.se

Structuring of materials on the nanoscale is a common way to increase performance in a wide variety of devises. Nanostructures are a challenge for transmission electron microscopy (TEM) sample preparation, as they typically extend from the substrate without surrounding material. One solution has been to remove the nanostructures from their substrate and place them on a TEM grid. Although being simple and time efficient preparation method, it is not always adequate. If the nanostructure-substrate interface is of interest, the structures have a thickness of more than a few hundreds nm, or a TEM projection direction along a long axis is required, an alternative preparation method is necessary.
Here we report on methods for FIB sample preparation for TEM analysis of GaAs-GaInP core shell nanowires. By using polymer resin as support and protection we are able to produce cross-sections both perpendicular to and parallel with the substrate surface with minimal damage. Consequently nanowires grown perpendicular to the substrates could be imaged both in plan and side view, including the nanowire-substrate interface in the latter case. The nanowires, which are roughly 1.5 µm high and 350 nm in diameter, were grown on a GaAs substrate using metalorganic vapour phase epitaxy.
For plan view cross-section, the nanowires were first casted in a tablet-shaped mold using Spurr’s epoxy. The tablet was mechanically polished on one side until a section of the GaAs substrate was exposed. A TEM lamella was extracted using standard in-situ lift out method.
For side view cross-section, the nanowires were first covered in polymer resin by spin coating. Lithography resist proved to be a suitable resin as recipes for precise thickness are available from the resist manufacturers. The viscosity proved low enough not the bend the nanowires during spin coating. After spin coating the resin was soft baked on a hot plate, which increased the viscosity without affecting the nanowires. Finally, a TEM lamella was extracted using standard in-situ lift out method.


Fig. 1: a) TEM overview of plan view cross-section lamella. b) TEM image of upper nanowire marked with a black arrow in a). Shell and core have {011} facets. c) STEM HAADF image of lower nanowire cross-sections marked with a black arrow in a). The shell have both {011} and {112} facets. d) STEM XEDS maps from the red rectangle in a) (scale bar is 100 nm).

Fig. 2: a) TEM overview of side view cross-section lamella. b) HRTEM image of interface marked in a). c) diffractogram of image b). d) DFTEM images, colors correspond to diffractions spots marked in c). e) TEM image from tip of nanowire marked in a). f) DFTEM image from a wurtzite diffraction spot. g) DFTEM image from a zincblende diffraction spot.

Type of presentation: Poster

IT-13-P-1667 Method For Improving FIB Prepared TEM Samples By Very Low Energy Ar+ / Xe+ Ion Mill Polishing

Kauffmann Y.1, Cohen-Hyams T.1, Kalina M.1, Kaplan W. D.1
1Department of Materials Science & Engineering, Technion IIT, Haifa, Israel
mtyaron@tx.technion.ac.il

The great progress in development of new transmission electron microscopes (TEM) during the last two decades has reached a point where the main limiting factor for obtaining fully quantitative and reliable information at the atomic scale is not the optics or the stability of the microscopes but rather the quality of the investigated specimen. The quality of a TEM specimen is determined by how thin and transparent it is to electrons, the surface roughness (variation in local thickness), and the amount of amorphization of the free surfaces (top, bottom and edges) of the specimen.

Well established methods for preparing TEM samples, such as mechanical polishing and electro-chemical polishing, are available. These methods provide very good quality samples when large structures or interfaces are present. When nano-scale site-specific investigation is needed, the best method available is the dual focused ion beam (FIB). This method uses a focused Ga+ ion beam to thin the area of interest down to few tens of nanometers. The ion bombardment of the specimen surface can introduce various artifacts, such as surface amorphization, Ga+ ion implantation, cratering and material re-deposition. These artifacts can be partially reduced by lowering the Ga+ ion energy down to 2 KeV.

For fully quantitative high resolution TEM studies, one needs to get the thinnest possible sample and remove completely all the artifacts introduced by the FIB. This can be achieved by further milling of the FIB sample using well controlled low energy (0.2-0.5 KeV) Ar+ / Xe+ ion milling.

Here we present a method to improve various FIB prepared TEM samples using low voltage ion polishing. This method provides a quick and fairly easy way to prepare high quality TEM samples for fully quantitative and reliable studies.


Type of presentation: Poster

IT-13-P-1686 Focused ion beam sample preparation for atom probe tomography

Cherezova V.1, Chelpanov V.1, Kurushin V.1, Filatov A.1
1Systems for Microscopy and Analysis LLC, Moscow, Russia
v.a.kurushin@gmail.com

In recent times nanostructured multicomponent materials for which spatial distribution of chemical elements is crucial have been widely adopted. The most acceptable method for these materials is atom probe tomography, which allows identifying atom nature and its position in the volume of interest providing 3D imaging with sub-nanometer resolution. During material researching by tomographic atom probe the evaporation of atoms from the sample surface takes place in the process of high electric field. For the achieving of true information the main requirements are preferred for sample geometry: needle-shaped, tip radios, uniform circular cross-section, etc. The focused ion beam sample preparation permits to meet the before-mentioned needs in the optimum way. Now we describe existing FIB-based needle preparation technique (“lift-out”) from particle-reinforced materials for atom probe tomography. This procedure includes following stages: protective layer deposition onto the region of interest (ROI); milling of two regular cross section patterns on both sides of the ROI; the tilt of the stage and cut the lamella almost free from the bulk sample. Then the preparation is proceeded with using in-situ micromanipulator: it is inserted to the ROI; the lamella is attached to the probe using ion beam assisted Pt deposition and cut completely; the lamella is then transferred to a pre-prepared micropost. The next stage includes the lamella attaching (in preferred orientation) to the micropost and milling to detach from the micromanipulator. The final step is thinning of the made sample in order to sharpen its tip to aimed parameters and processing of the needle with low-energy ion beam to remove amorphous layer.
The advantage of chosen method of samples preparation for atom probe tomography is an opportunity to both orient them in parallel with initial surface and cut the sample from the need thickness.


Fig. 1: SEM micrographs (Quanta 3D FEG) of lamella, attached to micromanipulator and maneuvered to the micropost

Fig. 2: SEM micrographs of work piece attached to the micropost

Fig. 3: SEM micrographs of the needle during initial thinning

Fig. 4: SEM micrographs of the needle after final thinning

Type of presentation: Poster

IT-13-P-1712 Sample preparation using Xenon IC-Plasma FIB: benefits and problems.

Audoit G1, Estivill R2, Mariolle D1, Blot X1, Grenier A1, Barnes J P1, Cooper D1
1CEA, LETI, MINATEC Campus, 17 rue des Martyrs, 38054 Grenoble Cedex 9, France., 2STMicroelectronics, 850 rue Jean Monnet, 38926 Crolles, France.
guillaume.audoit@cea.fr

For many types of specimens it is necessary to remove a large amount of material in order to provide an electron transparent region or a needle structure for atom probe tomography (APT). This is particularly true in the case of 3D Integrated Circuit development and manufacturing. This requirement has become sufficient to bring to the market a new generation of commercially available focused ion (FIB) tools that are equipped with inductively coupled Xenon plasma ion sources. This technology allows the generation of beam currents that are twenty times higher than those available using conventional FIBs that use a liquid metal (gallium) ion source. The use of xenon ions to prepare samples such as scanning electron microscope (SEM) cross-sections, transmission electron microscope (TEM) lamellae or even APT needles is attractive because of the theoretical reduced damage of xenon when compared to gallium. Figure 1 shows SRIM simulations of Xe, Ga and Ar ions with different energies in silicon at an angle of incidence of 5° normal to the specimen surface [1]. The simulations suggest that the range of the ions is significantly less for Xe ions. Although the simulations do not account for effects such as channeling in a crystalline specimen, the link between specimen damage and the SRIM simulations has been verified [2]. In this presentation we will introduce the Xe plasma milling system and present measurements of the implantation of Xe ions in Silicon compared to Ga and Ar. For example Figure 2 shows amorphisation and ion implantation profiles have been measured using TEM and APT measurements as a function of the accelerating voltage on silicon and compare to TRIM calculations. We will show the results of specimen preparation using Xe ions, for example of materials that are sensitive to gallium like GaAs and InP which tend to form eutectic compounds that precipitate under Gallium implantation and local heating. Scanning spreading resistance microscopy (SSRM) measurements on InP/GaAs samples cross sectioned with Xenon ions have been compared to Ga-FIB prepared samples in order to compare the sample surface in terms of roughness and dead layer for electrical measurements. Xenon Plasma-FIB specimen preparation also has drawbacks due to the large beam size diameter that has been quantified in this study. We believe that using Xe plasma FIB could open up applications for site specific time of flight (ToF)-SIMS, Auger and XPS analysis, for which the use of a Ga-FIB is impracticable for the production of large enough surface areas, i.e. few hundreds of micrometers.


We thank the Recherche Technologique de Base (RTB) program (national network of large facilities for Basic Technological Research) and the nanocharacterization platform (PFNC).

Fig. 1: Figure 1(a): SRIM simulation of Ar, Xe, and Ga range in silicon as a function of the energy for an angle of incidence of 5°. (b) Magnified plot indicated in (a) by the red shaded region.

Fig. 2: Figure 2: SEM view of an APT needle of silicon milled with 30kV xenon ions and a reconstructed volume showing the Xenon penetration.

Type of presentation: Poster

IT-13-P-1838 FOCUSED ION BEAM LITHOGRAPHY OF SUBWAVELENGTH PHOTONIC 3D-CHIRAL STRUCTURES

Artemov V. V.1, Rogov O. Y.1, Gorkunov M. V.1
1Institute of Crystallography RAS
artemov@ns.crys.ras.ru

The focused ion beam (FIB) milling is a powerful tool for fabricating nanoscale photonic structures [1]. As the next step after successful fabrication of lamellar optical gratings with subwavelength periods [2], we employ the FIB technique to produce truly 3D patterned nano-scale structures. The report describes fabrication and analysis of periodic subwavelength arrays of 3D-chiral holes in a freely suspended silver film.
In order to generate digital templates for FIB lithography (‘stream files’) a special numerical routine has been developed. The templates contain the ion beam waypoints’ coordinates and their ‘dwell time’. Accordingly, all the desired characteristics such as the form, dimensions and the etching depth of a single element can be set. Employing this method allowed us to fabricate periodic arrays of 3D-chiral holes in the freely suspended 200 nm thick silver film with a total processed area of 30x30 μm2 and one element size of 300 nm. Fig. 1a features a 3D lithography model of a single chiral element. The X, Y, and Z coordinates correspond to those of waypoints in the template and the dwell time, respectively. Fig. 1b shows a micrograph of the fabricated structure tilted by 52°, and Fig. 1c shows a normal view of the structure. The fabricated structures have proven to exhibit significant optical activity and circular dichroism.

[1] C. Enkrich, F. Perez-Williard, D. Gerthsen, J. Zhou, T. Koschny, C.M. Soukoulis, M. Wegener, S. Linden, Focused-ion-beam nanofabrication of near-infrared magnetic metamaterials, Adv. Mater. 17 (2005) 2547.
[2] M.V. Gorkunov, V.V. Artemov, S.G. Yudin, S.P. Palto, Tarnishing of silver subwavelength slit gratings and its effect on extraordinary optical transmission, Phot. Nanostr. Fund. Appl. (2013) http://dx.doi.org/10.1016/j.photonics.2013.10.001.


This research was financially supported by the RFBR No. 13- 02-12151 ofi_m and the RAS Presidium program 24. We are grateful to A. L. Vasiliev for the access to the FEI Helios microscope.

Fig. 1: 3D model of the chiral structure unit cell as implemented into the FIB milling digital template (a), SEM micrograph of the fabricated structure tiled by 52° (b), normal view of the fabricated structure (c). 

Type of presentation: Poster

IT-13-P-1861 Measurement of TEM lamella thickness and Ga implantation in the FIB

Lang C.1, Hiscock M.1, Dawson M.2, Hartfield C.2, Statham P. J.1
1Oxford Instruments NanoAnalysis, High Wycombe, UK, 2Oxford Instruments NanoAnalysis, Dallas, USA
matthew.hiscock@oxinst.com

Accurate control over sample thickness and quality is paramount in order to take full advantage of the ever increasing resolution in aberration corrected TEMs. For instruments combining a focused ion beam with an electron beam methods based on either back scattered electron contrast [1] or transmissivity of electrons [2] have been demonstrated for measuring the sample thickness. However, these methods only work on homogenous samples without compositional variations. They also don’t provide any information on the degree of ion implantation.

Here we show a method that uses X-rays generated by the electron beam - lamella interaction to accurately and rapidly measure the lamella composition and thickness. In order to measure the thickness and composition of the lamella, we used Oxford Instruments’ AZtec LayerProbe software [3] and X-Max 150 EDS detectors to acquire and process EDS spectra. LayerProbe refines a starting model of the sample structure against the EDS spectra to calculate the film thickness and composition of the layers. The first layer is defined as the material comprising the lamella. The top layer can be defined to contain the element used as the ion source (e.g. Gallium) to obtain a measure of the degree of ion implantation in the specimen.

Fig. 1a shows an electron image of a TEM lamella prepared from a Ni based superalloy . Fig. 1b shows a surface plot of the lamella thickness and Fig. 1c the Ga thickness calculated from a grid of EDS spectra. The thickness of the lamella is clearly decreasing from the area close to the weld towards the free end of the lamella with the lowest thickness of the lamella measured at around 75nm. The Ga thickness profile shows a different trend with an increase Ga thickness close to the left lower corner and also close to the weld. Fig. 2 shows an X-ray map of a TEM lamella prepared from a silicon semiconductor device. The device structures containing Cu and W are clearly visible in the X-ray maps. One of the Cu lines fades and disappears from the right side to the left of the lamella indicating that the line runs at an angle to the direction of the FIB cut. With LayerProbe it is possible to measure the projected Cu thickness and Si thickness from X-ray spectra reconstructed from the X-ray map. By comparing measurements taken from the right side of the lamella with measurements towards the left side we can see how the thickness increase of the lamella affects the ratio of device vs surrounding Si matrix for both the W and Cu rich device areas.

[1] A. R. Hall, Microscopy and Microanalysis 19 (2013), p. 740.
[2] U. Golla-Schindler, Conference Proceedings EMC (2008), p 667.
[3] C. Lang et al., Microscopy and Microanalysis 19 (2013), p. 1872.


Fig. 1: (a) shows an electron image of a TEM lamella of Ni superalloy 600 and the area for which the lamella thickness in (b) and the equivalent Ga thickness (c) have been calculated.

Fig. 2: The local lamella thickness as well as the contribution of different device layers to this thickness was calculated from spectra reconstructed from an X-ray map.

Type of presentation: Poster

IT-13-P-1915 Advances in ex situ lift out

Giannuzzi L. A.1
1EXpressLO LLC
lucille.giannuzzi@expresslo.com

The focused ion beam (FIB) ex situ lift out (EXLO) technique for scanning/transmission electron microscopy (S/TEM) specimen preparation was historically the first lift out technique developed [1]. EXLO is well known for its ease, speed, and reproducibility, and is perfectly suited for manipulation of electron transparent specimens to carrier devices developed for in situ S/TEM testing as shown in figure 1a. Using EXLO for manipulation to a conventional carbon coated grid limits the specimen from being further FIB milled and inhibits certain S/TEM techniques. The development of a patent pending grid design and technique called EXpressLO™ allows EXLO and manipulation without needing a carbon film support [2-4]. The specimen is lifted out and manipulated directly to a slotted S/TEM grid surface such that the specimen may be directly analyzed and/or further FIB milled, broad beam ion milled or plasma cleaned. Using this new grid design, a specimen can also be manipulated easily into a backside orientation which avoids curtaining artifacts after further FIB milling [3]. The Xe+ ion plasma FIB (PFIB) is capable of producing electron transparent specimens for S/TEM [5]. The EXpressLO™ method can also be used for manipulating large PFIB prepared specimens as shown in figure 1b where a 50 micrometer long specimen is manipulated to a grid [6]. The 1 micrometer thick PFIB specimen manipulated to the EXpressLO™ grid can be further milled using conventional Ga+ ion FIB or a PFIB. EXLO is now flexible and continues to be fast and reproducible which saves labor and FIB instrumentation time, ultimately reducing the cost per specimen.

References:

[1] L.A. Giannuzzi, J.L. Drown, S.R. Brown, R.B. Irwin, F.A. Stevie, Mat. Res. Soc. Symp. Proc. Vol. 480, Workshop on Specimen Preparation for TEM of Materials IV, (1997), Materials Research Society, p. 19-27.

[2] L.A. Giannuzzi, Microscopy and Microanalysis 18, Supp 2, (2012) 632-633.

[3] Lucille A. Giannuzzi, Proceedings of ISTFA, ASM International. (2012), 388-390.

[4] L.A. Giannuzzi, Microscopy and Microanalysis 19, Supp 2, (2013) 906-907.

[5] L.A. Giannuzzi and N.S. Smith, in press, Microscopy and Microanalysis 20, Supp 2, (2014)


Qiang Xu from DENSsolutions provided samples and the in situ MEMS carrier device shown in figure 1a. Noel Smith from Oregon Physics provided the PFIB prepared specimen shown in figure 1b.

Fig. 1: (a) electron transparent specimen manipulated to a DENSsolution in situ carrier device via EXLO. (b) PFIB specimen manipulated via EXLO using the EXpressLO™ method and grid.

Type of presentation: Poster

IT-13-P-1929 Influence of FIB milling on the determination of sp2/sp3 ratio of carbon materials

Zhang X.1, Schneider R.1, Müller E.1, Mee M.2, Meier S.2, Gumbsch P.1, 2, Gerthsen D.1
1Karlsruhe Institute of Technology (KIT), Karlsruhe, Germany, 2Fraunhofer Institute for Mechanics of Materials IWM, Freiburg, Germany
xinyi.zhang2@kit.edu

In focused-ion-beam (FIB) preparation of TEM samples, the energetic Ga+ beam may have damaged the original structure of both sides of the cross-section specimen. Damaged cover layers of amorphous structure are expected for FIB lamellae of carbon materials with modified bonding configurations. Quantitative ELNES technique is well-established for bonding-configuration analysis of carbon. However the sensitivity of this technique is limited by the inevitable FIB-induced damage. Here we propose a simple mathematical model to correct this damage influence on the determination of sp2/sp3 ratios of carbon materials. And the model is tested for HOPG and DLC films with different fractions of sp2 bonds.
The bonding configuration throughout the sampled material column can be considered as a linear combination of those of the damaged layers on both sides and the bulk. Assuming that the damaged cover layers are of uniform thickness and ignoring the local difference in the bonding configurations in the damaged layers and the bulk material, a linear relationship can be derived between the Iπ*/Iσ* ratio for HOPG (or sp2 % for DLCs) obtained from the C-K edge spectra and 1/t, where t indicates the relative thickness obtained from corresponding low-loss spectra. Consequently, the intercept is the real Iπ*/Iσ* ratio (or sp2 %) for the bulk.
FIB preparation for HOPG and DLC samples was followed by a standard lift-out technique. 30 keV Ga+-ions were used for thinning and during the final stage a high tension of only 5 kV was applied to minimize the damage. C-K edge EELS spectra were taken at magic angle (MA) conditions. The cleaved HOPG specimen was largely kept perfect in graphite crystallinity and thus provides as a standard for the FIB-prepared HOPG.
The difference between the FIB-prepared HOPG and the standard is reduced from as high as 20 % to 4 % after the correction. Fig. 1 demonstrates the original quantitative EELS results of two DLCs as a function of 1/t. The DLC (a-C:H) with high sp2 % (69 %) shows little discrepancy with the thickness variation (see red symbols in Fig. 1) and is in accordance with the Raman study (70 %). Therefore, it could imply that the damaged a-C layer contains the same fraction (~ 70 %) of sp2-hybridized C-atoms. Seen from the black symbols in Fig. 1, the ta-C film with lower fraction of sp2 bonding shows a larger dispersion of sp2 % from 39 % to 60 % with respect to t ranging from 0.4 to 1.4, however a linear relationship is indeed found and the sp2 % is corrected to 33 ± 1.3 % by the model. Further assuming that the sp2 % of the FIB-damaged layer is ~ 70 % for all carbon specimens, the damaging depth on each side are estimated to be ~ 15 nm for the HOPG lamella and ~ 10 nm for the ta-C one.


XZ acknowledges funding from China Scholarship Council (CSC) (No. 2010606030). PG acknowledges support from Deursche Forschungsgemeinschaft DFG (project grant Gu 367/30).

Fig. 1: MA-EELS quantification of sp2 % for the a-C:H and ta-C DLCs as a function of the reciprocal of the relative thickness (1/t). Dashed lines are linear fitting results for each sample.

Type of presentation: Poster

IT-13-P-2055 Optimized Detection Limits in FIB-SIMS by Using Reactive Gas Flooding and High Performance Mass Spectrometers

Wirtz T.1, Dowsett D.1, Philipp P.1, Eswara Moorthy S.1
1Department “Science and Analysis of Materials” (SAM), Centre de Recherche Public – Gabriel Lippmann, 41 rue du Brill, L-4422 Belvaux, Luxembourg
wirtz@lippmann.lu

FIB-based instruments play a crucial role in materials science and also in life science. While such FIB instrumentation is an ideal tool for high resolution imaging and nanofabrication, its analysis capability is currently limited. By contrast, Secondary Ion Mass Spectrometry (SIMS) is an extremely powerful technique for analyzing surfaces given its excellent sensitivity, high dynamic range, high mass resolution and ability to differentiate between isotopes. Adding SIMS capability to FIB instruments offers not just the prospect of obtaining SIMS information limited only by the size of the probe-sample interaction (~10nm) but also enables a direct correlation of such SIMS images with high resolution secondary electron images of the same zone taken at the same time (correlative microscopy).

Past attempts of performing SIMS on FIB instruments were rather unsuccessful due to unattractive detection limits, which were due to (i) low ionization yields of sputtered particles, (ii) extraction optics with limited extraction and collection efficiency of secondary ions and (iii) mass spectrometers having low duty cycles and/or low transmission. In order to overcome these limitations, we have investigated the use of reactive gas flooding during FIB-SIMS and we have developed compact high-performance magnetic sector mass spectrometers with dedicated high-efficiency extraction optics.

Our results show that the yields obtained with Ga+, He+ and Ne+ bombardment, which are intrinsically low compared to the ones found in conventional SIMS, may be drastically increased (up to 4 orders of magnitude) by using reactive gas flooding during analysis, namely O2 flooding for positive secondary ions and Cs flooding for negative secondary ions (Figure 1) [1-3]. The resulting detection limits vary from 10-3 to 10-6 for a lateral resolution between 10 nm and 100 nm (Figure 2).

The emitted secondary ions are extracted by dedicated optics which we have designed for several FIB-based instruments and injected into a specially designed compact high-performance magnetic sector double focusing mass spectrometer. The obtained extraction efficiency ranges from 40% to 100% while successfully avoiding any artefacts (broadening or distortion) regarding the primary ion beam. The specifications of the mass spectrometer include highest transmission (100%), high mass resolution (M/DM > 2000), full mass range (H-U) and parallel detection of several masses.

Here we will present the FIB-SIMS systems we have developed, give an overview of the obtained performances and present typical examples of applications.

References

[1] P. Philipp et al., Int. J. Mass. Spectrom. 253 (2006) 71

[2] T. Wirtz et al., Appl. Phys. Lett. 101 (2012) 041601

[3] L. Pillatsch et al., Appl. Surf. Sci. 282 (2013) 908


Fig. 1: Enhancement of secondary ion yields using reactive gas flooding under He+ and Ne+ bombardment.

Fig. 2: Detection limit using a Ga+ FIB with and without Cs0 flooding vs. minimum feature size: example for the detection of Si-.

Fig. 3: Ga+ FIB-SIMS image of the Ca distribution in skin cells (field of view: 50x50 µm2)

Type of presentation: Poster

IT-13-P-2068 Characterization of carbonaceous contamination and the cleaning capability of atomic hydrogen during focused ion beam processing

Steiger-Thirsfeld A.1, Basnar B.2, Tomastik C.3, Pongratz P.4, Lugstein A.2
1University Service Center for Transmission Electron Microscopy, Vienna University of Technology, Vienna, Austria, 2Institute of Solid State Electronics, Vienna University of Technology, Vienna, Austria, 3AC2T Research GmbH Austria, Wr. Neustadt, Austria, 4Institute of Solid State Physics, Vienna University of Technology, Vienna, Austria
steiger@ustem.tuwien.ac.at

During focused ion beam (FIB) milling undesired peculiar depositions (Fig. 1a), from mainly carbon compounds of the residual gas at the edge of ion beam exposed regions, occur. Surface diffusion of residual gas molecules and the reduced ion dose are causing enhanced residual gas deposition rates at the boundary area of ion beam exposed regions [1]. Especially for sputtered structures with dimensions in the nanometer range, these unwanted ion beam induced depositions (IBID) can be of the same order of magnitude as the intended structures, therefore they must be eliminated.

Preliminary, we have characterized the dynamics of this contamination growth. Crystalline (c-Si) and amorphous Si (a-Si) were irradiated by a 50 kV Ga+ ion beam, with various scanning parameters, under high vacuum conditions (2.7 × 10-7 mbar). A refresh time variation in several steps from 23 ms to 15 s was performed with a dwell time of 1 μs (Fig. 1b). Atomic force microscopy (AFM) measurements reveal that for refresh times longer than 1 s, IBID entirely inhibits the net volume loss by sputtering. The residual gas deposition rate saturates for a refresh time of about 5 s (Fig. 2a). Dwell time variations from 0.5 μs to 10 μs at a refresh time of 5 s show the expected decreasing in deposition rate. The pronounced indented shape of the deposits exhibits the influence of surface diffusion. A surface diffusion constant in the order of 10-9 cm2/s was roughly estimated. Auger electron spectroscopy (AES) depth profiles of the chemical composition of low dose ion irradiated c-Si show a mixture of Si, C, and O with decreasing C and O concentrations from top to the bottom of the deposits (Fig. 2b). The investigated depositions consist roughly of of 22 at% C and 2 at% O in addition to Si.

When an atomic hydrogen (H*) gas beam, generated by a thermal cracker source [2], has been delivered to a c-Si surface during FIB processing, reduced swelling heights in the range estimated from the difference in mass density of amorphous and crystalline Si were determined (Fig. 3a). Moreover, the H* delivery suppresses the residual gas deposition, even for refresh times in the range of seconds (Fig. 3b). We conclude that simultaneous Ga+ and H* bombardment removes adsorbed carbon compounds and the additional H* delivery impedes the surface diffusion of adsorbed residual gas molecules. Thus a significant build up of contamination can be avoided.

[1] J.B. Wang, A. Datta, Y.L. Wang; Applied Surface Science 135 (1998) 129-136
[2] K. G. Tschersich, J.P. Fleischhauer, H. Schuler; Journal of Applied Physics 104, 034908 (2008)


This work has been supported by the EC (FP6, CHARPAN, Contract no.: IP 15803-2).

Fig. 1: AFM images of ion beam induced residual gas depositions. 1 (a) Undesired IBID at the edge of a dot pattern. 1 (b) Indented IBIDs on c-Si for fluence values in the range of 6.9 × 1013 ions/cm2 to 2.1 × 1016 ions/cm2 with a refresh time of 5 s.

Fig. 2: Characterization of ion beam induced residual gas depositions. 2 (a) Deposition height dependence on ion fluence with the refresh time as parameter measured by AFM. 2 (b) AES depth profile of the chemical composition of low fluence (6.9 × 1014 ions/cm2) ion irradiated c-Si.

Fig. 3: Cleaning effect of atomic hydrogen during FIB processing. 3 (a) Comparison of height (depth) characteristics of FIB processed areas in c-Si and a-Si with and without additional H* delivery. 3 (b) Comparison of height (depth) characteristics of FIB processed areas at a refresh time of 5 s with and without additional gas delivery.

Type of presentation: Poster

IT-13-P-2075 Nanopatterning Plasmonic Structures Using Focused Ion Beam and E-Beam Lithography

Cohen Hyams T.1, Spektor G.2, Gal L.2, Orenstein M.2
1Department of Materials Science & Engineering, Technion, Haifa, Israel, 2Department of Electrical Engineering, Technion, Haifa, Israel
tzipic@technion.ac.il

The Focused Ion Beam (FIB) system uses a Ga+ ion beam to raster over the surface of a sample in a similar way as the electron beam in a scanning electron microscope (SEM). One of the capabilities of FIB is its ability to mill complex nanopatterns (including bitmapped images), making it the ideal tool for precise 3D maskless nanopatterning of a wide variety of materials. The FIB is the ideal tool for prototyping a wide range of devices in the R&D stage of product development, since it offers high reproducibility and scalable throughput. However, the ion bombardment of the specimen surface can introduce various artifacts, such as surface amorphization, Ga+ ion implantation, cratering and material re-deposition.

Electron-beam lithography is an alternative tool for 3D nanopatterning. E-Beam lithography uses a focused beam of electrons to “write” patterns on a surface covered with an electron sensitive film (resist). The main advantage of electron-beam lithography is that it can direct-write with nm resolution. This form of maskless lithography has high resolution and low throughput.

These two techniques are considered to be the best methods for fabricating structures for surface plasmon coupling and manipulation. The structures can be used to obtain confined longitudinally polarized plasmonic focal spots and other higher order effects.
In this study, we present a comparison between FIB nanopatterning and E-beam lithography to fabricate various nano engraved plasmonic structures in gold comprising different spiral types and their engagements.


Type of presentation: Poster

IT-13-P-2432 Optimization of the sample preparation method for semiconductor dopant contrast observation with SEM

Druckmüllerová Z.1,2, Kolíbal M.1,2, Vystavěl T.3, Šikola T.1,2
1Institute of Physical Engineering, Brno University of Technology, Brno, Czech Republic, 2CEITEC, Brno University of Technology, Brno, Czech Republic, 3FEI Company, Brno, Czech Republic
zdenadr@seznam.cz

Since semiconductor devices are being scaled down to dimensions of several nanometers, there is a growing need for techniques capable of quantitative analysis of dopant concentrations at nanometer scale in all three dimensions. Therefore we optimized the sample preparation methodology for imaging dopant contrast by scanning electron microscopy (SEM) at incident electron energies about 1keV [1], which enables to visualize and analyze dopant concentration changes. SEM analysis at such conditions became widely used providing promising results, but many unresolved issues hinder its routine application for device analysis, especially in case of buried layers where the site-specific sample preparation is challenging. We report on optimization of a site-specific sample preparation by the focused Ga ion beam (FIB) providing an improved dopant contrast in SEM. As a testing sample we used differently doped multilayer structure deposited on Si (see Fig. 1). Similarly to the lamella preparation for transmission electron microscopy by FIB, a polishing sequence with decreasing ion energy is necessary to minimize the thickness of the electronically dead layer [2]. We have achieved the contrast values comparable to the cleaved sample, being able to detect dopant concentrations down to 1x1016 cm-3.(see Fig. 2). A theoretical model shows that the electronically dead layer corresponds to an amorphized Si layer [3] formed during ion beam polishing. Our results also demonstrate that the contamination caused by electron beam scanning is significantly suppressed for focused-ion-beam treated samples compared to the cleaved ones.

References:
[1] Chakk Y. & Horvitz D. (2006). Contribution of dynamic charging effects into dopant contrast mechanisms in silicon. J. Mat. Sci. 41, 4554-4560.
[2] Giannuzzi L. A., Geurts R. nad Ringnalda J. (2005). 2kV Ga+ FIB milling for reducing amorphous damage in silicon. Microsc. Microanal. 11 suppl.2, 828-829.
[3] Kazemian P., Twitchett A. C., Humphreys J. C. & Rodenburg. C. (2006b). Site-specific dopant profiling in a scanning electron microscope using focused ion beam prepared specimens. Appl. Phys. Lett. 88, 212110.


Sample was provided by Cornelia Rodenburg, PhD., University of Sheffield, Great Britain. This work was supported by the Grant Agency of the Czech Republic (P108/12/P699) and by European Regional Development Fund – (CEITEC - CZ.1.05/1.1.00/02.0068). M. K. and Z.D. acknowledge the support of FEI Company.

Fig. 1: Description of the sample used: SEM image of the cleaved multilayer structure deposited on n-doped silicon substrate is in the middle, a schematic description of alternating layers of intrinsic silicon (150 nm) with p-doped layers in depicted on the left. The contrast profile normalized to n-doped substrate is shown on the right.

Fig. 2: SEM contrast improvement of sample cut by Ga FIB using decreasing final polishing energy.

Type of presentation: Poster

IT-13-P-2440 Channeling contrast: a cost effective alternative to EBSD orientation mapping in scanning focused probe instruments (SEM/FIB) ?

Langlois C. T.1, Yuan H.1, Douillard T.1, Van de Moortele B.2, Descamps-Mandine A.3, Blanchard N.4, Epicier T.1
1MATEIS Laboratory, INSA Lyon, France, 2LGL laboratory, Ecole Normale Supérieure Lyon, France, 3INL Laboratory, INSA Lyon, France, 4Light Matter Institute, Claude Bernard University of Lyon, France
cyril.langlois@insa-lyon.fr

Generally, for grains at the micron scale, orientation maps are obtained by Electron Back Scattered Diffraction (EBSD) in a Scanning Electron Microscope (SEM). For various reasons, orientation maps with the EBSD technique can be challenging: material properties (conductive or not, preparation problems), geometrical setup (impossibility to collect the signal), pseudo-symmetry Kikuchi diffraction patterns, nano-sized structures, large area mapping, etc. In this context, it is worth investigating other means to map the crystallographic orientations of grain.
Channeling contrast is a well-known phenomenon allowing grains of a polycrystalline sample to be distinguished, even if only one phase is present. Depending on the orientation of the crystallographic planes, the secondary and backscattered electron yields vary from one grain to the other, resulting in different intensities received by the detectors. For this reason, the contrast of each grain varies when the orientation of the sample is changed [3].
We show in this study how it is possible to use this channeling effect (with an electron or an ion beam) to obtain the three Euler angles characterizing the orientation at a given point of the sample surface. The main concept is to obtain an intensity profile at that point, to compare the intensity profile to a semi-empirical database of profiles and to find the best fit, i.e. the three Euler angles associated with this point. Comparing with EBSD measurements on the same area allowed us to determine the precision of the indexation, which is better than 5°. We discuss then the best way to obtain an intensity series, either by tilting or rotating the sample, with regards to acquisition stability and unicity of the indexation. The issue of acquisition time is also discussed, and an example of our indexation method based on a 30 sec movie over 360° is shown. We conclude by evaluating the pros and cons of using ions or electrons for such indexation purposes.
[1] Estimation of recrystallized volume fraction from EBSD data, J. Tarasiuk, Ph. Gerber and B. Bacroix, Acta Materialia (2002) 50 1467–1477
[2] Characterization of the Grain-Boundary Character and Energy Distributions of Yttria Using Automated Serial Sectioning and EBSD in the FIB, S.J. Dillon and G.S. Rohrer, J. Am. Ceram. Soc. (2009) 92 1580–1585
[3] Crystallographic orientation contrast associated with Ga+ ion channeling for Fe and Cu in focused ion beam method, Y. Yahiro, K. Kaneko, T. Fujita, W.-J. Moon and Z. Horita, J. Electron Microscopy (2004) 53 571–576


Fig. 1: Snapshot of the software written to extract intensity profiles from a tilt series and Euler angles from an EBSD map acquired on the same area

Fig. 2: Snapshot of the software written to compare the experimental profile (green), the semi-empirical profile corresponding to EBSD orientation (red), and the best fit semi-empirical profile found in our database (blue). In this case, the disorientation between the EBSD orientation and the ‘best fit’ orientation is around 3°only

Type of presentation: Poster

IT-13-P-2481 Cryogenic FIB Lift-out as a preparation method for damage-free soft matter TEM imaging 

Parmenter C. D.1, Fay M. W.1, Hartfield C.2, Amdor G.2, Moldovan G.2
1University of Nottingham, 2Oxford Instruments Nanoanalysis
christopher.parmenter@nottingham.ac.uk

We have demonstrated that it is possible to prepare and remove a thinned lamella and transfer to the TEM, whilst maintaining cryogenic conditions. Once further refined, this method offers the possibility of compression and stain artefact free imaging of soft matter samples (cells, tissues, plant samples, polymers, gels etc) preserved and maintained at cryogenic temperatures. Biological samples contain a high degree of water, which dehydrate under vacuum. Solutions are: critical point drying, resin impregnation with heavy metal stains or cryogenic fixation. Once stabilised the samples can be prepared with an ultramicrotome to yield electron transparent sections, however, they commonly suffer from compression and/ or knife artefacts. In addition, there is a desire to move away from staining or methods which can induce structural re-arrangement. The removal of a thinned lamella from a bulk sample for Transmission Electron Microscopy (TEM) analysis has been possible in the Focused Ion Beam Scanning Electron Microscope (FIB-SEM) for over 20 years either via in-situ (by use of a micromanipulator) or ex situ lift-out approaches [1]. Both are currently only applied to samples at room temperature as there are a number of technological and sample handling issues for cryogenic samples. Recent efforts have demonstrated cryo lift-out is possible for materials samples[2]. This work further extends the development of cryo lift-out to allow label and damage-free imaging of soft and biological structures. To preserve the vitreous nature of the water in cryo-preserved samples the temperature should be maintained below -140°C and the probe tip held by the manipulator cooled to at least -130°C. To achieve this, an OmniProbe 100 was modified with a thermal break and cooling braid, which was attached to the cold finger of the cryo stage (Quorum PPT 2000).

Prior to lamella extraction, an alginate-collagen hydrogel, was sputter-coated with platinum and a tungsten layer from a gas injector. The gel was milled using a modified TEM lamella protocol to approximately 2μm thickness, before the lamella was attached to the cooled tip by cryo-condensation of water via a gas injector (figure 1). The lamella was subsequently secured to a TEM (lift-out) support grid (figure 2) and further thinned to electron transparency (figure 3). The sample was transferred under liquid nitrogen to a cryo-TEM holder and imaged at 200 kV in both bright and dark field imaging (figure 4). 

[1] L Giannuzzi et al. in “Introduction to Focused Ion Beams: Instrumentation, Theory, Techniques and Practice”, ed. LA Giannuzzi and FA Stevie, (Springer, 2005) Chapter 10, p.201-228.

[2] N Antoniou et al, Conf. Proc. 38th Int. Symp. Testing and Failure Analysis (2012) p. 399-405.


Many thanks to Dr James Dixon (University of Nottingham, CBS) for supplying the hydrogel samples.

Fig. 1: Cryo-FIB milling of a bulk sample to prepare a thin lamella, scale bar 5 µm

Fig. 2: Extraction of the lamella by the cooled manipulator after attachment and release of lamella, scale bar 10 µm.

Fig. 3: Micrograph of the attached and thinned lamella, scale bar 5 µm

Fig. 4: Dark field TEM image of collagen fibrils in an alginate hydrogel matrix, scale bar 100 nm

Type of presentation: Poster

IT-13-P-2485 Characterization of Ga+ FIB Damage in Electron Beam Induced Deposited Platinum, Tungsten and Carbon Chemistries for In-situ S/TEM Sample Preparation

Van Leer B.1, Landin T.1, Wall D.1, Roussel L.1
1FEI
brandon.van.leer@fei.com

TEM/STEM sample preparation by focused ion beam (FIB) and SEM/FIB instrumentation has become routine in the last several years. Technology advances in automation and in-situ techniques have reduced preparation times for sub-50 nm lamellae to less than an hour and with state-of-the-art technology less than 30 minutes [1]. DualBeams (FIB-SEM) are also often used for micro- and nanoprototyping applications.  For S/TEM sample preparation electron or ion beam induced deposition (EBID, IBID) is required to planarize the region of interest to minimize artifacts generated by the FIB [2].

Many studies have investigated surface and sidewall lamella damage in Silicon by FIB [3]. In addition to sidewall damage by FIB for cross-sections or FIB processed S/TEM samples, surface damage must also be considered for FIB preparation especially when characterization of the sample surface is required. It has been shown that low energy electron beam induced deposition (EBID) imparts the smallest surface damage when compared to Ga+ ion beam induced deposition (IBID) [4]. However, the rate of deposition with EBID is ~ 20X slower than IBID, thus understanding the damage depth into EBID layer during the FIB deposition process will reduce process time for cross-section or S/TEM sample preparation.

Approximately 100 nm EBID C, W and Pt layers were deposited onto Si using 5 keV; 6.3 nA (C and W) and 2 keV; 6.3 nA (Pt). 30 keV Ga+ IBID C, W and Pt layers were deposited over the layers of interest and FIB prepared for STEM in SEM analysis. A 5 keV Ga+ FIB Pt IBID over 2 keV Pt EBID sample was also prepared. Each face of the lamellae was FIB milled using Ga+ ions at 30 keV and 88.5 degrees incident angle, followed by 5 keV at 85 degrees incident angle. Figures 1a and 1b are STEM in SEM images of the Pt and C experiments respectively. Measurements (Figure 2) reveal that IBID tungsten penetrated approximately 15 nm into the EBID tungsten while the IBID platinum penetrated more than 3X deeper into EBID platinum. This value decreased to approximately 16 nm when 5 keV IBID was employed for Pt.


[1] D. Wall, “Ultra-Fast In-Situ Sample Preparation.” FEI P/N 04AP-FR0111, FEI Company, 2007.
[2] Introduction to Focused Ion Beams, eds. L.A. Giannuzzi and F.A. Stevie, Springer (2005).
[3] L. A. Giannuzzi et al., Micros. Microanal., 11(Suppl 2) (2005), p. 828.
[4] B.W. Kempshall et al., J. Vac. Sci. Tech. B, 20(1) (2002) 286.

Fig. 1: 30 keV STEM in SEM images of surface 30 keV Ga+ FIB damage during Pt IBID into EBID Pt

Fig. 2: 30 keV STEM in SEM images of surface 30 keV Ga+ FIB damage during C IBID into EBID C

Fig. 3: Average FIB damage depth (nm) during IBID of Pt, W and C over EBID Pt, W and C

Type of presentation: Poster

IT-13-P-2498 FIB/SEM tomography for 3D visualization of virus infected fibroblasts with TEM-like resolution.

Villinger C.1, 2, Neusser G.3, Kranz C.3, Walther P.2, Mertens T.1
1Institute of Virology, University Medical Center Ulm, Germany, 2Central Facility for Electron Microscopy, Ulm University, Germany, 3Institute of Analytical and Bioanalytical Chemistry, Ulm University, Germany
clarissa.villinger@uni-ulm.de

Keywords: FIB/SEM tomography, high pressure freezing, freeze substitution

Focussed ion beam/scanning electron microscopy (FIB/SEM) tomography is a novel electron microscopy technique that is increasingly used within life sciences. Here we present its application for ultrastructural visualization of fibroblasts infected with human cytomegalovirus (HCMV). For that we employed optimized sample preparation protocols including high pressure freezing and freeze substitution. The result was an improved ultrastructural preservation and a high image contrast. Additionally, our well established embedding procedure of cell monolayers in Epon allows not only ultrathin sectioning but also sample preparation for FIB/SEM tomography. The detection of the secondary electron signal allows us to resolve cellular and viral ultrastructures down to the level of lipid bilayers (Fig. 1).

The main focus of our work is the egress of HCMV capsids from the nucleus (capsid formation and genome packaging) into the cytoplasm (further virion maturation). The diameter of virus capsids exceeds the diameter of nuclear pores. Hence, the virus has to find an alternative path to leave the nucleus. It is known that this process is made possible by an envelopment and de-envelopment process at the inner (INM) and outer nuclear membrane (ONM), respectively [1, 2]. We visualized a portion (z=5 µm) of an HCMV infected nucleus with FIB/SEM tomography. The resulting 3D reconstruction showed that the surface area of the INM was increased through large infoldings which can extend deep into the nucleoplasm (Fig. 2). DNA free and DNA filled capsids were both present within these infoldings (perinuclear capsids). This model has already been postulated based on 2D TEM images [3]. Our 3D data now confirm this model. Additionally, the slice thickness of 20 nm allowed imaging of every nuclear as well as perinuclear capsid within the analyzed volume. This gives the opportunity to analyze the capsid distribution in 3D, thus, making the results and interpretation of 2D TEM images more accurate. In conclusion, our FIB/SEM data provide a detailed image of the nuclear stages of HCMV morphogenesis, from capsid assembly and DNA packaging to capsid egress.

[1] Skepper JN et al. (2001).J. Virol. 75, 5697–5702.

[2] Mettenleiter TC et al. (2013). Cell. Microbiol. 15, 170–178.

[3] Buser C et al. (2007). J. Virol. 81, 3042–3048.


Fig. 1: Mature HCMV particle within a vesicle. The resolutions of the SEM and the TEM images are comparable. The two leaflets of the lipid bilayers are resolved in both images. The high resolution in the SEM image is gained by a primary acceleration voltage of 5kV and detection of the secondary electron signal. Diameter of virus particle approx. 200 nm.

Fig. 2: 3D reconstruction of HCMV infected nucleus. DNA free (orange) and DNA filled (green) capsids are enclosed by large infoldings of the INM (blue). Only the outline of the infolding is depicted in (A), (B+C) show complex internal membranes. (B) The infoldings can be spherical and/or tubular.

Type of presentation: Poster

IT-13-P-2590 Enabling future Nanotomography and Nanofabrication with Crossbeam technology

Schulmeyer I.1, Kienle M.1
1Carl Zeiss Microscopy GmbH, Carl-Zeiss-Str. 22, 73447 Oberkochen
ingo.schulmeyer@zeiss.com

A Focused Ion Beam (FIB) combined in one instrument with a Field Emission Scanning Electron Microscope (FE-SEM) has become a powerful instrument for numerous Standard and cutting edge applications in Research and Industry. The FIB is mainly used to open up the third dimension to a SEM. The ion beam is not only able to cut slices of a samples surface for tomographic imaging, but can also be used to create new materials or functional structures with superior properties. With increasing application maturity also the demand for faster systems, complete detection, more precise structuring and a wider application range rises.

3D-imaging and –analytics allow a complete characterization of a samples volume. It is widely used in Materials and Life Sciences and allows better understanding of compound materials, brain tissues, electronic devices and other samples. To achieve representative information of a sample, the analyzed volume needs to be sufficiently large and the resolution has to be high. The voxel resolution in tomography is mainly limited by the thickness of the slices cut with the ion beam [1]. The Crossbeam 540 allows the thinnest slices down to 5 nm and below. To keep the thickness homogenous over thousands of slices, a long-term stable FIB comes along with a sophisticated Software solution including adaptive slice thickness tracking (Figure 1). In interaction with a charged particle beam, each material behaves different. We introduce solutions to process, image and analyze all kinds of samples, including charging, outgassing or dirty and even magnetic samples. We will discuss our latest developments in detector technology to improve the signals especially at low acceleration voltages and the acquisition time.

Preparation of samples for TEM or other imaging techniques is one of the main applications for modern FIB-SEM instruments. Not only the quality, thickness and homogeneity of the prepared specimen matters, but also an fast, easy and integrated workflow. The quality of the results in TEM and STEM depend not only on the instrument, but to a high degree on the analyzed specimen. The unique X² method combined with a FIB that performs excellently at low voltages provides homogenous TEM lamellas with < 10nm thickness and minimum amorphous layer [2] (Figure 2).

Modern FIB-SEM systems cover not only the typical application range but are also used to host and integrate numerous instruments and components for advanced experiments. The Zeiss FIB-SEM allows integration of many components for In-Situ analysis like heating, cooling and tensile stages, SIMS, EBIC and CL, different manipulators, e.g. for probing and liftout, a Laser option for micropatterning or fast removal of large amounts of material and many more.


[1] L. Holzer and M. Cantoni, Review of FIB-tomography, Nanofabrication Using Focused Ion and Electron Beams: Principles and Applications (2011), p. 410ff

[2] L. Lechner, J. Biskupek and U. Kaiser, Improved Focused Ion Beam Target Preparation of (S)TEM Specimen - A Method for Obtaining Ultrathin Lamellae, Microscopy and Microanalysis 18 (2012), p. 379-384

Fig. 1: Innovative thermo-electrical generator transforming exhaust heat into electrical power (Mg2Sn-Mg2Si). 3D stack was acquired on Crossbeam running ATLAS3D. 8nm Voxel Size. The diffusion zone thickness varies between 10 to 25µm

Fig. 2: large 50µm x 20µm TEM lamella prepared automatically in 25 min (left). Lamella thinned using the X²-method (centre). HRTEM image of a lamella demonstrating atomic resolution at 20kV. The spotty contrast variations in the image are caused by strong dynamic diffraction contrast that cannot be avoided at 20 kV [2] (right).

Type of presentation: Poster

IT-13-P-2616 In situ FIB/SEM micro-compression tests of layered crystals

Schweizer P.1, Niekiel F.1, Butz B.1, Spiecker E.1
1Center for Nanoanalysis and Electron Microscopy (CENEM), Department of Materials Science and Engineering, University of Erlangen-Nürnberg, Cauerstr. 6, 91058 Erlangen, Germany
peter.schweizer@ww.stud.uni-erlangen.de

Keywords: micro-compression, layered crystals, FIB, transition metal dichalcogenides

With the recent surge of interest in two-dimensional materials other than graphene, transition metal dichalcogenides (TMDCs) have gained a lot of attention because of their outstanding properties ranging from superconductivity to the formation of charge density waves [1, 2]. However not only monolayers of TMDCs are of interest, since the bulk form also shows strongly anisotropic behavior in most physical properties, which makes TMDCs promising candidates for many applications such as solid lubrication [3]. The anisotropic behavior in these layered crystals arises from the contrast between the weak Van der Waals bonding between the layers and the strong ionic/covalent bonds within the layers. Despite the amount of attention that this class of material has gained, the mechanical properties are largely unknown. This is due to the difficulty of preparing samples that are suitable for traditional mechanical testing. Consequently compression tests of micro-pillars are emerging as a novel way to measure the mechanical properties of materials on a micro scale [4].
In this contribution we show the preparation of micro-pillar samples from layered crystals, choosing vanadium diselenide (VSe2) as model material system. An FEI Helios Nanolab 660 DualBeam has been equipped for both sample preparation and in situ compression of the studied pillars. The pillars have been prepared using the focused ion beam (FIB), compression has been performed using a Kleindiek micromanipulator while imaging in situ using SEM. Force measurement is enabled using a Kleindiek SpringTable. Image correlation is used to determine the deflection of a cantilever, which corresponds to a force via a known spring constant. Deformation is measured tracking the difference in displacements between substrate and pillar. The complete indentation setup is shown in Figure 1.
Figure 2 shows the compression of an exemplary VSe2 pillar, cut at an angle of 20 degrees from the basal planes. The inset shows the resulting force displacement diagram. As expected a preferential slip along the basal planes is clearly visible. The sample preparation and micro-compression testing route established in this work on the example of VSe2 is anticipated to provide a deeper insight on the mechanical properties of TMDCs and other more complex layered crystals , like, e.g., misfit layer compounds.

References
[1] Q H Wang et. al., Nat. Nanotechnol., 2012, 7, 699–712
[2] X Huang et. al., Chem. Soc. Rev., 2013, 42, 1934
[3] L Rapoport et. al., J. Mater. Chem., 2005, 15, 1782–1788
[4] M D Uchic et. al., Science, 2004, 305, 986


The Authors gratefully acknowledge financial support by the German Research Foundation (DFG) via the research training group 1896 „In situ microscopy with electrons, X-rays and scanning probes”.

Fig. 1: The Kleindiek indenter setup inside the chamber of the focused ion beam.

Fig. 2: SEM-image of a vanadium diselenide pillar (prepared at an angle of 20° relative to the basal planes) being compressed by a diamond flat punch and the resulting force displacement curve (inset).

Type of presentation: Poster

IT-13-P-2701 High performance nanomachining using the new analytical FIB-SEM system

Jiruše J.1, Havelka M.1, Haničinec M.1, Polster J.1, Hrnčíř T.1
1TESCAN Brno, s.r.o., Brno, Czech Republic
jaroslav.jiruse@tescan.cz

A new analytical tool GAIA, combining high performance Focused Ion Beam (FIB) column with ultra-high resolution Scanning Electron Microscope (SEM), has been developed. The SEM resolution has been improved down to 1 nm at 15 kV and 1.4 nm at 1 kV, see Figure 1, thanks to a new objective lens of a single-pole immersion type [1, 2]. It creates a strong magnetic field surrounding the sample and decreasing optical aberrations. Intermediate lens enables to work in the magnetic-field-free mode suitable for analysis, magnetic sample imaging and observation during FIB machining. FIB milling process can be controlled by SEM imaging simultaneously, because of two independent scanning generators and sophisticated TESCAN detection system. Besides chamber detectors for detection of secondary (SE), backscattered (BE), transmitted (TE) electrons and secondary ions (SI), InBeam SE and InBeam BE detectors placed in the column give the free space around the sample.

The new non-magnetic Cobra FIB column with high resolution of 2.5 nm at 30 kV [3] and great performance at high currents has been designed to protect from the influence on the magnetic field of the immersion SEM objective lens. DrawBeam software allows drawing of patterns in CAD-like GUI for electron and ion beam lithography. The patterns are generated by ultra-fast scanning generator with pixel dwell time down to 20 ns. The novel milling strategy is included in DrawBeam software for 2.5 times faster FIB cross sectioning, see Figure 2. The technique is based on the correction of the intended shape to maximize the milling rate and to minimize the redeposition effects. The new AutoSlicer software for the automated cross sectioning and TEM lamella preparation increases FIB performance even further, see Figure 3. FIB column control is greatly simplified by using TESCAN In-Flight Beam TracingTM technology, which newly enables to compute and optimize FIB column settings.

GAIA instrument is prepared for fabrication and observation of non-conductive samples. Charge accumulation on the surface caused by FIB milling can be neutralized by integrated electron flood gun or SEM electron beam. SEM imaging without charging artifacts can be performed at critical energy, below 4 keV, with improved resolution. TESCAN beam deceleration technology, applying negative voltage on the sample, allows automatic control of the electron landing energy down to 50 eV (manually down to 0 eV) and it further improves SEM resolution at low beam energies, see Figure 1. The new control module provides sample discharge and touch-alarm protections.

References:

[1] Z Shao et al, Rev. Sci. Instrum. 60 (1989) p. 3434.

[2] J Jiruše et al, Microsc. Microanal. 19 (Suppl 2) (2013) p. 1302.

[3] A Delobbe et al, EFUG 2011, http://www.imec.be/efug


The research leading to these results has received funding from the European Union Seventh Framework Program [FP7/2007-2013] under grant agreement No. 280566, project UnivSEM.

Fig. 1: Ultra-high resolution at low beam energies: TiO2 nanotubes at 1 kV (left) and polymer nanofibers at 20 V (right).

Fig. 2: Different approaches for cross section preparation: (a) basic top-down “staircase” strategy, (b) a single-pass cross sectioning with high redeposition, (c) a fully optimized cross section object with 2.5 times higher milling speed, the highest depth and the best shape without redeposition. Milled using 12 nA FIB current at 30 keV.

Fig. 3: Automated TEM lamella preparation utilizing innovative fast cross section milling approach.

Type of presentation: Poster

IT-13-P-2704 Focused ion beam patterning of boron-doped diamond electrodes: Influence of patterning parameters on the heterogeneous electron transfer behavior

Eifert A.1, Langenwalter P.1, Higl J.2, Lindén M.2, Nebel C.3, Mizaikoff B.1, Kranz C.1
1Institute of Analytical and Bioanalytical Chemistry, University of Ulm, Ulm, Germany, 2Institute of Inorganic Chemistry II, University of Ulm, Ulm, Germany, 3Fraunhofer Institute for Applied Solid State Physics, Freiburg, Germany
alexander.eifert@uni-ulm.de

Diamond with its large band gap of 5.49 eV can be transformed into a metallic-like semiconductor by doping with very high boron concentrations. This new electrode material outmatches many common electrode materials concerning potential window, chemical inertness and signal-to-noise-ratio. The superior chemical and physical properties in combination with the possibility to fabricate microelectrodes or microelectrode arrays for example with FIB [1, 2], renders them highly suitable for bioanalytical applications. The electrochemical behavior of BDD electrodes depend on a variety of parameters, such as doping level, defects, carbon impurities, crystal orientation, surface termination/modification and grain boundaries [3]. Patterning of BDD with FIB technology leads to damages due to the ion bombardment such as amorphization and hence changes the overall electrochemical behavior.
In this contribution we present the influence of different FIB patterning parameters on the electrochemical properties such as heterogeneous electron transfer rate constant and peak separation of the obtained BDD microelectrode arrays. Post fabrication electrochemical treatments will restore to a certain extent the electrochemical properties. Next to electrochemistry also Raman spectroscopy was applied to characterize the ion irradiated sample, which will also be presented.

References
[1] J. Hees, R. Hoffmann, A. Kriele, W. Smirnov, H. Obloh, K. Glorer, B. Raynor, R. Driad, N. Yang, O. A. Williams, C. E. Nebel, ACS nano 5 (2011) 3339–3346.
[2] A. Eifert, W. Smirnov, S. Frittmann, C. E. Nebel, B. Mizaikoff, C. Kranz, Electrochem. Commun. 25 (2012) 30–34.
[3] K. B. Holt, A. J. Bard, Y. Show, G. M. Swain, J. Phys. Chem. B 108 (2004) 15117–15127.


The Focused Ion Beam Center UUlm, which is supported by FEI Company (Eindhoven, Netherlands), the German Science Foundation (INST40/385-F1UG), and the Struktur- und Innovationsfonds Baden-Württemberg are greatly acknowledged. This work was supported by the project "Methoden für die Lebenswissenschaften P- LSMeth/23" funded by the Baden-Württemberg Stiftung.

Fig. 1: Rate constants calculated from recorded CVs after electrochemical treatment. Electrochemical measurements were recorded in a solution containing 10 mmol/l K4[Fe(CN)6], 0.016 mol/l Tween 20 and 0.1 mol/l KCl. The influence of dwell time dependency was investigated at an accelerating voltage of 30 kV and a beam current of 15 nA.

Fig. 2: Raman spectra recorded after different fabrication steps of the microelectrode arrays. Raman spectra were acquired using an excitation laser with 532 nm and 10 mW power. For detection a 100x objective with a 0.9 numerical aperture was used.

Type of presentation: Poster

IT-13-P-2724 Novel strategies towards faster and smoother FIB cross-sectioning

Hrnčíř T.1, Dluhoš J.1, Haničinec M.1, Hrachovec V.2
1TESCAN Brno, s.r.o., Libušina třída 1, Brno, Czech Republic, 2ON Semiconductor Czech Republic, 1. máje 2230, Rožnov pod Radhoštěm, Czech Republic
tomas.hrncir@tescan.cz

The Focused Ion Beam (FIB) and the Scanning Electron Microscopy (SEM) are essential techniques for many applications. FIB modifies the sample by the milling or, when accompanied by the Gas Injection System (GIS), by the deposition; SEM is used for imaging of resulting shapes at the high resolution, for charge compensation, or as a source of electrons for other analytical techniques. Common FIB-SEM instruments allow creation and imaging of a broad range of shapes. The most important shape is the cross section, which is used both for sectioning the sample and TEM sample preparation, by milling two cross sections at both sides of TEM sample [1]. Two parameters of the cross section are crucial – the fast milling rate and the high quality of the surface, with no damage or artifacts.

The milling rate depends on the sample material and on the beam incidence angle (Fig. 1). An optimized scanning strategy for the cross section is introduced, to keep the slope of the sample surface under the ion beam closer to the maximum rate. Apart from increasing the milling rate, this method also reduces the redeposition artifacts to avoid obstacles limiting the effective depth. The resulting cross section shape is greatly improved and around 2.5 times deeper when compared to classical stair shape (Fig. 2). This shortens the cross-sectioning and TEM sample preparation time significantly.

The common width and depth of the cross section, milled by Ga FIB, are ~10 µm. For larger cross sections with dimensions up to ~100 µm, Xe plasma FIB is much more efficient [2]. As the surface milled by Xe plasma FIB is often not as smooth as the surface milled by Ga FIB, the quality has to be improved by tilting the sample and milling from several directions [3]. This makes cross-sectioning more difficult and less accurate. To overcome these drawbacks, a novel method was developed to greatly improve the surface quality, while keeping the milling process easy and accurate. Commonly used eucentric sample stages allow the tilting only around the axis perpendicular to both FIB and SEM columns. The novel multi-tilt sample stage allows an additional tilting also in the plane of the cross section. Unlike the solution described in [3], where the eucentric stage was used, the proposed method allows to control the whole milling/tilting process by SEM imaging, which is essential for the precise end-point detection (Fig. 3). Greatly improved results were obtained on polycrystalline material samples and on semiconductor samples, like solder bumps (Fig. 3), packaged ball-bond Au wires (Fig. 4) and TSV.

References
[1] MHF Overwijk et al, J. Vac. Sci. Technol. B 11 (1993), 2021.
[2] T Hrnčíř et al., ISTFA Conf. Proc. (2012), 26.
[3] F Altmann, et al., ISTFA Conf. Proc. (2012), 39.
[4] http://www.srim.org


The research has been supported by the Technological Agency of Czech Republic *TE 01020233 (AmiSpec).

Fig. 1: Ga FIB milling rate on Si when changing the incident angle, modeled by SRIM [4]. Beam energy is 30 keV.

Fig. 2: Increasing cross-sectioning rate by optimizing the scanning strategy. A) Classical stair strategy, which gives a nice shape but it is slow; B) One-pass polishing, where the shape is corrupted by redeposition artifacts; C) Novel optimized strategy, with the nicest shape and the highest depth of the cross section.

Fig. 3: Xe plasma FIB polishing of the solder bump by alternating the stage tilt in the cross section plane. Arrows point in FIB direction and the polishing process is controlled by SEM imaging, allowing to stop the process in the center of the bump accurately.

Fig. 4: Cross section through the packaged semiconductor sample (PWM controller from ON Semiconductor) by Xe plasma FIB, which was practically impossible to perform by Ga FIB previously.

Type of presentation: Poster

IT-13-P-2741 Ultra-Fast Three Dimensional Microanalysis Using the Scanning Electron Microscope Equipped with Xenon Plasma Focused Ion Beam

Dluhoš J.1, Petrenec M.1, Peřina P.1, Reinauer F.2, Kopeček J.3, Hrnčíř T.1, Jiruše J.1
1TESCAN Brno, s.r.o., Brno, Czech Republic, 2EDAX-AMETEK GmbH, Wiesbaden, Germany, 3Institute of Physics ASCR, v. v. i., Prague, Czech Republic
jiri.dluhos@tescan.cz

The three dimensional microanalysis became a widely accepted analytical method during the past few years, gaining from the ability to describe the materials structure and composition as it exists in real components.

A high resolution scanning electron microscope (SEM) combined with a focused ion beam (FIB) is used for a high precision tomographical method based on FIB slicing and SEM observation of the slice. The FIB-SEM can be further equipped by analytical methods like energy-dispersive X-ray spectrometer (EDS) for elemental composition or electron backscattered diffraction analyzer (EBSD) for crystal orientation mapping, resulting in 3D microanalysis, i.e. 3D EDS and 3D EBSD [1].

However, the main limitation of this tomographic method so far has been the speed of data acquisition. This influences also the volume which can be analyzed in reasonable time of several hours. A novel solution for rapid 3D microanalysis is introduced in this paper using a high performance Xe+ plasma focused ion beam. Such a system allows FIB-SEM tomography on objects with dimensions of hundreds of microns easily within few hours [2, 3], newly combined also with high speed EDS and EBSD. Utilizing the recently developed “static position” approach [4], the speed of 3D EDS and 3D EBSD acquisition can be maximized.

The conventional Ga+ FIB systems have a limitation of maximum beam current of about 60 nA. For practical FIB-SEM tomography, the volume is limited to units or maximum several tens of microns. Contrary to that, the Xe+ plasma source FIB incorporated in the TESCAN’s FERA3 FIB-SEM allows ion beam currents up to 2 µA [5]. Together with the higher sputtering yield of accelerated Xe ions it reaches about 50 times higher milling rates than Ga+ based FIB.

Examples of 3D EDS and 3D EBSD obtained using the Xe+ plasma FIB-SEM are shown. The 3D EBSD was acquired on a Cu wire commonly used in microelectronics, see Fig. 1. A total volume of 100×100×30 µm was analyzed in about 2.5 hours. Data acquisition time was about 1 min for FIB slicing at 30 keV beam energy with 1 µA beam current and 4 min for 200×200 points EBSD map acquisition for each of the 30 slices.

The 3D EDS example in Fig. 2 shows the volume of 100×70×45 µm of Sn60Pb40 solder processed in about 2.3 hours. Acquisition of 45 slices was done at SEM beam energy 20 kV with lateral resolution 0.5 µm. EDS maps were stored with full spectra at each point. Elemental ROI maps using Sn Lα, and Pb Mα peaks were used for 3D visualization.

References:

[1] S Zaefferer, Book of abstracts EMAS Workshop (2009) p 123.

[2] T Hrnčíř et al, Microsc.Microanal. 19 (Suppl 2) (2013) p. 860.

[3] T Hrnčíř et al, 39th ISTFA Proceedings (2013) p. 27.

[4] Patent CZ 301692 (2009).

[5] T Hrnčíř et al, 38th ISTFA Proceedings (2012) p. 26.


The research leading to these results has received funding from the European Union Seventh Framework Program [FP7/2007-2013] under grant agreement No. 280566, project UnivSEM.

Fig. 1: 3D EBSD reconstruction of a copper wire used in microelectronic. The volume 100x100x30μm was analyzed and all data were acquired within 2.5 hours using FERA3 Xe+ Plasma FIB-SEM equipped with EBSD by EDAX . Crystal orientation was mapped using a color coded inverse pole figure (IPF-Z).

Fig. 2: 3D EDS reconstruction of a Sn60Pb40 solder the FERA3 Xe+ Plasma FIB-SEM. Volume of 100x70x45µm was analyzed and all data were acquired in about 2.3 hours. a) Reconstruction of a composite 3D elemental map using Sn Lα, and Pb Mα peaks. b) Separation of Pb phase reveals formation of dendritic structure.

Type of presentation: Poster

IT-13-P-2874 Targeted 3D-CLEM workflow on cultured cells

Steyer A. M.1, Schieber N. L.1, Duishoev N.1, Pepperkok R.1, Kirmse R.2, Schertel A.2, Schwab Y.1
1EMBL Heidelberg, Meyerhofstraße 1, 69117 Heidelberg Germany, 2Carl Zeiss Microscopy GmbH, Oberkochen
anna.steyer@embl.de

Correlative light and electron microscopy (CLEM) experiments uniquely provide a highly accurate link between the imaging of living cells and their 3D ultrastructure (4). However, CLEM generally suffers from a low throughput. The major hurdles include tracking the object throughout the different imaging modalities, the tedious procedures for sample preparation and the lack of automation in the data acquisition by electron microscopy. We aim to overcome these issues by developing an automated correlative workflow that links live cell imaging to high resolution 3D electron microscopy using FIB-SEM. The approach we will develop enables collecting statistically significant data from a large number of cells in a heterogeneous population, clearing the way to statistical analysis of important mechanisms in cell biology. Aiming to develop a flexible and versatile approach, we foresee applications of our method in other biological areas such as pharmacology, developmental biology or virology.

Currently, our workflow is developed and tested on cultured cells and employed to in a high-throughput study of Golgi apparatus organization. In a tight collaboration with the team of R. Pepperkok (EMBL Heidelberg), we visualize at the ultrastructural level how specific mutations influence the morphology of the Golgi apparatus. Using an automated light microscopy platform, large genome wide siRNA knockdown screens have led to the identification of key genes for the morphogenesis and function of the Golgi apparatus (3). Light microscopy was utilized to screen for specific phenotypes, fluorescent microscopy to find cells of interest and electron microscopy to look at the ultrastructure.

References:

1) Briggman K. L. and D. D. Bock (2012). "Volume electron microscopy for neuronal circuit reconstruction" Curr Opin Neurobiol 22(1): 154-161.

2) Colombelli J., et al. (2008). “A correlative light and electron microscopy method based on laser micropatterning and etching.” Methods Mol Biol. 457:203-213

3) Simpson J. C., et al. (2012). "Genome-wide RNAi screening identifies human proteins with a regulatory function in the early secretory pathway" Nat Cell Biol 14(7): 764-774.

4) Spiegelhalter C., et al. (2010). “From dynamic cell imaging to 3D ultrastructure: novel integrated methods for high pressure freezing and correlative light-electron microscopy” PLOS One 5(2): 203-213

5) Villinger C., et al. (2012). "FIB/SEM tomography with TEM-like resolution for 3D imaging of high-pressure frozen cells" Histochem Cell Biol 138(4): 549-556.


Many thanks to Team Schwab, Team Pepperkok and our collaboration partner Zeiss.

Fig. 1: Correlative light and electron microscopy workflow using a Focused Ion Beam-Scanning Electron Microscope i) 10x transmitted light image, j) 40x transmitted light image with laser cuttings around the area of interest, k) 40x fluorescent microscopy image, m) SEM image, o) Golgi network model acquired with a FIB-SEM

Type of presentation: Poster

IT-13-P-2936 Characterisation of the FIB Induced Damage in Diamond

Rubanov S.1, Suvorova A.2
1University of Melbourne, Melbourne, Australia, 2The University of Western Australia, Perth, Australia
sergey@unimelb.edu.au

Despite diamond’s extreme properties a TEM sample from diamond can be relatively easy prepared using FIB milling [1]. However, FIB milling results in formation of amorphous damage layers [2-3]. In addition, the rearrangement of broken diamond sp3 bonds into graphitic sp2 bonds is possible.
To study the initial stages of the damage formation (001) diamond sample was irradiated with 30 keV Ga ions with doses ranging from 2×1014 to 1016 ions/cm2. Continuous milling effect has been studied using rectangular trenches 4×4 µm2 and 2 µm deep formed using 100 pA beam current. The near surface regions of the trenches contained two types of damage: the bottom-wall, where the ion beam was normal to the surface and the side-wall, where it was at low angle to the trench walls.
For the dose 2×1014 ions/cm2 the point defect density was below amorphisation threshold and implanted region remains crystalline. For the dose 4×1014 ions/cm2 and above most of implanted region had defect density above amorphisation threshold and became amorphous (Fig. 1). The bottom part of the implanted layer remains crystalline but distorted due to still large number of point defects (Fig. 1b). EELS examination showed the presence of both sp2 and sp3 bonding in the damage corresponding to two different chemical states of carbon. The swelling of the amorphous damage layer shown in Fig. 1a is related to diamond’s sp3 bonds conversion to sp2 bonds with significant decrease in density. Using a mass balance calculation the density of the amorphous layers was determined. The density decreased with ion dose increased, and reached density of graphite (2.24 g/cm3, 80% sp2) for highest dose. For continues milling the thicknesses of the amorphous damage layers were measured to be 16 nm for side-walls and 44 nm for the bottom-walls (Fig. 2a). Concentration of Ga atoms was found to be 20 and 32 at.% for side-wall and bottom-wall damage layers. The thickness of the initial amorphous damage layers exponentially grows with ion dose (Fig. 2b) and has a tendency to saturate at the value which was measured for continuous milling.
The FIB induced damage in diamond comprises amorphous and crystalline components and is a result of complex process of ion penetration, swelling and sputtering. Amorphisation in diamond results in transition of sp3 bonds to sp2 corresponding to two different chemical states of carbon with accompanying density reduction. High concentration of Ga atoms is a result of accumulation of implanted atoms in damage layers due to short penetration depth and low sputtering yield in diamond.
References
[1]    S. Rubanov, AMTC Letters 2 (2010) p. 104.
[2]    J.F. Walker and R.F. Broom, Inst.Phys.Conf.Ser. 157 (1997)  p. 473.
[3]    S. Rubanov  and. P.R. Munroe, J. Microsc. 214 (2004) p. 213.


Fig. 1: TEM image showing damage in diamond after implantation of 4×1015 Ga/cm2 (a) and mag-nified TEM image of amorphous-diamond interface (b).

Fig. 2: (a) TEM image showing damage in diamond after continues milling; (b) the measured thick-ness of the amorphous damage layers as a function of the implantation dose.

Type of presentation: Poster

IT-13-P-2986 High quality site specific TEM cross section preparation of structured materials

Graff A.1, Hübner S.1, Simon-Najasek M.1, Altmann F.1
1Fraunhofer Institute for Mechanics of Materials / Center for Applied Microstructure Diagnostics (CAM), Halle, Germany
andreas.graff@iwmh.fraunhofer.de

A requirement for every TEM investigation is the preparation of electron transparent samples. TEM preparation by focused ion beams (FIB) is nowadays widely used to produce site specific sections from the region of interest. Various approaches for TEM sample preparation using FIB have been developed. The most flexible are lift-out techniques where a lamella is directly made from the original sample, transferred to a support grid and finally thinned to electron transparency.
In this paper, we demonstrate an improved workflow for TEM sample preparation by FIB for extremely thin and distortion-free lamellas. By using a special TEM grid with clamps and an active griper for sample transfer the welding and cutting procedure necessary for the standard lift out technique (Fig. 1 left) can be avoided [1]. An additional advantage is the fixation of the TEM lamella on both sides, thus reducing the bending of the lamella in the final stages of the thinning. The special design of the holder allows preparing the sample from different directions without damaging the clamps. A TEM grid holder construction with an additional rotation axis perpendicular to the sidewalls of the TEM lamella is presented where the angle of incidence can be varied independently for the front and the backside (Fig. 1 right). The result of this kind of preparation is an electron transparent window inside a mechanically stable bar of the specimen (Fig. 2). The transparent window has a trapezium shape with adjustable angles between 0 and 90 degrees. A possible variation of the angles can be used to control and reduce curtaining effects often appearing in structured multi-material systems. To control the residual thickness of the lamella inside the window, thickness measurements are performed during thinning by electron backscatter imaging using a cross beam instrument (NVision, Zeiss). Thus, the plan parallel shape and the thickness of the sample can be controlled during the final milling to reach a well-defined homogeneous lamella thickness [2]. TEM investigations of the samples prove the reduction of curtaining and wrapping in the ultra-thin transparent window (Fig. 3).
The workflow was successfully applied to different material systems which are discussed in the present contribution. The efficiency of the process and the high quality of the TEM samples are shown.
References:
1. Altmann F., et al., Microscopy and Microanalysis, Volume 17, Supplement S2, 2011, pp 626-627
2. Salzer R., Lechner L., Microscopy and Microanalysis, Volume 18, Supplement S2, 2012, pp 654-655


Fig. 1: Left: SEM image of the lamella transfer into the special clamp holder by active gripper. Right: Sample holder mounted on a Kleindiek RotTip to realize a second tilt axis.

Fig. 2: Left: Color coded thickness map of a semiconductor sample. Right: TEM overview of the sample. Curtaining is hardly visible.

Fig. 3: Left: Bright field image of the grain structure of a Tungsten plug. Right: HRTEM image of the silicon SiO2 interface in ultrathin area of the lamella (Thickness less than 40nm).

Type of presentation: Poster

IT-13-P-3033 About the accuracy of post-mortem alignment methods in FIB/SEM nano-tomography

Yuan H.1,2, Van de Moortèle B.2, Epicier T.1,3
1University of Lyon, MATEIS, umr CNRS 5510, Bât. Blaise Pascal, INSA Lyon, 69621 Villeurbanne Cedex, France, 2University of Lyon, LGLTPE, umr CNRS 5276, ENS Lyon, 69364 Lyon Cedex 07, France, 3University of Lyon, IRCELYON, umr CNRS 5256, 2, Av. A. Einstein, 69626 Villeurbanne Cedex, France
bertrand.van.de.moortele@ens-lyon.fr

Within the last decade, FIB/SEM tomography has become a commonly used tool for 3D microstructural investigation of materials at a sub-micrometer level [1, 2]. It consists in a true tomographic approach; in a Focused Ion Beam (FIB) microscope coupled with a Scanning Electron Microscope (SEM), the ion beam is used to mill the sample and prepare fresh surfaces which are successively imaged with SEM. The stack of acquired SEM images is then further aligned in order to restore the analysed volume of matter. By alignment, one intends both the drift correction during the acquisition itself [3] and a final post-mortem numerical alignment. The present contribution focuses on this final step: the post-mortem alignment before 3D reconstruction.

Basically, typical post-mortem alignment of the image stack is performed using cross-correlation based algorithms allowing the successive images to be aligned with respect to a reference image. Several methods have thus commonly used [1-5]: the most intuitive method is to define a Region of Interest (ROI) as a reduced frame in the field of view, with the option to locate this ROI near a lateral edge of the scanned area, or at the interface with the top surface of the sample in order to take benefit of an assumed fixed feature which will improve the accuracy of the alignment (fig.1). Further refinements introduced markers, i.e. holes or small trenches machined with the ion beam which will serve as markers facilitating the alignment. It will be demonstrated here that these alignment routines may easily fail, although the final reconstructions generally look correct. The principal issue of all methods mentioned above is due to the fact that they rely on the assumption that the markers, or the microstructure itself, are isotropic, intrinsically ‘fixed’ (with respect to the bulk sample) and not deformed, modified or erased during the 3D acquisition. This is generally not true and this will be demonstrated by dedicated test examples. As illustrated by fig.2 and 3, we will investigate alternative methods in order to improve the alignment and consequently the reliability of the reconstructed volumes. Among these methods, a promising one consists in a correlation with the top surface of the sample which can be reconstructed by stereoscopy prior to the 3D FIB-SEM acquisition and matter removal.

References

[1] L. Holzer,et al., J. of Microscopy 216 (2004), 84.

[2] M. Schaffer, et al., Ultramicroscopy 107 (2007), 587.

[3] H. Yuan in “EMC 2012: Proc.15th EMC, Vol. 2”, ed. D.J. Stokes and J.L. Hutchison, RMS: London, (2012), p. 135.

[4] H. Iwaia, et al., J. Power Sources 195 (2010), 995.

[5] M.D. Uchic et al., Ultramicroscopy 109 (2009), 1229.

[6] K. Lepinay et al., M and M, 19, (2013) 85.


We kindly acknowledge the financial support of Carl Zeiss SAS. Thanks are due to the CLYM (http://www.clym.fr) for the access to the Nvision 40 FIB instrument.

Fig. 1: Front view of the sample (serpentine) during the 3D-FIB analysis. Different ROIs (numbers, frames and arrows) are used for the post-mortem alignment (see fig. 2): ROI3 consists in FIB-milled lateral trenches filled with W (white contrast), whereas ROI4 corresponds to the top surface underlined by a W layer.

Fig. 2: Reconstruction of the top surface topography from a stereoscopic SEM analysis using 3 images taken at -10°, 10° (shown) and 0° with the help of the MEX program (Alicona SARL, Les Ulis, France).

Fig. 3: Superimposition of the top surface as reconstructed by the MEX© procedure (in green) and from the 3D-FIB analysis (in red) for the 4 ROI defined in figure 1. The best correspondence is by far obtained from the alignment based on the surface topography (case 4).

Type of presentation: Poster

IT-13-P-3258 FIB as Fabrication Tool for Advanced Analytical Infrared Sensing Schemes

Sieger M.1, Neusser G.1, Schaedle T.1, Kranz C.1, Mizaikoff B.1
1Institute of Analytical and Bioanalytical Chemistry, University of Ulm, 89081 Ulm, Germany
gregor.neusser@uni-ulm.de

Within the last decade, focused ion beam (FIB) technology emerged as a universal tool, for maskless prototyping via highly localized milling and deposition processes. FIB prototyping is particularly interesting for analytical devices such as novel optical sensing structures for the mid infrared (MIR) regime and functionalized scanning probes [1].

The MIR (3-12µm) spectral range is particularly interesting for biosensing applications, since it provides inherent molecular selectivity. MIR photons interact with most organic and inorganic molecules by excitation of vibrational und rotational modes [2-5].

Quantum cascade lasers and semiconductor thin-film waveguides facilitate highly sensitive optical biosensors for the MIR regime. Among those biosensors, the Mach-Zehnder interferometer (MZI) is one of the most promising sensor as it can be fully integrated in a lab-on-a-chip microsystem and provides high sensitivity.

The developed MIR-MZIs are chip-integrated solid-state devices based on GaAs/Al0.2Ga0.8As technology wave guide fabricated via conventional optical lithography and reactive ion etching (RIE). Since optical lithography is limited to a resolution of about 2 µm, the microfabricated devices were further structured via FIB milling for minimizing scattering losses at the Y-junction, and to increase the optical throughput.The radius of the Y-junction edges was reduced from about 2 µm to less than 100nm, thereby leading to a throughput enhancement of more than 30% of incident light in comparison to structures without FIB milling (Fig. 1)[6].

Grating couplers are another way for minimizing coupling losses, especially for wave guide with micro- to nanoscale dimensions. FIB milling of grating couplers is a reproducable and reliable strategy for launching radiation into a wave guide structure and improving the overall performance of the optical device [7].

References:

[1] C. Charlton, B. Temelkuran, G. Dellemann, B. Mizaikoff, Appl. Phys. Lett., 86, 194102 (2005).

[2] C. Charlton, M. Giovannini, J. Faist, B. Mizaikoff, Anal. Chem., 78, 4224-4227 (2006).

[3] C. Young, S.-S. Kim, Y. Luzinova, M. Weida, D. Arnone, E. Takeuchi, T. Day, B. Mizaikoff, Sens. &Act. B, 140(1), 24-28 (2009).

[4] X. Wang, S.-S. Kim, R. Roßbach, M. Jetter, P. Michler, B. Mizaikoff, Analyst, 137, 2322-2327 (2012).

[5] A.Eifert, W. Smirnov, S. Frittmann, C.E. Nebel, B. Mizaikoff, C. Kranz, Electrochem. Commun. 25 (2012) 30-34

[6] M. Sieger, F. Balluff, X. Wang, S.-S. Kim, L. Leidner, G. Gauglitz, B. Mizaikoff, Anal. Chem., 85, 3050-3052 (2013).

[7] T. Schädle, A. Eifert, C. Kranz, Y. Raichlin, A. Katzir, B. Mizaikoff Applied Spectroscopy, 67, 1057-1063 (2013).

[8] X. Wang, M. Sieger, B. Mizaikoff, Proc. SPIE 8631, Quantum Sensing and Nanophotonic Devices X, 86312M (2013)


The authors gratefully acknowledge support for parts of this work by the Kompetenznetz Funktionelle Nanostrukturen Baden Wuerttemberg, Germany. Finally, the Focused Ion Beam Center UUlm is thanked for providing prototyping and characterization facilities during the present studies.

Fig. 1: Figure 1. Scanning electron microscopy (SEM) images of GaAs/AlGaAs MIR-MZI waveguides with different opening angles α and a constant distance between the arms d: (a) top view of a Y-junction before, and (b) after FIB milling process [8].

Type of presentation: Poster

IT-13-P-3393 USING THE FOCUSED ION BEAM MICROSCOPE TO DEVELOP DIFFRACTION GRATINGS FOR QUANTUM WELL INFRARED PHOTODETECTORS

Rodrigues W.1, Schmidt W.1
1Microscopy Center from Federal University Minas Gerais
wesller@hotmail.com

Quantum well infrared photo detectors (QWIP) has several advantages when compared to other kinds of IR detectors, but it requires an optical coupling in order to work due to the quantum selection rule. These detector must has a Electrical Component of radiation (TE) parallel with structures grown direction. Diffraction gratings are one of the several types of optical coupling suitable to be implemented on these detectors. In this work we opt by diffractions gratings because the best optical coupling if compared with other techniques (1).

Due to the wavelengths of interest in this work (4,1um) and the refraction index the detector's material (n = 3.2) the features of the such gratings are almost sub-micrometric (2). The processes of optical lithography currently available in Brazil do not allow the fabrication of these diffraction gratings, essential to the functioning of these IR detectors, with the theoretical dimensions recommended, especially for near and mid infrared (3). The Focused Ion Beam microscope (FIB) allows the fabrication of these structures at resolutions more than enough (4) (5) (6).

The study and fabrication of these gratings are the objects of this work. FIB was used in the manufacture of gratings with several dimensions, from the theoretical dimensions until the optical lithographic dimensions (Figure 1) on QWIP mesa devices fabricated by usual optical lithography, enabling the investigation of the possible loss of electrical response of the QWIP´s detectors if they were textured with the available optical lithographic processes currently in Brazil.

To characterize the dimensions of arrays was used Atomic Force Microscopy (Figure 2). After fabrication diffraction gratings was deposited In each individual mesas a thin Gold film (180nm) to make the electric contacts by gold wire bonding (Figure 2a,b ).

This chip was assembly in header where the electric response the individual mesas will be measured and compared (Figure 3). The features of dimensions these gratings and their effect on the electric response of the devices will be presented.


Fig. 1: Figure 1 – SEM image of the photodetector mesas after fabrication of diffraction gratings

Fig. 2: Figure 2a – AFM Image of a diffraction grating. Figure 2b – SEM image of one already textured mesa with the deposited ohmic contact

Fig. 3: Figure 3 – SEM image of the chip assembled on the header and SEM image of the wire bonded mesas on the header

Type of presentation: Poster

IT-13-P-5975 LIFE CELL IMMUNOGOLD LABELLING OF RNA POLYMERASE II visualized by Focused Ion Beam Scanning Electron Microscopy (SEM-FIB)

Orlov I.1, Schertel A.2, Zuber G.3, Drillien R.1, Weiss E.4, Schultz P.1, Spehner D.1
1Integrative Structural Biology, IGBMC, UMR 7104, 67404 Illkirch, France, 2Carl Zeiss Microscopy GmbH, ASC, D-73447, Oberkochen, Germany, 3LCAMB UMR 7199, Faculty of Pharmacy, 67401 Illkirch, France, 4Biotechnology and Cell Signaling, UMR 7242, IREPS, 67412, Illkirch, France
spehner@igbmc.fr

The intracellular localization and dynamics of proteins involved in cellular processes are often studied in living cells at light microscopy resolution by monitoring proteins fused to fluorescent tags. These methods have nevertheless intrinsic drawbacks. First, the level of expressed fusion proteins is difficult to match with endogenous levels and since the endogenous protein is generally not extinct, the fusion protein acts on top of its native counterpart. The second restriction comes from the limited spatial resolution of light microscopy which, despite the spectacular development of super-resolution light microscope modalities, does not attain molecular dimensions.
Electron microscopy is an invaluable method to improve spatial resolution and to describe the cellular context of proteins of interest but relies on electron dense probes or reagents to detect the labeled macromolecule. The most successful and widely used method consists in conjugating primary or secondary antibodies to gold particles (Faulk and Taylor, 1971) which have a high electron scattering power and create an easily recognizable highly contrasted round shape.
Here we exploited the ability of cells to internalize macromolecules with a method named “live cell immunogold labeling” which takes advantage of lipid-based protein delivery agents compatible with cell viability to internalize the probes (Futami et al., 2012, Freund et al., 2013).
We used this method to label RNA polymerase II with electron dense labels suitable for EM localization studies and demonstrate for the first time that antibodies coupled to 0.8 nm ultrasmall gold particles (Van de Plas P. and Leunissen J.L., 1993) can enter the nucleus and be detected after amplification. Cells grew normally for more than 8 hours after probe uptake. The label was detected, after silver enhancement, by transmission electron microscopy and by scanning electron microscopy coupled to Focussed Ion Beam slicing (SEM-FIB)(Schroeder-Reiter E.et al.,2009). These methods open the new possibility to label nuclear or cytoplasmic antigens in living cells, and to immunolocate them in the whole cell volume using the SEM-FIB technology.
Faulk, W. P. & Taylor, G. M. Immunochemistry 8, 1081-1083 (1971)
Futami, M. et al. Bioconjug Chem 23, 2025-2031, doi:10.1021/bc300030d (2012)
Freund, G. et al. MAbs 5, 518-522, doi:25084 [pii]
Van de Plas, P. & Leunissen, J. L. Methods Cell Biol 37, 241-257 (1993)
Schroeder-Reiter E. et al., J. Struct. Biol., 165, Issue 2, (2009)


Type of presentation: Poster

IT-13-P-6027 Structural characterization of hybrid-organic nanocomposites via focus ion beam preparation and electron microscopy

Ankah G. N.1, Büchele P.2, Tedde S. F.2, Adam J.1, Torrents O.1, Pfaff M.1, Poulsen K.3, Koch M.1, Gimmler C.3, Schmidt O.2, Kraus T.1
1INM - Leibniz Institute for New Materials, Campus D2.2, 66123 Saarbruecken, Germany, 2Siemens AG Corporate Technology, Günther-Scharowsky-Str. 1, 91058 Erlangen, Germany, 3Centrum für Angewandte Nanotechnologie (CAN) GmbH, Grindelallee 117, 20146 Hamburg, Germany
GenesisNgwa.Ankah@inm-gmbh.de

Nanocomposites of conductive polymers and functional nanoparticles have recently been employed in applications such as photodetectors [1-3]. The particles convert light or x-rays into charge-carrier combinations that travel through a polymer blend to the contacting electrodes. In order to correlate device performances with the distribution of the nanoparticles in the organic polymer matrix, it is necessary to perform structural investigation at particle-level resolution. Focused ion beam (FIB) sample preparation is a prerequisite when specific regions are to be analyzed at such resolutions. In this work, composites of P3HT, PCBM and PbS nanoparticles or other inorganic nanoparticles were prepared by spray coating. Trenches and lamellae were prepared via FIB at several positions of the sprayed area. Their microstructures were analyzed using transmission electron microscopy (TEM) and scanning electron microscopy (SEM).
The distribution of nanoparticles inside the nanocomposite affects the properties of electronic materials. Conductive pathways, optical adsorption lengths and optical scattering depend on particle arrangement and affect sensor performance. Agglomerates and fully demixed particle phases make it harder for charge carriers to enter the polymer blend. Voids in the composite hinder transport and scatter light. We discuss the occurrence of such defects depending on processing.
FIB cuts through the soft polymer matrix that intersect hard nanoparticles can cause artefacts in the microstructure such as ridges, grooves, etc. The high energy ion bombardment may lead to local melting or the creation of amorphous layers. We discuss ion-beam related artefacts and their dependence on the preparation. We studied the extent of beam damage by comparing FIB cuts with samples prepared differently (e.g. using a microtome).


[1] Rauch T., Böberl M., Tedde S. F., Fürst J., Kovalenko M. V., Hesser G., Lemmer U., Heiss W., Hayden O., Nature Photonics 3, (2009) 332.
[2] Wagner B. K., Kang Z., Nadler J., Rosson R., Kahn, B., Proc. of SPIE 8373, (2012).
[3] Saunders B. R., Turner M. L., Advances in Colloid and Interface Science 138, (2008) 1.


Type of presentation: Poster

IT-13-P-6033 Structural and mechanical characterization of human dental tissues across multiple scales at the Oxford Multi-Beam Laboratory for Engineering Microscopy (MBLEM)

Sui T.1, Ying S.1, Lunt A. J.1, Baimpas N.1, Landini G.2, Korsunsky A. M.1
1MBLEM, Department of Engineering Science, University of Oxford, UK, 2The School of Dentistry, College of Medical and Dental Sciences, University of Birmingham, St Chad's Queensway, Birmingham B4 6NN, UK
alexander.korsunsky@eng.ox.ac.uk

The combination of focused ion and electron beams within the vacuum chamber of the Tescan LYRA3 XM instrument at the Oxford Multi-Beam Laboratory for Engineering Microscopy (MBLEM) provides a set of versatile tools for structural and mechanical characterisation of natural and engineered materials. In the present study we report the use of SEM for imaging, and FIB for ion beam milling and TEM lamella preparation of samples of human dentine and enamel, the two principal tissues used by nature to build teeth. We also pay particular attention to the dentine-enamel junction, the DEJ, a functionally and structurally graded interface that accommodates significant change in the degree of mineralisation (nearly two-fold), stiffness (nearly five-fold) and hardness (two-fold). We report using the combination of synchrotron X-ray diffraction with FIB ring-core milling to determine the variation of the lattice parameter of the mineral content, the nanocrystalline hydroxyapatite (HAp) particles. The use of FIB-DIC allows the separation of lattice parameter variation into chemical changes and mechanical effects (residual elastic strains) in the vicinity of the DEJ. In addition, the analysis of TEM lamella extracted from the sample made it possible to visualise the fractures observed in the dental tissues, and to explain the toughening mechanisms that operate at the nano-scale in these materials. The combination of synchrotron X-ray diffraction with FIB ring-core milling was also used to determine the variation of the hydroxyapatite (HAp) lattice parameter and elastic strains in the vicinity of the DEJ.


We acknowledge the support of colleagues at Tescan in UK and CZ: Ray Codd, Zora Strelcova, Jiri Dluhos and many others.

IT-14. Scanning probe microscopy and near-field microscopies

Type of presentation: Invited

IT-14-IN-2450 Advances in quantitative and three-dimensional mapping of soft matter by bimodal force microscopy

Garcia R.1
1Instituto de Ciencia de Materials de Madrid, CSIC
r.garcia@csic.es

Force microscopy is considered the second most relevant advance in Materials Science since 1960. Despite the success of AFM, the technique currently faces limitations in terms of three-dimensional imaging, spatial resolution, quantitative measurements and data acquisition times. Atomic and molecular resolution imaging in air, liquid or ultrahigh vacuum is arguably the most striking feature of the instrument. However, high resolution imaging is a property that depends on both the sensitivity and resolution of the microscope and on the mechanical properties of the material under study. Molecular resolution images of soft matter are hard to achieve. In fact, no comparable high resolution images have been reported for very soft materials such as those with an effective elastic modulus below 10 MPa (isolated proteins, cells, some polymers). Similarly, it is hard to combine the exquisite force sensitivity of force spectroscopy with molecular resolution imaging. Simultaneous high spatial resolution and material properties mapping is still challenging. This presentation reviews some of the above limitations and some recent developments based on the bimodal operation of the AFM to address and overcome them.

Recent References

E.T. Herruzo, A. P. Perrino and R. Garcia, Nat. Comm. 5, 3126 (2014)
R. Garcia and E. T. Herruzo, Nat. Nanotechnol. 7, 217-226 (2012).
E. T. Herruzo, H. Asakawa, T. Fukuma, R. Garcia, Nanoscale 5, 2678 (2013)
H. V. Guzman, A.P. Perrino, R. Garcia, ACS Nano 4, 3198 (2013)


We thank the financial support from the Ministerio de Economía y Competitividad, project CSD2010-00024 and the European Research Council, project 340177- 3DNanoMech.

Fig. 1: Bimodal AFM 3D images of solid-water volumes. a, 3D map of a mica-water interface. The stripes are associated to the presence of hydration layers. b, 3D map of a GroEL patch-water interface. The side view shows a slightly rough landscape with variations of the amplitude of about 1 nm.

Type of presentation: Invited

IT-14-IN-3211 Theory of Near-Field SEM: lateral resolution, scaling properties and compact equations

Xanthakis J. P.1
1National Technical University of Athens, Athens, 15700 Greece
jxanthak@central.ntua.gr

A new form of lenseless microscopy called near-field emission scanning electron microscopy (NFESEM) has been devised at ETHZ. In this form of microscopy the emitting tip is placed at a distance d of a few nm away from the sample (anode) with no focusing device in between. Besides the obvious advantages of price, NFESEM can perform DOS analysis (as STM) but also chemical identification by looking at the backscattered (or secondary reemitted) electrons.

The vertical resolution of this form of microscopy is as good as that of other forms of microscopy but the lateral resolution is about 3 nm at d=25nm and can be improved to 1 nm at smaller d. The latter constitutes a surprisingly good result considering the absence of a lens but it is not understood with conventional field emission theory. In this paper we present ab initio 3-dimensional WKB calculations applied to an ellipsoidal emitting tip that can explain the good lateral resolution capabilities of NFESEM. In particular, we show that the electron trajectories converge faster to the vertical direction compared to spherical emitting tips. This process begins in the classically forbidden (or tunneling) region where, contrary to accepted wisdom, electron paths may bend (see figure 1). As an end result of our calculations the beam spot size as a function of tip-anode distance d is obtained (figure 2)

The current-voltage (I-V) characteristics of this device are obviously d dependent. However for d>R=radius of curvature of the emitter, these characteristics show a remarkable property: all I-V curves fall onto each other when the applied voltage V is scaled (multiplied) by a single scaling function S(d), see figure 3. The explanation of this is rudimentary for image rounded linear potentials but for non-linear tunneling potentials U as those of nanoscopic emitters it demands a thorough investigation of the non-linear terms of U which we give.

Finally, in the course of our investigations, we have managed to produce- by strict mathematical proof- a generalized Fowler-Nordheim equation which is valid for any shape of surface. The accuracy of our equation – estimated by comparison to ab initio calculations- depends on R. For R>10nm it gives excellent results, for 5nm<R<10nm it gives satisfactory results and for R<5nm it does not work so well but still it gives much better results than the traditional FN equation, see figure 4.


This work would not have been possible without the substantial contributions of my former postdoc Dr Gerry Kokkorakis and my present PhD student Andreas Kyritsakis

Fig. 1: Comparison of beam-spot diameters by usuing straight line paths and curved paths calculated by a 3-Dimensional WKB method

Fig. 2: Lateral resolution of an ellipsoidal NFESEM emitter as a function of tip-anode distance d. The various curves are for different combinations of the large and small radii of the elliptical tip.

Fig. 3: Scaling properties of NFESEM voltage-current (I-V) characteristics.Main figure: scaled I-V curve, Inset: I-V curves for different d

Fig. 4: Comparison of Fowler-Nordheim plots calculated by our compact generalized FN equation and those obtained by ab initio WKB  calculations.

Type of presentation: Oral

IT-14-O-1627 Using SPM nanomanipulation to discover new materials and properties

Barboza A. P.1, Guimaraes M. H.1, Oliveira C. K.1, Massote D. V.1, Fernandes T. F.1, Archanjo B. S.2, Lacerda R. G.1, Batista R. J.3, Oliveira A. B.3, Mazzoni M. S.1, Chacham H.1, Neves B. R.1
1Universidade Federal de Minas Gerais, Belo Horizonte, Brazil, 2Inmetro, Duque de Caxias, Brazil, 3Universidade Federal de Ouro Preto, Ouro Preto, Brazil
berneves@gmail.com

In this work, Scanning Probe Microscopy (SPM) was employed for matter manipulation at the nanoscale in ambient conditions. More specifically, the SPM nanomanipulation potential is illustrated by two recent works of our group: in the first one, a new material – diamondol – is proposed and its realization is evidenced via SPM experiments [1]. In the second one, the SPM nanomanipulation was used to both induce and discover a general property of solid lubricants: a negative dynamic compressibility [2].
According to our ab initio calculations, the diamondol, or hydroxylated diamond, would be a new 2D material, formed via compression-induced diamondization of two, or more, graphene layers stabilized by hydroxyl ions (see Fig. 1a). The experimental observation of diamondol was carried out in a series of SPM experiments on mono-, bi-, and multilayer graphene in a controlled environment (humidity and temperature). Using electric force microscopy (EFM) to both inject and monitor charges and to apply pressure on the sample [3] (see Fig. 1b), we observed a pronounced inhibition on the charging efficiency for bilayer and multilayer flakes as the tip pressure increased, while monolayer charging was pressure-independent (Fig. 1c). The influence of the water content on the sample surface was tested in a series of charge injection experiments carried out at different temperatures (25°C and 120°C). The ensemble of experimental results can be well accounted for by the diamondol hypothesis, thus giving strong evidence of its experimental realization.
In the second study, a novel mechanical response of solid lubricants (few-layer graphene, h-BN, talc and MoS2) to the simultaneous compression and shear by a SPM tip is observed. The response is characterized by the vertical expansion of these 2D layered materials upon compression (see Figs. 2a-d). Such effect is proportional to the applied load, leading to vertical strain values (opposite to the applied force) of up to 150% (Fig. 2d). The effect is null in the absence of shear, increases with tip velocity, and is anisotropic. It also has similar magnitudes in these solid lubricant materials, but it is absent in single-layer graphene and in few-layer mica and Bi2Se3 (non-lubricant layered materials). We propose a physical mechanism for the effect where the combined compressive and shear stresses from the SPM tip induce dynamical wrinkling on the upper material layers, leading to the observed flake thickening (Fig. 2e). The new effect (and, therefore, the proposed wrinkling) is reversible in all four solid lubricants where it is observed.


Financial support from Fapemig, Capes, CNPq, Rede Nacional de Pesquisas em Nanotubos de Carbono, and INCT/Nano-Carbono is acknowledged.

Fig. 1: (a) Upon compression and in the presence of hydroxyl ions, two graphene layers undergo a sp3 re-hybridization, creating a layer of diamondol (b) SPM setup used to make and identify the diamondol. (c) Graph of the amount of injected charges as a function of the tip compression force which indicates diamondol creation.

Fig. 2: Contact Mode AFM images of a graphene flake under increasing tip loads: (a) 10 nN, (b) 195 nN, and (c) 391 nN. (d) Height profiles from (a), (b), and (c) evidencing the negative strain of graphene. (e) Upon compression and shear by the AFM tip, the topmost layers of a solid lubricant slide and wrinkle, creating a net expansion of the material.

Type of presentation: Oral

IT-14-O-1830 In-situ Dynamic SPM Studies of Organic Semiconductor Thin Film Growth on Silicon Oxide

Chiodini S.1, Straub A.1, Donati S.1, Borgatti F.1, Albonetti C.1, Biscarini F.2
1Consiglio Nazionale delle Ricerche, Istituto per lo Studio dei Materiali Nanostrutturati (CNR-ISMN), Bologna, Italy, 2Life Science Dept., Università di Modena e Reggio Emilia, Modena, Italy
schiodini@bo.ismn.cnr.it

The full text of the abstract is not available. Please contact the presenting author.


Type of presentation: Oral

IT-14-O-1831 Exploring the Local Photovoltaic Mechanisms in Organic Bulk Heterojunction Nanostructures by means of Scanning Probe Microscopy

Sebaihi N.1, Moerman D.1, Letertre L.1, Douhéret O.2, Lazzaroni R.1,2, Leclère P.1
1Laboratory for Chemistry of Novel Materials, Center for Innovation and Research in MAterials and Polymers (CIRMAP), Research Insitute for Materials Science and Engineering, University of Mons (UMONS), Mons, Belgium, 2Materia Nova R&D Center, Mons, Belgium.
philippe.leclere@umons.ac.be

Recent research and progress in organic photovoltaic (OPV) repeatedly insist on the importance of the molecular organization of the compounds forming the active bulk-heterojunction (BHJ) blends. The morphology of the blend has been to tremendously affect both the charge transfer at the donor-acceptor interface and the carrier transport to the electrodes. And still, for each material combination, much remains to be understood to fully assess its specific and ultimate morphology. For this purpose, high resolution characterization methods are of primary interest to locally depict the different electrical mechanisms ruling the photovoltaic process. Conductive Atomic Force Microscopy (C-AFM) and Kelvin Probe Force Microscopy (KPFM) have already proven to be of significant help to yield nanoscale two-dimensional mapping of electrical properties.
C-AFM and related PeakForce TUNA emerged as powerful technique to electrically evidence phase separation in blends. An additional key feature lies in local I-V curve providing useful information about the charge transport mechanisms within the materials forming the blends. Quantitative measurements leading to local determination of hole mobility have already been reported. It appears that upon illumination the technique has shown to be sensitive to photocurrent. With photoconductive-AFM (pc-AFM), a dedicated external calibrated module  has been recently introduced allowing full quantitative mapping of photovoltaic mechanisms. In this study, we will present and discuss the obtained results on two well-known samples: (i) poly(3 hexyl thiophene)(P3HT):[6,6]-phenyl-C61-butyric acid methyl ester (PCBM) and (ii) poly[2-methoxy-5-(3′,7′-dimethyloctyloxy)-1,4-phenylenevinylene])(MDMO-PPV):[6,6]-phenyl-C61 -butyric acid methyl ester (PCBM).
KPFM is mainly used to delineate phase separation and potential variations at interfaces. Upon illumination, photovoltage can also be evidenced. Yet, in organic electronics, KPFM still suffers from harsh operating environment (ultra-high vacuum and low temperature) to reach satisfactory spatial resolution and lacks for modeling for quantitative measurements. Augmenting KPFM with the PeakForce TappingTM technology allows a drastic improvement in the spatial resolution for KPFM measurements in ambient conditions. With the additional external calibrated illumination module, mapping of photovoltage in BHJ blends can be obtained, opening the doors of local characterization of charge transfer at donor-acceptor interfaces, where crucial processes are occurring in photovoltaic devices.


Research on conjugated polymers in Mons is supported by the Science Policy Office of the Belgian Federal Government (BELSPO PAI 7/05), the OPTI2MAT Excellence Program of Région Wallonne, and FNRS-FRFC. Ph.L. is Research Associate of F.R.S.-FNRS (Belgium).

Type of presentation: Oral

IT-14-O-2891 Gaseous nanobubbles on immersed surfaces: Properties and imaging by AFM in situ and ex situ

Janda P.1, Tarábková H.1
1J. Heyrovský Institute of Physical Chemistry ASCR v.v.i., Prague, Czech Republic
pavel.janda@jh-inst.cas.cz

Hydrophobic surfaces immersed in water are often densely occupied by gaseous nanodomains – nanobubbles and nanopancakes, with the size ranging between 10 – 100 nm, but it can exceed 1000 nm. The first direct proof of their existence came in the year 2001 in a form of in-situ atomic force microscopic (AFM) image made by J.W.G. Tyrrell and P. Attard [1] . Nevertheless, nanobubbles were still considered as artifacts, mostly due to their rather peculiar behavior, seeming disobeyance of Young Laplace law and hence nonexistence of plausible physical model. As various fields became affected by nanobubble existence, including interfacial physical chemistry, biophysics, microbiology, material sciences, nanofluidics, heterogeneous (electro) catalysis, immersion lithography and others, etc, nanobubbles started to attract growing attention. Their full impact is however, still to be disclosed.
Recent work performed in our laboratory [2] [3] [4] revealed the nanobubble ability to significantly rearrange solid surfaces which they adhere to. Nanobubbles can act as surface nanopatterning elements, changing in a significant manner its nanomorphology even at very mild conditions - in deionized water, at room temperature and pressure variations not exceeding 10 kPa. Besides nanobubble imaging by AFM in situ and distinguishing nanobubbles from solid nano-objects, we are presenting our novel technique, which can be called „nanobubblegraphy“[4], due to its ability to allow ex situ recognition and ex-post imaging of nanobubbles on dried surface as imprints developed after relatively short (sub-second) exposition in polymeric matrix (Figs 1 - 4).
Nanobubble properties and their utilization for imaging in situ and ex situ are discussed in relation to current physical models.

References

[1] J. W. G. Tyrrell and P. Attard: Phys. Rev. Lett. 87, 176104 (2001)
[2] P. Janda, H. O. Frank, Z. Bastl, M. Klementová, H. Tarábková, L. Kavan, Nanotechnology 21 (2010) 095707 (7pp)
[3] Viliam Kolivoška, Miroslav Gál, Magdaléna Hromadová, Štěpánka Lachmanová, Hana Tarábková, Pavel Janda, Lubomír Pospíšil, Andrea Morovská Turoňová: Colloids and Surfaces B: Biointerfaces 94 (2012) 213– 219
[4] H. Tarábková, P. Janda: J. Phys.: Condens. Matter 25 (2013) 184001


Acknowledged project support GACR P208/12/2429

Fig. 1: AFM image of polystyrene film as received by spin-coating. 

Fig. 2: AFM image of polystyrene film after exposition to nanobubbles in deionized water. Net-like nanopattern corresponds to 2D nanofoam imprinted into polystyrene matrix.

Fig. 3: Profile analysis of polystyrene film as received by spin coating (data from Fig. 1)

Fig. 4: Profile analysis of polystyrene film after exposition to nanobubbles in deionized water (data from Fig. 2)

Type of presentation: Poster

IT-14-P-1654 Investigating electrical charged samples by scanning probe microscopy: the influence to atomic force microscopy and magnetic force microscopy image artifacts.

Costa C. A.1, Lanzoni E. M.1, Piazzetta M. H.1, Galembeck F.1, Deneke C. F.1
1Laboratório Nacional de Nanotecnologia (LNNano), Centro de Pesquisa em Energia e Materiais (CNPEM), Campinas, São Paulo, Brazil.
carlos.costa@lnnano.cnpem.br

The electric state of a surface is of great importance for its chemical and physical properties, e.g. bounding of molecules, charge transfer or dielectric properties. Various scanning probe microscopy techniques based on atomic force microscopy (AFM) map electric surface potentials, electrostatic forces as well as the topography to distinguish electrostatic from van der Waals forces. However, the signal acquired from van der Waals forces can also show artifacts arising from electrostatic forces. Even so such effects are of great importance to interpret image obtained from charged surfaces, the effect is not discussed in detail in literature.

In this work, AFM artifacts resulting from electrical charged surfaces are investigated. In a detailed and systematic study, the influence of an electric field gradient above sample to topography, phase and magnetic force microscopy (MFM) images is investigated. Images were acquired with a commercial AFM using a lithographical patterned Kelvin force microscopy (KFM) calibration sample (Fig. 1). Our results show that electrical charges give rise to a signature in topography (Fig. 2) and phase signal. In order to minimize these artifacts, they are studied in regard of various acquisition parameters. We find that either using a low relative set point or high free vibration amplitude during images acquisition reduces the influence to the AFM measurements. Both approaches can sufficiently negate the effect by increasing the tip/sample interaction (either by getting the tip closer to the sample surface or by larger tip vibration amplitudes). As a trade off to these approaches, the sensitivity to topography features is reduced.

Finally, commercial metalized MFM cantilevers are studied in regard their sensitivity to electrical charge present on the sample surface. We observe the appearance of a MFM contrast for the non-magnetic KFM test structure (Fig. 3) for such conditions. The electrical charges give rise to a MFM signature indistinguishable from a magnetic signature exhibiting a strong correlation to KFM images obtained from our sample. The results indicate that great care has to be taken in the interpretation of topographic, phase contrast or magnetic images, when electrical fields are present on the sample surface.


This research was financially supported by Ministério da Ciência, Tecnologia e Inovação (MCTI) - Brazil.

Fig. 1: Illustration of KFM test sample consisting of interdigitated Al stripes on a glass substrate.

Fig. 2: AFM topography images (a) 0V, (b) 5V and (c) 10V applied between the finger pairs.

Fig. 3: (a)Topographic image and “MFM” phase shift for our KFM test sample (b) 0 V and (c)5V applied between the finger pairs.

Type of presentation: Poster

IT-14-P-2107 High-Sensitivity Imaging in Liquid by Torsional Resonance Mode Atomic Force Microscopy utilizing Lorentz Force Actuation

Yang C. W.1, Ding R. F.1, 2, Lai S. H.1, Liao H. S.1, Lai W. C.1, Huang K. Y.2, Chang C. S.1, Hwang I. S.1
1Institute of Physics, Academia Sinica, Taipei, 11529, Taiwan, 2Department of Mechanical Engineering, National Taiwan University, Taipei, 10617, Taiwan
yangcw@phys.sinica.edu.tw

Atomic force microscopy (AFM) [1] is a widely used technique for characterizing the structures and mechanical properties of material surfaces. However, operation in aqueous solution is still very challenging because the force sensitivity of dynamic AFM modes in liquid is usually much reduced compared with that in air or in vacuum. It has been shown that torsional resonance (TR) mode of vibrating cantilever is only affected by the tip-sample lateral force gradient, and not sensitive to the long-range normal forces [1]. The resonance characteristics (amplitude, phase, or frequency) start to change only when the tip gets in contact with the sample, which allows clear detection of the contact point and maintaining a soft contact between the tip and the sample. However, up to now, only very few types of cantilevers can be excited with pure torsional resonance in water [1,2].

Here we present a design based on Lorentz force induction to excite pure torsional resonances of different types of micro-cantilevers in air and in water [3]. Figure 1 shows the schematic of a rectangular cantilever actuated by the Lorentz force. An oscillating current passes through a cantilever which is mounted near two permanent magnets. The induced Lorentz force is equal and opposite on the two cantilever beams, resulting in a net torque to excite the torsional resonance [4]. With this actuation, pure torsional resonance of different types of cantilevers can be excited in air as well as in water, as shown in figure 2. To demonstrate the imaging capability, phase-modulation torsional resonance (PM-TR) mode is employed to resolve fine features of purple membranes in a buffer solution, as shown in figure 3. Most importantly, as shown in figure 4, force-versus-distance curves using a relatively stiff cantilever (k~ 40 N/m) can clearly detect hydration layers at a water-mica interface, indicating high force sensitivity of the torsional mode. Thus, the high resonance frequencies and high quality-factors for the tosional mode may be of great potential for high-speed and high-sensitivity imaging in aqueous environment. Moreover, this mode has a good potential and application for in-plane material characterization.

[1] Yang C. W.; Hwang I. S.; Nanotechnology 2010, 21, 065710
[2] Mullin N.; Hobbs J., Appl. Phys. Lett. 2008, 92, 053103.
[3] Yang C. W et al, Nanotechnology 2013, 24, 305702
[4] Byeonghee Lee et al, Nanotechnology 23, 055709 (2012).


This research is supported by the National Science Council of ROC (NSC99-2112-M-001-029-MY3 and NSC101-2221-E-002-022) and Academia Sinica.

Fig. 1: Schematics of Lorentz force actuation

Fig. 2: Torsional resonance curves of different cantilevers in air and in water

Fig. 3: AFM images of purple membrane taken with PM-TR mode in a buffer solution.

Fig. 4: Force curves of the phase lag and the amplitude of a vibrating cantilever versus the tip–sample distancemeasured on a freshly cleaved mica surface in DI water

Type of presentation: Poster

IT-14-P-1878 Thermal conductivity reduction measurement on Si and P3HT nanowires : diameter size effect

Grauby S.1, Munoz Rojo M.2, Martin Gonzalez M.2, Dilahire S.1
1Univ. Bordeaux, LOMA, UMR 5798, F-33400 Talence, France, 2Instituto de Microelectrónica de Madrid, IMM-CSIC, 8 PTM 28760 Tres Cantos, Madrid, Spain
stephane.grauby@u-bordeaux.fr

Nanostructuration has induced a renewed interest for thermoelectric materials whose performance can be evaluated through their figure of merit ZT=(S2σT)/λ where S is the Seebeck coefficient, σ the electrical conductivity and λ the thermal conductivity. Indeed, in a nanostructured material, the dimension reduction could possibly induce a thermal conductivity reduction and consequently an increase of their figure of merit. This can be partially ascribed to phonon boundary scattering appearing when one dimension of the nanostructured material becomes smaller than the phonon mean free path. The nanomaterial behaves than as a phonon glass and an electron crystal.

The work presented here deals with the study of the variation of the thermal conductivity of nanowires when reducing their diameter size due to confinement effects. The thermoelectric device is actually made of nanowires embedded in a matrix. We have studied two different kinds of nanowires with varying diameters: on one hand inorganic semiconductor Si nanowires in a SiO2 silica matrix, on the other hand poly(3-hexylthiophene) (P3HT) nanowires, an organic semiconductor polymer, which have been proven to have good thermoelectric properties[1], in an alumina matrix.

Measuring the thermal conductivity of individual nanowires embedded in a matrix is still challenging and nowadays there are not many techniques able do it[2]. Nevertheless, we have developed a technique based on an AFM associated to a thermoresistive tip and called 3ω-Scanning Thermal Microscopy (3ω-SThM). The thermoresistive tip is used both as a heater and a sensor. A current passing through it heats the tip. Depending on the thermal conductivity of the scanned material, the heat flux passing from the tip to the scanned sample varies, inducing a tip temperature variation. Then, the tip resistance changes, which induces a tip voltage variation. As a consequence, measuring the tip voltage variation enables to deduce the material thermal conductivity[3]. This mode enables to simultaneously obtain a topographical image and a thermal conductivity contrast image.

We show that a thermal conductivity reduction is observed when reducing the diameter of the nanowires for both silicon and P3HT nanowires. The thermal conductivity is reduced by 4 for P3HT nanowires when their mean diameter is reduced from 350nm to 120nm and up to a factor 10 for silicon nanowires when the mean diameter is reduced from 300nm to 50nm.

References:

1. C. Bounioux et al, Energy & Environmental Science, 2013, 6, 918-925.

2. M. Muñoz Rojo et al, Nanoscale, 2013.

3. E. Puyoo et al, J. Appl.Phys., 2011, 109, 024302-024309.


Fig. 1: Si nanowires: Scanning Electron Microscopy images respectively before (a) and after (b) encapsulation in the SiO2 matrix; 5µmx5µm 3ω-SThM images (c) topographical image, and (d) thermal conductivity contrast image.

Type of presentation: Poster

IT-14-P-1892 Conduction and Dissipation of Electrostatic Charges: Fundamental Study by Scanning Probe Microscopy

YIN J.1, NYSTEN B.1
1IMCN-Institute of Condensed Matter and Nanosciences (BSMA - Bio and Soft Matter Division), Université Catholique de Louvain, Belgium
jun.yin@uclouvain.be

Static electricity is a well-known and often observed physical phenomenon. It can cause dangerous problems in many applications, such as dust filters, chemistry, sophisticated electronics, cable insulation, charge based data storage, etc. Although many contributions have been done, the understanding of conduction and dissipation behaviors, charge transfer to, and retention on, surface or charges leakage over surfaces is far from being completed. Since most of studies are at the macroscopic scale, a microscopic and systematic study is of importance to understand these phenomena. Thanks to the development of scanning probe microscopy, a number of new electrical modes using a conductive probe have been developed and used to characterize the different microscopic electrical properties, such as Current-Sensing AFM (CS-AFM), Kelvin Probe Force Microscopy (KPFM). The objective of our project is to make a microscopic, detailed and systematic study of the phenomena of electrification, charge, discharge, conduction and dissipation mechanisms of electrostatic charges. The materials studied are two different kinds of fibers used in antistatic filters: polyester fiber and stainless steel conductive fiber commercially named Bekinox® fiber.

Surface properties of stainless steel conductive fiber are first studied (Fig.1). The surface topography and surface roughness are studied by standard AFM, the surface electrical resistance distribution is measured by CS-AFM, and the surface potential distribution is measured by KPFM. I-V spectroscopy is performed statistically to investigate the different charge transport mechanisms from different surface states. Second, the mechanisms responsible for charge conduction and dissipation between two fibers are studied. It can be noted that when a conductive fiber is put in non-galvanic contact with an other polarized conductive fiber, a dynamic behavior of surface potential variation can be measured by KPFM on the first conductive fiber (Fig.2). The quantified charging and discharging curves can be fitted to obtain relaxation times. Different contact systems, including different types of fibers, galvanic and non-galvanic contacts, are investigated systematically in order to deeply understand the mechanisms of conduction and dissipation.


This research is supported by FRIA (Fonds pour la Formation à la Recherche dans l’Industrie et dans l’Agriculture) of FRS-FNRS (Fonds de la Recherche Scientifique).

Fig. 1: Topography, surface current distribution, surface resistance distribution images obtained at the same location on Bekinox® fiber (a) topography, (b) (c) (d) electrical current distribution at 2 V, -2 V and 2 V again, (e) (f) (g) electrical resistance distribution at 1 V, 3 V, 4 V and 5 V, respectively. Image size is 5x5 µm².

Fig. 2: Topography and successive KPFM images on  Bekinox® fibers in non-galvanic while changing the applied voltage from 0 to 8 V (a) topography, (b) (c) successive KPFM images, the white arrows present the scanning direction (d) (e) two profiles from KPFM images, the average value of profile (d) is higher than profile (e).

Type of presentation: Poster

IT-14-P-2022 A ferrule-top optomechanical probe to collect topographic and near field information about a sample at the nanometer scale

van Hoorn C. H.1, Chavan D. C.1, Tiribilli B.2, Margheri G.2, Mank A. J.3, Ariese F.1, Iannuzzi D.1
1Faculty of Sciences, LaserlaB, Vrije Universiteit, Amsterdam, The Netherlands, 2Institute of Complex Systems, National Research Council, Sesto Fiorentino, Italy, 3Philips Innovation Services, Eindhoven, The Netherlands
c.h.van.hoorn@vu.nl

Atomic force microscopy (AFM) is an excellent technique for obtaining high-resolution images of the topography of a sample. The impact of AFMs in nanotechnology could be even more significant if the imaging capabilities were supported by an accurate mapping of the optical field in the close proximity of the surface. Scanning near field optical microscopy (SNOM), for example, has been shown to be able to combine AFM imaging with the possibility to collect optical information at the nanoscale. Yet, because of the complexity of its working principle, SNOM has been so far only used in specialized research laboratories and has been mostly limited to the analysis of surfaces in dry environments. To solve this limitation, in a previous work [1] some of us have proposed an all-optical device obtained by carving a tipped cantilever on top of an optical fiber. The opposite end of the fiber can be coupled to a readout system that was shown to be able to detect any tiny movement of the cantilever and to collect the SNOM signal coming from a prism illuminated under total internal reflection conditions. Here, we present another similar all-optical probe for AFM+SNOM imaging. The probe is based on ferrule-top technology [2-4], which relies on the possibility to fabricate a small cantilever on top of a ferruled fiber. This design keeps the overall advantages of the previous version (small dimensions, ease of use, easy integration in harsh environments) while significantly reducing the fabrication costs. Using this probe, we were able to obtain the SNOM profile and, simultaneously, the topographic image of a test grating (NT-MDT SNG01) kept in air and illuminated from below via an evanescent field. The AFM height resolution and the SNOM lateral resolution resulted to be comparable to conventional SNOM systems. Interestingly, measurements were repeated in water, with no major deterioration of the overall performance. This result paves the way for AFM+SNOM imaging on biological samples.

References:

1. B. Tiribilli, G. Margheri, P. Baschieri, C. Menozzi, D. C. Chavan, D. Iannuzzi, Journal of Microscopy 2011, 242, 10–14.

2. G. Gruca, K. Heeck, J. H. Rector, and D. Iannuzzi, Optics Letters 2013, 38, 1672.

3. G. Gruca, D. C. Chavan, J. H. Rector, K. Heeck, and D. Iannuzzi, Sensors and Actutators 2013, A190, 77.

4. D. C. Chavan, G. Gruca, S. de Man, M. Slaman, J. H. Rector, K. Heeck, D. Iannuzzi, Review of Scientific Instruments 2010, 81, 123702.


This work was funded by the European Research Council and by NanonextNL.

Fig. 1: Fabrication procedure of a ferrule-top probe. A glass ribbon is glued on top of a ferrule (a, b). The length of the cantilever is adjusted using a laser ablation machine (c). A tipped fiber is glued to the cantilever and the cantilever is released by focussed ion beam milling (FIB). Finally, a fiber is inserted into the borehole (d).

Fig. 2: A schematic view of the experimental setup. A blue (473 nm) laser beam impinges the top surface of a prism to create an evanescent field. The probe was scanned over the sample surface, in order to obtain a topographic and SNOM image simultaneously.

Type of presentation: Poster

IT-14-P-2058 High-sensitivity high-resolution elemental 3D analysis by in-situ combination of SIMS and SPM

Fleming Y.1, Eswara Moorthy S.1, Wirtz T.1, Gerard M.1, Gysin U.2, Glatzel T.2, Meyer E.2, Maier U.3
1Department “Science and Analysis of Materials” (SAM), Centre de Recherche Public – Gabriel Lippmann, 41 rue du Brill, L-4422 Belvaux, Luxembourg, 2Department of Physics, University of Basel, Klingelbergstr. 82, Basel, Switzerland, 3Ferrovac GmbH, Thurgauerstr. 72, CH-8050 Zürich, Switzerland
wirtz@lippmann.lu

Owing to its excellent sensitivity, its high dynamic range and its good depth resolution, Secondary Ion Mass Spectrometry (SIMS) constitutes an extremely powerful technique for analyzing surfaces and thin films. In recent years, considerable efforts have been spent to further improve the spatial resolution of SIMS instruments. As a consequence, new fields of application for SIMS, e.g. nanotechnologies, biology and medicine in particular, are emerging [1-2].

State-of-the-art SIMS instruments allow producing 3D chemical mappings with excellent sensitivity and spatial resolution. However, several important artifacts arise from the fact that the 3D mappings do not take into account the sample’s surface topography. The traditional 3D reconstruction assumes that the initial sample surface is flat and the analyzed volume is cuboid. The produced 3D images are thus affected by a more or less important uncertainty on the depth scale and can be distorted. Moreover, significant field inhomogeneities arise from the surface topography as a result of the distortion of the local electric field. These perturb both the primary beam and the trajectories of secondary ions, resulting in a number of possible artifacts, including shifts in apparent pixel position and changes in intensity.

In order to obtain high-resolution SIMS 3D analyses without being prone to the aforementioned artifacts and limitations, we developed an integrated SIMS-SPM instrument, which is based on the Cameca NanoSIMS 50 [2]. This instrument, an in-situ combination of sequential high resolution Scanning Probe Microscopy (SPM) and high sensitivity SIMS, allows topographical images of the sample surface to be recorded in-situ before, in between and after SIMS analysis. Hence, high-sensitivity high-resolution chemical 3D reconstructions of samples are possible with this extremely powerful analytical tool [3-4].

In addition, this integrated instrument allows a combination of SIMS images with valuable AFM (Atomic Force Microscopy) and KPFM (Kelvin Probe Force Microscopy) data recorded in-situ in order to provide an extended picture of the sample under study. The known information channels of SIMS and AFM/KPFM are thus combined in one analytical and structural tool, enabling new multi-channel nanoanalytical experiments. This opens the pathway to new types of information about the investigated nanomaterials.

This paper will present the prototype instrument with dedicated software, its performances and some typical examples of application.

References:

[1] Y. Fleming et al., Appl. Surf. Sci. 258 (2011) 1322-1327

[2] T. Wirtz et al., Surf Interface Anal. 45 (2013) 513-516

[3] T.Wirtz et al., Rev. Sci. Instrum. 83 (2012) 063702

[4] C. L. Nguyen et al., Appl. Surf. Sci. 265 (2013) 489-494


Fig. 1: Combined SIMS-SPM 3D reconstruction of an Al (100nm) / Si sample exposed to a plasma streamer (Field of view: 15x15 µm2): (a) Al- signal (b) Si- signal

Fig. 2: PS/PMMA blend (Field of view: 22.3x17.3 µm2): (a) Combined SIMS-SPM 3D reconstruction of the 12C- secondary ion signal. (b) Combined SIMS-SPM 3D reconstruction of the 16O- signal, which is characteristic of PMMA [3].

Type of presentation: Poster

IT-14-P-2144 Imaging nanostructures of nitrogen/oxygen molecules at HOPG-water interfaces with different atomic force microscopy modes

Hwang I.1, Yang C.1, Lu Y.1, Fung C.1, Ko H.1
1Institute of Physics, Academia Sinica, Taipei, 11529, Taiwan
ishwang@phys.sinica.edu.tw

Nanobubbles, cap-shaped soft nanostructures, and micropancakes, quasi-2D layered structures, have been reported at the interfaces between hydrophobic solid surfaces and water based on atomic force microscopy (AFM) studies [1]. Previous studies have indicated that these interfacial structures contain gases because they are formed under water saturated or supersaturated with gases. They were considered as novel gaseous states by many researchers. However, there are several puzzles about them, such as the high stability, the nature, the rather flat morphology, etc.


We have investigated these interfacial structures on highly ordered pyrolytic graphite (HOPG) surfaces in pure water with different atomic force microscopy (AFM) modes, including the frequency-modulation (FM), the tapping, and PeakForce techniques. The FM mode provides more accurate measurement of the surface profiles of nanobubbles than the other two imaging modes (Fig.1). The height obtained with PeakForce mode is smaller than the true height of nanobubbles due to a snap-in when the tip touches a nanobubble, as shown in the force vs the tip-sample separation curve (Fig. 2a). This is because a positive peak force is required to achieve stable imaging. The resonance frequency shift vs the tip-sample separation curve (Fig. 2b) shows a sharp increase in the resonance frequency when the tip touches a nanobubble, thus the snap-in has little effect on the height measurement in the FM mode. Similar force curves are seen on micropancakes. Combining AFM images obtained with these modes, models for nanobubbles and micropancakes are proposed, which can provide a better explanation for the high stability of these interfacial structures.

[1] Seddon J. R. T. and Lohse D.; J. Phys: Condens. Matter 2011, 23, 133001.


This research is supported by the National Science Council of ROC (NSC96-2628-M-001-010-MY3 and NSC99-2112-M-001-029-MY3) and Academia Sinica.

Fig. 1: Topographic imaging of nanobubbles on HOPG in DI water taken with (a) the FM mode, (b) PeakForce mode, (c) the tapping mode (TM). (d) Horizontal height profiles across the center of nanobubble 3 taken with the FM PF, and TM modes.

Fig. 2: Approaching force curve measured on nanobubbles at HOPG-water interfaces. (a) Force vs the tip-sample separation curve. (b) Resonance frequency shift vs the tip-sample separation curve.

Type of presentation: Poster

IT-14-P-2422 Instrument induced artifacts in scanning probe microscopy.

Lanzoni E. M.1, Costa C. A.1, Barboza V. A.1, Galembeck F.1, Deneke C. F.1
1Laboratório Nacional de Nanotecnologia (LNNano), Centro de Pesquisa em Energia e Materiais (CNPEM), Campinas, São Paulo, Brazil.
evandro.lanzoni@lnnano.cnpem.br

In the last three decades, scanning probe microscopy (SPM) techniques have been established as the major way to directly probe the 3 dimensional structure of a sample surface. The images are acquired by accurate movements of a sharp tip (probe) above the sample surface, controlled by a scanning electronic. For topography images van der Waals interaction between tip and sample is generally used as feedback mechanism. The  two major classes of topography images artifacts are: the tip-sample convolution resulting in a broadening of the observed structures as well as digitalization artifacts arising from the analog-digital conversion carried out during image acquiring.

In this work, we carry out a detailed analysis of the tip-sample convolution artifacts to topographic image formation in regard of the finite resolution implied by the analog-digital conversation. We discuss possible ways to identify these artifacts and wrote a software module to identify them in obtained images. As shown in Fig. 1, the resolution in the X-Y is limited by the tip-sample surface convolution depending on the geometry of the probe-scan-plane and sample-surface-plane. Commonly, the real tip-sample contact occurs on the tip side and not at the tip apex. Furthermore, the tip scans over the surface and the microscopy converts the obtained analog signal to a digital image with a certain number of points. Hence, the lateral resolution depends on the number of points for a given scan size as illustrated in Fig. 2. As illustrated in Fig. 3, for a small enough pixel size, the tip-sample convolution dominates the maximal obtainable resolution as the real contact point is not the tip apex. Furthermore, the tip-sample convolution in conjunction with the finite pixel size results in an interpolation of the surface, which is shallower than the real surface feature and is determined by the tip geometry.

We implemented a software module in the free SPM software Gwyddion that analysis the sample surface inclination in regard of such sample-tip convolution gradients. By assuming a certain tip radius and tip slope, we mark areas in the obtained topographic image, which are most likely exhibiting the wrong topographic information. The artifact analysis allows a better understanding of the instrument or acquisition parameters, i.e. tip radius needed to obtain artifact free images (e.g. use of super sharp tips), inclination of scan/surface planes, number of points needed for a image, or dynamic scanner range.


This research was financially supported by the Ministério da Ciência, Tecnologia e Inovação (MCTI) - Brazil.

Fig. 1: Illustration of tip apex /sample surface geometry convolution. The black line shows the surface profile obtained due to the convolution.

Fig. 2: AFM topography images of InGaAs surface in the same area scanned with (a) 32 X 32 pixels and (b) 256 X 256 pixels.

Fig. 3: Illustration of a profile interpolation resulting from the pixel size.

Type of presentation: Poster

IT-14-P-2536 Surface potential investigation of AlGaAs/GaAs heterostructures by Kelvin Force Microscopy

Pouch S.1, Chevalier N.1, Mariolle D.1, Triozon F.1, Niquet Y. M.2, Melin T.3, Borowik Ł.1, Delaye V.1
1CEA, LETI, MINATEC Campus, 17 rue des Martyrs, 38054 GRENOBLE Cedex 9, France, 2CEA, INAC, MINATEC Campus, 17 rue des Martyrs, 38054 GRENOBLE Cedex 9, France., 3Institut d’Electronique de Microélectronique et de Nanotechnologie, CNRS-UMR 8520, Avenue Poincaré, BP 60069,59652 Villeneuve d’Ascq Cedex, France.
sylvain.pouch@cea.fr

The Kelvin force microscopy (KFM) provides a spatially resolved measurement of the surface potential, which is related to the energetic band structure of a material. The goal of this work is to investigate the surface potential measured by KFM on AlGaAs/GaAs heterostructures.

For this study, we selected a certified reference sample BAM-L200 [1] composed of epitaxial layers of AlGaAs and GaAs, with a decreasing thickness (600 to 2 nm) and an uniform Si (n) doping (5.1017 cm-3). The surface potential measurement is performed with an Omicron XA VT AFM, under ultra-high vacuum (of 10-11 mbar). Two scanning modes are used: the amplitude modulation (AM-KFM), sensitive to the electrostatic force and the frequency modulation (FM-KFM), sensitive to its gradient. [2] Three kinds of tips have been used for this study: Platinum coated silicon tips (BudgetSensors), Au nanoparticles coated silicon tips (Next Tip) and super sharp silicon tips (Nanosensors).

We will present the measurements obtained with these different tips on the sample area containing the narrowest layers. The relevant result is the fact that the contrast decreases with diminution of the layer thickness. With Pt-coated Si tips, a maximum contrast of about 270 meV was observed, whereas for super sharp Si tips the maximum contrast equals 290 meV [Fig. 1]. This contrast vanishes when layer thickness becomes thinner than 5 nm for Pt-coated SI tips and 3 nm for super sharp Si tips. This loss of contrast can be explained primarily by the resolution limit of our instrument but also the band bending length scale at the AlGaAs/GaAs interface, related to the dopant concentration. The contribution of band bending between the layers to the measured potential is evaluated by a self-consistent simulation of the electrostatic potential, accounting for the free carriers distribution inside the sample and for the surface and interface dipoles. As shown in Fig. 2, the electric fields of the narrow layers recover each other, resulting in the partial or total loss of the intrinsic sample structure. Simple comparison of simulation with KFM surface potential measurements provides information that KFM measurements represent real values and are not influenced by KFM resolution limit.

All measurements were made on the CEA Grenoble nanocharacterization platform (PFNC).

[1] M. Senoner, T. Wirth, W. Unger, W. Österle, I. Kaiander, R. L. Sellin and D. Bimberg, BAM-L002 - a new type of certified reference material for length calibration and testing of lateral resolution in the nanometre range, Surface and Interface Analysis 36, 1423-1426 (2004)
[2] S. Sadewasser and T. Glatzel, Kelvin Probe Force Microscopy (2012)


Fig. 1: KFM images obtained on BAM-L200 with super sharp silicon tips, with respect to sample schema: (1) Topographic image; (2) Surface potential images.

Fig. 2: Theoretical simulation of the electric potential at 5 nm above the surface of the sample, and averaged section of previous KFM image.

Type of presentation: Poster

IT-14-P-2566 DESIGN, IMPLEMENTATION AND CHARACTERIZATION OF A 3D-PRINTED AFM HEAD WITH PIEZOTUBE AND ELECTROMAGNETIC ACTUATORS FOR BIOMOLECULAR APPLICATIONS

Sevim S.1, Özer S.2, Feng L.2, Crawford K.2, Karaca O.2, Torun H.2
1Department of Mechanical Engineering, Boğaziçi University, Bebek/Istanbul, Turkey, 2Department of Electrical and Electronics Engineering, Boğaziçi University, Bebek/Istanbul, Turkey
semih.sevim89@gmail.com

An atomic force microscope (AFM) has been developed for biomolecular force spectroscopy. Design, implementation and characterization of the AFM head are described here. The head is portable and was manufactured at low cost using stereolithography. A flexible software-based controller was implemented that can be adapted easily for different applications. The AFM head, made of a rigid polymer material (Rigid Opaque, Stratasys, Ltd., MN, USA) is shown in Fig. 1(a). The head houses a piezotube actuator, a laser diode, a quadrature photodetector and an AFM cantilever. The cantilever is mounted to the piezotube using a customized holder. The incident laser beam (fiber pigtailed, Oz Optics, Ottowa, Canada) is directed to the cantilever using a kinematic mount. Reflected laser light is directed to the quadrature detector (Pacific Silicon Sensor, Westlake Village, CA, USA) via a mirror, which has a one degree-of-freedom of rotation. The photodetector is placed on a translational microstage with two degrees-of-freedom. The space below the cantilever plane is empty so that the head can be integrated with an inverted microscope. In addition, an electromagnet was employed with the head for direct cantilever actuation and for other magnetic applications. A one-dimensional actuation coil is integrated to the head as shown in Fig. 1(b). The system allows cantilever actuation using the piezotube actuator and the electromagnet. Fig. 2(a) shows a sample force curve obtained using the piezo actuator. The drive signal in various waveforms was generated by the customized software-based controller to actuate a commonly used AFM cantilever (SNL-10D, Bruker Probes, Santa Barbara, CA) on a silicon sample at various frequencies, from 10 mHz to 1 kHz. In addition, Fig. 2(b) shows a typical current signal input to the electromagnet the corresponding photodetector signal. A commercially available MFM cantilever (MESP, Bruker Probes, Santa Barbara, CA) was actuated by a square wave in air at 1 kHz. The head was designed and optimized for force spectroscopy experiments. The force noise of the system using typical AFM cantilevers has been characterized as 6.8 pN within a bandwidth of 1 kHz. Finally, a biomolecular force spectroscopy experiment to probe interactions between FGF-2 and heparin was performed using the piezotube actuation. Fig. 3(a) shows a typical force curve, exhibiting an unbinding event with a force strength of 500 pN, whereas in Fig. 3(b) there is another force curve indicating no adhesion/rupture events.


Authors would like to acknowledge funding from the EC (ICT FET-Open) under the MANAQA Project.

Fig. 1: Fig. 1 (a) The photograph of assembled AFM Head, manufactured using stereolitography. (b) The photograph of AFM head employed with an electromagnet to actuate the MFM cantilevers.

Fig. 2: Fig. 2 (a) Typical force curves, taken with 10 Hz piezo actuation over a silicon wafer. (b) Electromagnetic actuation of a MFM cantilever in air at 1 kHz.

Fig. 3: Fig. 3 (a) A typical force curve exhibiting an unbinding event between FGF-2 and Heparin molecules. (b) A force curve which is indicating no adhesion/rupture event.

Type of presentation: Poster

IT-14-P-2717 Integration of SPM module inside FIB-SEM instrument

Rudolf M.1, Sedláček L.1, Jiruše J.1
1TESCAN Brno, s.r.o., Brno, Czech Republic
libor.sedlacek@tescan.cz

In the nanotechnology field, SPM (Scanning Probe Microscope) integrated into a SEM (Scanning Electron Microscope) offers a completely new opportunities [1]. Recently, TESCAN integrated a SPM with the lateral scan range up to 50 µm and Z scan range 8 µm [2]. Its compact construction is optimized for operating in a confined space of a FIB-SEM chamber without affecting the performance of both the electron or ion beams. Such dedicated SPM design allows the investigation of the same place on the sample by SPM, FIB (Focused Ion Beam) and high resolution SEM with a spot size down to 1 nm simultaneously without the need to perform an additional sample re-positioning.

We developed an intuitive software module to simplify the SPM navigation on the sample. It is possible to save the region of interest to a memory and recall it later on either by SPM, the SEM or both. Saved positions are shown along with their title in a live SEM window, see Fig. 1. The past experience has proved that the SPM module inside the FIB-SEM microscope is useful for several applications, such as process optimization of electron and ion beam lithography [3, 4] or TOF-SIMS (time-of-flight secondary ion mass spectroscopy) where the SPM can provide complementary information about the depth profile.

Fig. 2 shows an example how such a combination of the SEM, FIB and SPM was utilized. Hydrogenated Diamond-like Carbon (H:DLC) layer prepared by PECVD (Plasma Enhanced Chemical Vapor Deposition) method on a Si wafer was locally milled by FIB in order to uncover the Si-DLC interface, and the thickness of the DLC layer was measured using SPM. The damage caused by FIB milling is only local and it has no disturbing effects on measurements performed later on the same sample.

In another example we utilized in-situ cooperation of SEM/FIB/TOF-SIMS/SPM techniques together. FIB and TOF analyzer were used to create a concentration depth profile of elements contained in H:DLC layer deposited on an Si wafer, and the SPM navigated by the SEM provided an additional information about the true depth profile, see Fig. 3.

References:

[1] W Heichler, Microsc. Microanal. 19 (suppl. 2) (2013) p. 350.

[2] M&M 2011 trade show. See also [online]. [cit. 2014-03-03]. <http://www.specs.de/cms/upload/PDFs/SPECS_Prospekte/new_design/20130312_Curlew_brochure_final_web.pdf>

[3] J Jiruše et al, MC Proceedings Part 1 (2013) p. 154.

[4] J Jiruše et al, Proceedings 57th EIPBN (2013) p. 01-07.


The authors acknowledge funding from the European Union Seventh Framework Program [FP7/2007-2013] under grant agreement No. 280566, project UnivSEM.

Fig. 1: Positions on the sample surface are saved to the memory and shown in a live SEM window. Position in a memory can be recalled either by SMP, SEM, or both.

Fig. 2: DLC layer deposited on the Si wafer sputtered by the FIB. The DLC layer thickness is determined by the SPM module. (a) SEM image, (b) AFM image, (c) AFM profile.

Fig. 3: Concentration depth profile of carbon and silicon in H:DLC layer measured by TOF-SIMS with the depth information provided by SPM. From the SPM depth measurements and the sharp increase in the silicon signal intensity the thickness of a DLC layer can be obtained.

Type of presentation: Poster

IT-14-P-2845 In situ characterization of the growth mechanism of PEDOT films with electrochemical atomic force microscopy

Reggente M.1, Passeri D.1, Angeloni L.1, Rossi M.1,3, Tamburri E.2, Orlanducci S.2, Terranova M. L.2
1Department of Basic and Applied Sciences for Engineering, Sapienza University of Rome, Rome, Italy, 2Department of Chemical Science and Technology - MINASlab, University of Rome 'Tor Vergata', Rome, Italy, 3Center for Nanotechnology Applied to Engineering of Sapienza (CNIS), Sapienza University of Rome, Rome, Italy
melania.reggente@uniroma1.it

Conductive polymers (CP) belong to an attractive class of materials with plastic-like mechanical properties and electric conductivity typical of metals, which have awakened an increasing interest for several applications, e.g. sensing, electronic and energy. One of the most studied CP is the Poly(3,4-ethylenedioxythiophene) (PEDOT) due to its well-known properties and the advantage of being synthesized as thin-film directly on the substrates of interest. PEDOT thin films can be realized by electrochemical synthesis performed in an aqueous media containing a small quantity of the monomer 3,4-ethylenedioxythiophene (EDOT) and a suitable supporting electrolyte [1-2]. It is known that the process parameters influence both the structural and the electrical properties but further studies on the nucleation and growth mechanisms of the film formation are still required. Thus, it can be useful to monitor in real-time the synthesis process of PEDOT films in order to tune the process parameters and produce films with reproducible specific properties. Up to now, atomic force microscopy (AFM) based techniques have been employed to underlying the morphological features of the films at different steps of the deposition process and their related conductive properties but results of an in situ AFM investigation are not yet reported.
In this work, the growth mechanism of electrodeposited PEDOT films are investigated by using the electrochemical atomic force microscopy (EC-AFM) in order to determine the correlation between their morphological features and the electrochemical parameters of the process. In EC-AFM, a standard AFM apparatus is equipped with a three-electrode electrochemical cell whose working electrode is the sample surface where the electrodeposition takes place. Thus, a real time study of the electrochemical reactions occurring at the surface of the sample is achieved and the in situ surface morphology evolution is monitored by using an unbiased AFM probe. In particular, the electropolymerization of EDOT is observed performing a cyclic voltammetry and controlling the evolution of current flowing through the electrode surface, together with a standard AFM image. By varying the supporting electrolyte concentration, the voltammetry scan rate and the working electrode surface, the nucleation and growth mechanisms of the film are investigated and the results are compared with the already hypothesized growth model.
Overall, this work demonstrates the capability of the EC-AFM to deepen the growth mechanism of electrodeposited polymeric films with tunable and reproducible properties.
[1] E. Tamburri et al., Synthetic Metals, 159 (2009) 406–414.
[2] V. Castagnola et al., Synthetic Metals, 189 (2014) 7–16.


Type of presentation: Poster

IT-14-P-2847 Atomic force microscopy techniques for the detection of nanomaterials incorporated in biological systems

Reggente M.1, Passeri D.1, Angeloni L.1, Scaramuzzo F.1, De Angelis F.2, Barteri M.2, Rossi M.3
1Department of Basic and Applied Sciences for Engineering, Sapienza University of Rome, Rome, Italy, 2Department of Chemistry, University of Rome Sapienza,Rome, Italy, 3Center for Nanotechnology Applied to Engineering of Sapienza (CNIS), Sapienza University of Rome, Rome, Italy
melania.reggente@uniroma1.it

The capability of detecting nanomaterials (NMs) in biological samples represents one of the main challenges in bionanoscience, as it would allow the monitoring of cell-NMs interactions at the nanoscale, which is of primary importance in several fields of application, from drug delivery to nanotoxicology. To this aim, Atomic Force Microscopy (AFM) has been proposed as a versatile platform for the detection of NMs in biological matrices. By detecting the inhomogeneity of the mechanical, electrical or magnetic properties of the host-guest systems, the presence of NMs can be revealed [1].
In this work, different AFM-based techniques are employed to investigate the interactions between Magnetic Nanoparticles (MNPs) and cells. We show the possibility to reveal the presence of MNPs in biological system by detecting both the magnetic and the mechanical properties of the samples. First of all Magnetic Force Microscopy (MFM), a two-pass AFM-based technique which requires a tip coated with a magnetic film to obtain images reflecting the local magnetic properties of the samples, is employed for the imaging of MNPs internalized in cells. In addition to this, buried MNPs in soft biological matrices are visualized using three different AFM-based techniques in which the contrast reflects the non homogeneous mechanical properties of the host and guest systems. In particular, images of the local sample stiffness are obtained using the AFM force-volume imaging mode allowing the detection of force-distance curves. Moreover, the Contact Resonance Frequencies (CRFs) of the cantilever in contact with the sample surface, which are related to the local elastic modulus of the sample, are recorded employing the Atomic Force Acoustic Microscopy (AFAM). Following an offline procedure, the semi-quantitative CRFs maps are then converted into quantitative indentation modulus maps by assuming a proper model for the cantilever-tip-sample system. Furthermore, online maps of the local indentation modulus of the samples are recorded using Torsional Harmonics AFM (TH-AFM) which allows the evaluation of the local sample stiffness by acquiring force-distance curves in tapping mode. Finally, the mechanical properties evaluated with these three different techniques are compared and the influence of the penetration depth of each technique on the results is discussed and rationalized.
A careful comparison between the images obtained using all these techniques based both on the magnetic and the mechanical contrast allows to detect NMs incorporated in biological matrices and represents a clear indication of the AFM powerfulness in the field of nanobiotechnology.
[1] Atomic Force Microscopy in Cell Biology, B.P. Jena, J.K.H. Hörber (Eds), Academic Press, San Diego California USA (2002).


Type of presentation: Poster

IT-14-P-2848 Bacterial adhesion force measurements by microbial cell probe Atomic Force Microscopy

Angeloni L.1, 2, Passeri D.1, Reggente M.1, Pantanella F.3, Schippa S.3, Mantovani D.2, Rossi M.1
1Department of Basic and Applied Sciences for Engineering, University of Rome Sapienza, Via A. Scarpa 16, 00161 Rome, Italy, 2Lab. for Biomaterials and Bioengineering (CRC-I), Dept. Min-Met-Materials Eng. & University Hospital Research Center, Laval University, Quebec City, Canada , 3Department of Public Health and Infectious Diseases, University of Rome Sapienza, Piazzale A. Moro 5, 00185 Rome, Italy
livia.angeloni@uniroma1.it

The ability of bacteria to adhere to solid surfaces, proliferate and make a biofilm is the primary cause of food contamination, hospital infections and failures of long-term biomedical implants. Consequently, there is an increasing interest in developing surfaces with antibacterial properties in many areas such as food processing and health-related fields like medicine and dentistry.
In order to design antibacterial surfaces it is necessary to understand the physical and molecular interactions governing the bacterial adhesion to a surface, which is the first crucial step of biofilm formation. Several methods have been developed to evaluate these mechanisms and to identify the main influencing parameters. Static adhesion assays can provide experimental samples suitable for the qualitative or semi-quantitative measurements of bacterial adhesion. Fluid shear systems have been used to simulate the in vivo dynamic mechanical stress and to obtain global probabilistic measurements of the bacterial adhesion strength [1].
Nevertheless, the physical interactions involved in bacterial adhesion have not been understood in detail and the development of experimental procedures for further investigations is essential.
For a more focused investigation on the adhesion mechanisms, techniques comprising the manipulation of single bacteria are more appropriate.
Atomic Force Microscopy (AFM) can be used to obtain local information about the first physicochemical interaction phase of bacterial attachment to a surface, by the measurement of force-distance curves. The process can be studied by two different experimental approaches: i) by measuring the interaction forces between bacterial cells and a standard AFM tip [1] or ii) by measuring the interaction forces between bacteria on AFM tip and different surfaces [2].
In this work we develop an experimental procedure to obtain quantitative measurements of bacterial adhesion to different surfaces, by recording force-distance curves using probes coated with different bacterial species. Force-distance curves measurements are carried out, in air and in liquid, on substrates with different properties (chemical composition, hydrophobicity, charge)
The influence of the experimental conditions on the results is analyzed with the aim of assessing the most appropriate procedure.
Also, the results are discussed focusing on the influence of different physicochemical properties of bacteria and surfaces in the adhesion mechanism.
Overall, this work represents a preliminary study on the capability of AFM force distance curves measurements to investigate the physicochemical mechanism involved in the bacterial adhesion to abiotic surfaces.
[1] M. Katsikogianni et al., Eur Cell Mater 8 (2004) 37-57
[2] Y.J. Oh et al., Ultramicroscopy 109 (2009) 874–880


Type of presentation: Poster

IT-14-P-2849 Quantitative characterization of magnetic nanoparticles properties by Magnetic Force Microscopy

Angeloni L.1, 2, Passeri D.1, Reggente M.1, Marianecci C.3, Mantovani D.2, Rossi M.1
1Department of Basic and Applied Sciences for Engineering, University of Rome Sapienza, Via A. Scarpa 16, 00161 Rome, Italy, 2Lab. for Biomaterials and Bioengineering (CRC-I), Dept. Min-Met-Materials Eng. & University Hospital Research Center, Laval University, Quebec City, Canada, 3Department of Drug Chemistry and Technologies, University of Rome Sapienza, Piazzale A. Moro 5, 00185 Rome, Italy
livia.angeloni@uniroma1.it

The development of techniques for the characterization of magnetic nanomaterials has great interest by reason of the specific properties that occur in magnetic materials when their dimensions are reduced to the nanoscale. These properties, coupled with the nanometric size, make magnetic nanomaterials suitable for several biomedical applications. Magnetic nanoparticles (MNPs) can be used as carriers for drug delivery systems, mediators for magnetic hyperthermia treatments, contrast agents for Magnetic Resonance Imaging (MRI), markers for cell labeling [1].
The design of these techniques requires a detailed knowledge on the magnetic and structural properties of the adopted nanomaterials. For example the magnetic hyperthermia heating effect, the translational force exerted on drug delivery carriers, the drag force in cells magnetic separation systems are strongly dependent on the size and the magnetic properties of the nanoparticles, like the magnetic susceptibility χ, the saturation magnetization Ms, and the magnetic dipole m.
Standard techniques, like Superconducting Quantum Interference Devices (SQUID) or Vibrating Sample Magnetometer (VSM), allow the detection of global magnetic properties of nanoparticles populations. But the detection of magnetic properties of single particles is not possible and the evaluation of these properties in dependence of the particles size is not explicit.
In this work we develop an experimental procedure to obtain quantitative measurements of nanoparticles magnetic properties (χ, Ms, m) and to directly relate these characteristics with the particles size, by using Magnetic Force Microscopy (MFM). MFM is a particular non-contact scanning probe technique, based on the detection of the magnetostatic interaction between a magnetic AFM probe and a magnetic sample [2]. Thanks to its nanometric lateral resolution and its capability to detect weak magnetic fields, MFM is a powerful tool for the characterization of single nanoparticles dimensions and magnetic properties. However, MFM measurements are also affected by non magnetic tip-sample interactions. Consequently the quantitative magnetic characterization of nanomaterials requires the accurate analysis and interpretation of MFM data. In this respect, the study is also focused on the evaluation of the influence, on the MFM measurements, of non-magnetic tip-sample interactions, like electrostatic forces, with the aim of assessing an experimental procedure to detect only magnetic tip-sample interactions.
Overall, this work represents a preliminary study on the applicability of MFM technique in the quantitative measurement of properties of magnetic nanomaterials.
[1] Q. A. Pankhurst et al., J. Phys. D: Appl. Phys. 36 (2003) R167–R181
[2] P. Grutter, Ultramicroscopy, 47 (1992) 393-399


Type of presentation: Poster

IT-14-P-3249 Water meniscus investigated at nanoscale contacts with a heated atomic force microscope (AFM) cantilever probe

Assy A.1, Lefèvre S.1, Chapuis P. O.1, Gomes S.1
1Université Lyon 1, CETHIL, UMR5008, F-69621 Villeurbanne cedex, France
ali.assy@insa-lyon.fr

While working under ambient conditions, Scanning Probe Microscopy (SPM) techniques face up to a liquid meniscus when the probe gets into contact with the sample. The meniscus is formed due to the capillary condensation of the ambient environment. This meniscus could be a barrier and prevents exploiting the performed measurements or an advantage for some applications like the “Dip-Pen Nanolithography”. In the case of some applications where BioMEMS or NEMS/MEMS are involved, the meniscus problem is referred as to the stiction. In our case, the stiction is the large lateral force required to initiate relative motion between the probe and the sample. In order to find out a solution to this problem, we present an investigation of the volume and the radii of the water meniscus at different temperatures of the probe. The probe is mounted on an atomic force microscope (AFM) for its 3D positioning and displacement and for controlling the force between the probe and the sample. A resistive element is located at the tip apex and serves to heat the probe depending on the electrical current. The probe temperature is verified through a Wheatstone bridge and is maintained constant during the approach of the probe to the sample. The variations of the capillary forces are measured at different probe temperatures on different hydrophilic and hydrophobic samples. The measurements as a function of the probe temperature show a progressive evaporation of the meniscus. Moreover, and simultaneously to these variations, the heat conductance to the sample is measured. A correspondence between the thermal signal and the capillary forces is evidenced as shown in Figure 1. Based on theoretical models found in the literature, the meniscus interaction radius is evaluated from the capillary forces. Afterwards, the heat conductance at different probe temperature levels is linked with the evolution of the capillary forces, e.g. the meniscus volume. The experimental results obtained with different probes are compared and in accordance with literature values. The effect of roughness on the capillary forces is shown for different samples. For each used probe, we introduce a model that takes into account all the heat transfer mechanisms that operate simultaneously between the probe and sample. The transposition of these results could be interesting for many related applications such as BioMEMS and NEMS/MEMS.


Fig. 1: An example of the correspondence between the pull-off forces and the heat flux ratio as a function of the probe mean temperature (Tm). The measurements shown here are between a Pt/Rd microprobe and a Ge sample. (IT) and (DT) stand up for increasing temperature and decreasing temperature respectively.

Type of presentation: Poster

IT-14-P-3424 Surface Potential Distribution on Quartz Crystal Surfaces by AFM Silica-Probing

Yelken G. O.1, Polat M.1
1Izmir Institute of Technology, Department of Chemical Engineering, Urla, Izmir, Turkey
gulnihalyelken@iyte.edu.tr

Interaction forces between colloidal particles play an important role in numerous physicochemical systems in mineral, ceramic, and environmental sciences since they determine stability, rheology, and forming characteristics. Control and manipulation of these properties depend on detailed analysis of the interactions among the particles. Interparticle interactions can be divided into two main categories; van der Waals(vdW) and Electrical Double Layer [1,2].
Proper use of these theories and their comparison with the experimentally measured force values require knowledge of such material properties as Hamaker constants and charge on the interacting surfaces. The surface potential at the point of measurement could then be determined from the electrostatic component. Multiple AFM force measurements on carefully selected locations on the surface could be used to generate a surface charge/potential map of the surface using appropriate theories [3,4,5].
In this study we used a powerful surface analysis tool, AFM, to determine the surface charge or surface potential on quartz single crystal surfaces in aqueous solutions. This use of AFM is new and novel and requires insightful use of theory and experiment. Using AFM to map the charge distribution on surfaces in solution is different than the EFM measurements in air since measuring surface potential in air or in vacuum is a straightforward process which has been used for years using different devices. The methodology, we used is basically depends on a point by point comparison of measured interaction force between a surface and the AFM tip of known characteristics with the theoretical force predicted for the same system. The results were confirmed with separate electrokinetic measurements of all surfaces.

References
1. [1] Derjaguin, B.V., L. Landau, Physicochim, URSS, No:14, 633,1941.
2. Verwey, E.J.W., J.T.G. Overbeek, Theory and Stability of Lyophobic Colloids, Elsevier, Amsterdam,1948..
3. Sader, J.E., Chon, J.W.M., Mulvaney, Calibration of rectangular atomic force microscope cantilevers, Rev. Sci. Instrum., No: 70, 3967-3969, 1999.
4. Polat, M., H. Polat, Analytical solution of Poisson–Boltzmann equation for interacting plates of arbitrary potentials and same sign, J.of Colloid and Interface Science, 341,1, 178-185, 2010.
5. Yelken,G. O., Polat, M., Determination of electrostatic potential distribution by atomic force microscopy (AFM) on model silica and alumina surfaces in aqueous electrolyte solutions, Applied Surface Science, 2014. http://dx.doi.org/10.1016/j.apsusc.2014.02.022


The support from The Scientific and Technological Research Council of Turkey (TUBITAK) under the project grant TUBITAK 109T695 is acknowledged.

Fig. 1: Surface potential distributions on a 5 μm × 5 μm portion of the quartz (0001) surface at  pH=2 values in 10−3 M KCl solution.

Type of presentation: Poster

IT-14-P-5952 Photothermal Excitation for Reliable and Quantitative AFM

Johann F.1, Labuda A.1, Walters D.1, Bocek D.1, Rutgers M.1, Cleveland J.1, Proksch R.1
1Asylum Research, an Oxford Instruments Company
florian.johann@oxinst.com

Since the advent of atomic force microscopy, cantilevers have predominantly been driven by piezos for AC imaging and data acquisition. The ease of use of the piezo excitation method is responsible for its ubiquity. However, the well-known “forest of peaks”, which is clearly observed while tuning a cantilever in liquids, renders AC imaging in liquids problematically because the peaks move around with time (see Figure 1). Effectively, these shifting peaks result in a setpoint that changes with time causing stability problems while AFM imaging. Furthermore, the same “forest of peaks” prevents the quantitative interpretation of forces in liquids[1], air[2], and vacuum environments[3], even if the cantilever tune looks clean. Dissipation studies in all these environments have especially suffered due to piezo excitation of the cantilever.

Photothermal excitation is an alternative method for exciting a cantilever by heating/cooling the base of the cantilever to drive the cantilever. Photothermal excitation results in repeatable, accurate and time-stable cantilever tunes, as seen in the Figure. Therefore, the setpoint remains truly constant while imaging, preventing tip crashes, or unwanted tip retractions. True atomic resolution images of calcite in water, shown in Figure 2, were made for hours with no user intervention, testifying to the stability of photothermal excitation. Unlike other specialized drive methods, photothermal excitation is compatible with almost any cantilever and with all AFM techniques. The introduction of a blue laser into the AFM also enables several other functionalities, such as tuning the temperature of the cantilever. Furthermore, because the photothermal tune represents the true cantilever transfer function, existing AFM theories can be applied to accurately recover conservative and dissipative forces between the tip and the sample. This is especially important for force spectroscopy, dissipation studies, as well as the frequency modulation AFM techniques.

Our recent developments in perfecting photothermal excitation and its benefits to the AFM community will be shown.

[1] A. Labuda, K. Kobayashi, et al. AIP Advances 1, 022136 (2011)
[2] R. Proksch and S. V Kalinin, Nanotechnology 21, 455705 (2010)
[3] A. Labuda, Y. Miyahara, et al. Phys. Rev. B 84, 125433 (2011)


Fig. 1: Since the amplitude and phase do not drift with time, blueDrive delivers stable imaging. Piezo drive, on the other hand, has a time varying amplitude and phase which requires constant intervention by the user to maintain stable imaging conditions.

Fig. 2: A freshly cleaved crystal face of calcite was imaged in ultrapure water during three hours. No user intervention was necessary throughout the experiment, because the drive amplitude was remained stable.

IT-15. X-ray, neutron and other microscopies

Type of presentation: Invited

IT-15-IN-1716 Chemically selective spectromicroscopy by soft x-ray scanning transmission x-ray microscopy

Hitchcock A. P.1
1Dept. of Chemistry & Chemical Biology McMaster University, Hamilton, Canada
aph@mcmaster.ca

Soft X-ray scanning transmission X-ray microscopy (STXM) uses natural near edge X-ray absorption spectral contrast for chemical speciation and quantitative chemical & orientation mapping (geometric & magnetic) in 2d & 3d with <20 nm spatial resolution. Recently STXM capabilities have been expanded to include electron detection for surface studies, X-ray fluorescence for enhanced sensitivity, and ptychography. STXM is ideal for wet samples since soft X-rays readily penetrate a few microns of water. I will outline instrumentation, data analysis, and capabilities of soft X-ray STXM. Examples will include:

Biomagnetism. STXM with circularly polarized light [CLS 10ID1 or ALS 11.0.2] measure magnetism by X-ray magnetic circular polarization (XMCD). We use this to study magnetotactic bacteria [1] which biomineralize intra-cellular chains of ~50 nm magnetite single crystals. In most cases all magnetic moments in a chain point in the same sense. Recently we found cases where there is internal reversal -the magnetic field of one part points opposite to other parts of the chain (Fig. 1). The gap region exhibits an Fe L3 spectrum similar to that of magnetite but without XMCD [2]. These are situations where either magnetite bio-mineralization has failed or the chain is in the act of growing. Our studies provide insights into biomineralization. Use of ptychography to measure XMCD with improved spatial resolution (<10 nm) will be described.

Automotive hydrogen fuel cells. Polymer electrolyte membrane fuel cells (PEM-FC) are being developed for near-future mass production automotive applications. The performance, efficiency and lifetime of PEM-FC depend on composition and nanostructure of electrodes. Optimization is critical for the cathode where the rate limiting oxygen reduction reaction takes place. STXM is a powerful tool to study a wide range of issues in PEM-FC optimization including mapping ionomer in cathodes [3,4]. Most studies to date have been carried out on dry, microtomed samples at ambient temperature (25 C, 0 % RH) which are very different from typical operating conditions of PEM-FC (70 C, 80 % RH). The nanostructure change with temperature and hydration. Instrumentation and methods to examine PEM-FC under more realistic conditions are needed. We have developed an environmental cell for in situ STXM measurements under controlled relative humidity (0-100%) and temperature (-30 - 80 C) (Fig. 2). We study water saturation in cathode and membrane [5] and changes on freezing.

1. K.P. Lam, et al. Chem. Geology 270 (2010) 1101; S. Kalirai et al. ibid 300 (2012) 14.
2. S. Kalirai et al. PLOS One 8 (2013) e53368.
3. V. Berejnov et al PCCP 14 (2012) 4835.
4. V. Berejnov, et al. ECS Trans., 50 (2012) 361.
5. V. Berejnov et al., ECS Trans. 41 (2011) 395.


Research supported by NSERC, CFI, OIT, Canada Research Chair funding and AFCC. CLS is supported by NSERC,CIHR, NRC and U. Saskatchewan. ALS (LBNL) is supported by BES, DoE.

Fig. 1: Internal magnetic reversal in a magnetotactic bacterium (MTB). (a) Fe L3 STXM-XMCD of an MV-1 MTB measured with circular polarization parallel (green), anti-parallel (red) to magnetic vector of chain. (b) TEM of cell with interrupted chain. (c) XMCD spectra of 3 sub-chains. (d) STXM image at 710 eV. (e) color coded XMCD signal.

Fig. 2: (a) cartoon of the in situ STXM environmental cell. (b) photo in ALS 5322 STXM. (c) O 1s spectra of 3 phases of water. (d) color coded composite (cathode, PFSA, liquid water) from O 1s stack of a PEM-FC membrane electrode assembly at 85% RH. (e) color coded composite (water vapor, PFSA, liquid water) from the same stack.

Type of presentation: Invited

IT-15-IN-3325 Imaging live cells by X-ray laser diffraction

Nishino Y.1, Kimura T.1, Joti Y.2, Bessho Y.3
1Research Institute for Electronic Science, Hokkaido University, Sapporo, Japan, 2Japan Synchrotron Radiation Research Institute/SPring-8, Hyogo, Japan, 3Academia Sinica, Taipei, Taiwan
yoshinori.nishino@es.hokudai.ac.jp

Coherent imaging is a growing field in optical science. It requires no lens for image formation, but instead numerically reconstructs object images from the coherent diffraction data. It is, therefore, advantageous for x-rays, for which it is difficult to fabricate lenses with a high numerical-aperture. Coherent imaging has been demonstrated to be a powerful tool to visualize cells and organelles using synchrotron radiation [1,2]. Recently emerging X-ray free-electron lasers (XFELs) further extends the ability of coherent imaging to achieve spatial resolution beyond the conventional radiation-damage limitation. Because the pulse duration of XFELs is in the femtosecond range, X-ray interaction with the sample occurs before radiation damage becomes obvious. XFELs also allow us to image samples in solution in close-to-natural conditions [3]. We are developing a method which we refer to as pulsed coherent x-ray solution scattering (PCXSS). We performed PCXSS experiments using a Japanese XFEL facility SACLA. We will present some early results of our PCXSS experiments performed for inorganic and biological samples [4].

References:
[1] “Imaging whole Escherichia coli bacteria by using single-particle x-ray diffraction”: J. Miao, K. O. Hodgson, T. Ishikawa, C. A. Larabell, M. A. Le Gros & Y. Nishino, Proc. Natl. Acad. Sci. U.S.A. 100, 110–112 (2003).
[2] “Three-Dimensional Visualization of a Human Chromosome Using Coherent X-Ray Diffraction”: Y. Nishino, Y. Takahashi, N. Imamoto, T. Ishikawa & K. Maeshima, Phys. Rev. Lett. 102, 018101 (2009).
[3] “Advances in X-ray scattering: from solution SAXS to achievements with coherent beams”: J. Pérez & Y. Nishino, Curr. Opin. Struct. Biol. 22, 670–678 (2012).
[4] “Imaging live cell in micro-liquid enclosure by X-ray laser diffraction”: T. Kimura, Y. Joti, A. Shibuya, C. Song, S. Kim, K. Tono, M. Yabashi, M. Tamakoshi, T. Moriya, T. Oshima, T. Ishikawa, Y. Bessho, &Y. Nishino, Nature Commum. 5, 3052 (2014).


This research was partially supported by the X-ray Free Electron Laser Priority Strategy Program from MEXT; CREST from JST; KAKENHI Grant Numbers 23651126, 22310075, 22540424, and 23860001 from JSPS; and the Cooperative Research Program of ‘Network Joint Research Center for Materials and Devices’. We thank the operation and engineering staff of SACLA for helping perform the PCXSS experiments.

Fig. 1: Schematic of pulsed coherent X-ray solution scattering (PCXSS)

Fig. 2: Coherent x-ray diffraction pattern from a living Microbacterium lacticum cell exposed to a single XFEL pulse

Fig. 3: Reconstructed image of a living Microbacterium lacticum cell 

Type of presentation: Oral

IT-15-O-2379 Lessons learnt from four years of experience with diffractive imaging at X-ray Free Electron Laser sources

Strueder L.1, Hartmann R.1, Holl P.1, Huth M.1, Schmidt J.1, Soltau H.2
1PNSensor, Munich, Germany, 2PNDetector, Munich, Germany
lothar.strueder@pnsensor.de

Fourth generation accelerator-based light sources, such as VUV and X-ray Free Electron Lasers (FEL), deliver ultra-brilliant (~1012 -1013 photons per bunch) coherent radiation in femtosecond (~10 fs to 100 fs) pulses and, thus, require novel focal plane instrumentation in order to fully exploit their unique capabilities. As an additional challenge for detection devices, existing FELs (FLASH, Hamburg, LCLS, Menlo Park; SACLA, Hyogo) cover a broad range of photon energies from the EUV to the X-ray regime with significantly different bandwidths, intensities and pulse structures.

In order to meet these challenges, a novel, large area, broadband (50 eV to 25 keV), high-dynamic-range, intensity and spectroscopic imaging X-ray detector based on the pnCCDs has been established [1]. The sensor covers an area of 60 cm2 with 1024 x 1024 pixels and 10.000 x 10.000 spatial resolution points, including a hole in the center for the non-scattered X-rays. They have been operated up to 120 Hz in a full frame high resolution mode. The pnCCD detectors have been used in experiments from 30 eV (FLASH) up to 9.5 keV (LCLS, SACLA). The sensitive thickness of the fully depleted, fully sensitive CCDs is 450 µm. As the detectors are back-illuminated, an ultra-thin radiation entrance window has been developed to achieve clean energy spectra and high quantum efficiency for the lowest to the highest energies. Some of the detectors are equipped with integrated light blocking filters to avoid signal deterioration through visible light (see Fig. 1).

Different classes of experiments have been performed, each going towards the physical limits of measurement precision of the detectors: highest energy resolution (see Fig. 2) (atomic physics), the highest dynamic range (nano-crystallography), imaging of biological samples and X-ray scattering experiments (Bond orientational order of liquid and supercooled water) requiring a position resolution well below 10 µm. For all of the above experiments optimizations have been realized to fulfill the experimental requirements. The deep subpixel resolution and the controlled extraction mode of the detectors have already been demonstrated at the light sources [2]. Fig.3 shows the improvement of the charge handling capacity from 3x105 electrons per pixel to more than 1.5x106 measured at LCLS. The better understanding of the detector physics and data analysis leads to an optimization of operation modes for specific experiments, enabling for the development of new and more precise measurement methods. Detectors of this type will be used in X-ray microscopy this summer. Measurements from this application will be shown equally.

[1] L. Strüder et al., Nucl. Instr. and Meth.A 614 (2010),483-496
[2] S. Send et al., Nucl. Instr. and Meth.A711(2013)132-142


Fig. 1: Image of the pnCCD detectors on a 150 mm Si-wafer. The central chip has an area of 60 cm2, a pixel size of 75x75 µm2 and a format of 1024x1024. The sensitive thickness is 450 µm. It has a center hole for the passage of the non-scattered X-rays.

Fig. 2: X-ray emission of highly excited Xe –atoms (35+) at an LCLS experiment. The excitation energy was 1.5 keV, the energy resolution is approx. 100 eV (FWHM) at 1.5 keV, integrated over an area of 60 cm2 with a frame rate of 120 Hz of a 1024 x 1024 format of the pnCCD spectroscopic X-ray imaging array.

Fig. 3: Left: charges over flooding neighboring pixels. The max. charge handling capacity (CHC) in approx. 3x105electrons per pixel for this standard setting. Right: The same scattering process with a CHC of approx. 1.5x106 electrons per pixel due to different operating conditions of the same detector.

Type of presentation: Oral

IT-15-O-2550 Diffraction Contrast Tomography as an Additional Characterization Modality on a 3-D Laboratory X-ray Microscope

Feser M.1, Merkle A.1, Holzner C.1, Fahey K.1, Lauridsen E.2, Reischig P.2, Poulsen H. F.2
1Carl Zeiss X-ray Microscopy Inc., Pleasanton CA 94566, USA, 2Xnovo Technology ApS, 4600 Køge, Denmark
michael@feser.org

We introduce a novel method to add grain position, orientation and size information to absorption 3-D x-ray microscope imaging for poly-crystalline samples. This imaging modality will be available on a commercial x-ray microscope and will open the way for routine, non-destructive studies of time-evolution of grain structure to complement destructive EBSD end-point characterization. Grain sizes down to 40 micrometers can be studied using this non-destructive image modality.

Crystallographic imaging (i.e. imaging of crystallites/grains in polycrystalline materials) are primarily known from electron microscopy, and particularly the introduction of the electron back-scattering diffraction (EBSD) technique in the early 1990’s, has made it a routine tool for research and/or development related to metallurgy, functional ceramics, semi-conductors, geology etc. The ability to image the grain structure in such materials is instrumental for understanding and optimization of material properties and processing. However, the destructive nature of 3D EBSD prevents the technique from directly evaluating the microstructure (and grain-orientation) evolution when subject to either mechanical, thermal or other environmental conditions. Non-destructive x-ray diffraction imaging methods allow for such ‘4D’ time dependent studies, and to date have been primarily the domain of a limited number of synchrotron facilities.

Here, we present a novel method to acquire, reconstruct and analyze grain orientation and related information from polycrystalline samples on a commercial laboratory x-ray microscope (ZEISS Xradia 520 Versa) that utilizes a synchrotron-style detection system. Known as laboratory diffraction contrast tomography (DCT), this technique may be efficiently coupled to in situ environments within the microscope or subject to an extended time evolution experiment (across days, weeks, months), which remains a unique strength of laboratory (non-synchrotron) experiments. Following an evolution experiment, the sample may be sent to the electron microscope or focused ion beam (FIB-SEM) for destructive but complementary investigation of the same volume of interest.

We will show a selection of results of laboratory DCT, discuss the boundary conditions of such a method, and point to the future to discuss ways in which this can be correlatively coupled to related techniques for a better understanding of a materials structure evolution in 3D at multiple length scales.


Type of presentation: Oral

IT-15-O-2821 Three-Dimensional architecture of Hepatitis C virus replication factory studied by soft X-ray cryo-tomography

Perez-Berna A. J.1, Rodriguez M. J.2, Friesland M.2, Sorrentino A.1, Chichon F. J.2, Carrascosa J. L.2, Gastaminza P.2, Pereiro E.1
1ALBA Synchrotron Light Source, MISTRAL Beamline – Experiments Division, 08290 Cerdanyola del Vallès, Barcelona, Spain, 2Centro Nacional de Biotecnología-Consejo Superior de Investigaciones Científicas (CNB-CSIC), Campus Cantoblanco, 28049 Madrid, Spain
ajperezberna@gmail.com

Hepatitis C virus (HCV) is a major cause of chronic liver disease, with an estimated 170 million people infected worldwide. Low yields, poor stability and inefficient infection systems have severely limited the analysis of the HCV life cycle and the development of effective antivirals and vaccines. HCV is a positive strand RNA that replicates its genome in intracellular membranes forming a complex membranous web. Nevertheless, the three-dimensional structure of this membranous web in whole infected cells is still unknown.

In this study we have performed full-field cryo soft X-ray tomography (cryo-SXT) in the water window photon energy range (Schneider et al. 2010; Chichon et al. 2012) to investigate in whole, unstained cells, the morphology of the membranous rearrangements induced by the HCV replicon in conditions close to the living physiological state. We have obtained the first complete cartography of the dramatic cellular modification caused by the stable subgenomic HCV replicon transfected in cell culture (Kato et al, 2003). Moreover, in order to identify the viral proteins allocation in the different subcellular compartments, we have correlated the three-dimensional structure obtained with X-rays with electron microscopy immunelabeling and confocal immunofluorescence. The morphology of the membranous HCV factory web is a cytoplasmic accumulation of large and small heterogeneous vesicles, mitochondria and lipid droplets.

The understanding of the membranous replication factory of HCV provides a powerful tool for the analysis of host-virus interactions that should facilitate the discovery of antiviral drugs and vaccines for this important human pathogen.

Schneider G, Guttmann P, Heim S, Rehbein S, Mueller F, Nagashima K, Heymann JB, Müller WG, McNally JG. Three-dimensional cellular ultrastructure resolved by X-ray microscopy. Nature Methods 2012, 7: 985-987

Chichón FJ, Rodriguez MJ, Pereiro E, Chiappi M, Perdiguero B, Guttmann P, Werner S, Rehbein S, Schneider G, Esteban M, Carrascosa JL. Cryo X-ray nano-tomography of vaccinia virus infected cells. J. Struct. Biol. 177, 202-211 (2012).

Kato T, Date T, Miyamoto M, Furusaka A, Tokushige K, Mizokami M, Wakita T. Efficient replication of the genotype 2a hepatitis C virus subgenomic replicon. Gastroenterology. 2003 Dec;125(6):1808-17.


These experiments were performed at MISTRAL beamline at ALBA Synchrotron Light Facility with the collaboration of ALBA staff.

Type of presentation: Oral

IT-15-O-3352 Laboratory Full-Field Transmission X-ray Microscopy and Applications

Dehlinger A.1,2,3, Seim C.1,2, Legall H.1,2,3, Stiel H.1,3, Rancan F.4, Meinke M.4, Rehbein S.5, Wiesemann U.6, Kanngießer B.1,2
1Berlin Laboratory for innovative X-ray technologies (BLiX), Berlin, Germany, 2Technische Universität Berlin, Institut für Optik und Atomare Physik, Berlin, Germany, 3Max-Born-Institut, Berlin, Germany, 4Charité Berlin, Clinical Research Center for Hair and Skin Science, Berlin, Germany, 5Helmholtz-Zentrum für Materialien und Energie, Berlin, Germany, 6Bruker ASC GmbH, Cologne, Germany
aurelie.dehlinger@mbi-berlin.de

X-Ray microscopy in the water window allows imaging with resolutions in the nanometer regime as well as a high natural contrast between carbon and oxygen. Hence, it is possible to examine aqueous biological samples with up to 10 µm thickness in their natural state. Apart from cryo fixation of the specimen, which is usually required in order to avoid radiation damage, extensive sample preparation is not necessary. The use of highly brilliant laboratory X-ray sources has allowed the transfer of this technology, previously limited to synchrotron facilities, into the laboratory. This transfer inures to the benefit of a broader scientific community for applications in various fields such as medicine, biology and environmental sciences.

We introduce the plasma driven laboratory full-field transmission X-ray microscope (LTXM) located at the Berlin Laboratory for innovative X-ray technologies [1]. The half-pitch resolution of ∆x = (31 ± 3) nm is comparable to resolutions achieved at synchrotron facilities. The used wavelength at 2.478 nm is close to the absorption edge of oxygen and thus offers the best contrast within the water window. The large penetration depth and the short exposure times of less than one minute reached by the microscope, make soft X-ray cryo tomography feasible. An overview of first applications, like measurements on cryo-frozen yeast cells and human skin slices, will be given.

References

[1] H. Legall, G. Blobel, H. Stiel, C.Seim et al., “Compact x-ray microscope for the water window based on a high brightness laser plasma source,” Opt.Express, 20(16), 18369-18369 (2012).


This project was funded by the BMBF (#13N8913) and the BMVBS (WTW #03WWBE106).

Type of presentation: Poster

IT-15-P-1449 Single flash imaging of live hydrated biological cells by a contact soft x-ray microscope coupled with an intense laser-plasma soft x-ray source

Kado M.1, Kishimoto M.1, Tamotsu S.2, Yasuda K.2, Aoyama M.2, Shinohara K.1
1Quantum Beam Science Directorate, Japan Atomic Energy Agency, 2Division of Natural Science, Nara Women's University
kado.masataka@jaea.go.jp

We have developed a contact soft x-ray microscope combined with an intense laser-plasma soft x-ray source to achieve flash imaging of live hydrated biological cells. Laser-plasma soft x-ray source produced by a high power pulsed laser is extremely bright and very suitable for biological x-ray microscopy to capture an image of living specimens for which require a single flash exposure to avoid imaging any damages on the specimens. We also have invented to use a fluorescent microscope to identify the cellular organelles in the images obtained with the soft x-ray microscope. The biological cells were cultivated directly onto the PMMA photo resists and observed with the soft x-ray microscope and the fluorescent microscope at the same time. The obtained soft x-ray images and fluorescence images of the cells were directly compared and each cellular organelle such as mitochondria, actin filaments, and chromosomes in the soft x-ray images were clearly identified. Since the soft x-ray microscope has higher spatial resolution than that of the fluorescent microscope, fine structures of the cellular organelles in the hydrated biological cells were observed.
Shown in figure 1 are the soft x-ray image (a) and the fluorescence image (b) of the live biological cells. Appearing blue in the fluorescence image were chromatin, red were mitochondria, and green were actin filaments. The both images were clearly identical and each cellular organelle in the soft x-ray image could be identified directly comparing with the fluorescence image.
Shown in figure 2 are the soft x-ray images of one of the cells (a) shown in Fig.1 and enlarged images of surrounding area of the nucleus (b) and mitochondria (c) in the same cell. The cellular organelles such as chromatin and mitochondria in the images were identified comparing directly with the fluorescence image. All of the bright spots surrounding the nucleus in Fig. 2(a) were recognized to be mitochondria. Shown in Fig. 2(c) is the soft x-ray image of a single mitochondrion picked up from the Fig. 2(b) and detailed structure of the mitochondrion was obtained.


This research was partially supported by the Ministry of Education, Science, Sports and Culture, Grant-in-Aid for Scientific Research (C), 25390134, 2014.

Fig. 1: Soft x-ray image (a) and fluorescence image (b) of live hydrated biological cells. Appearing blue in the fluorescence image were nuclei, red were mitochondria and green were cytoskeletons.

Fig. 2: Soft x-ray images of one of the cells (a) shown in Fig. 1 and enlarged images of surrounding area of the nucleus (b) and mitochondria (c) in the same cell. The cellular organelles such as chromatin and mitochondria in the images were identified comparing directly with the fluorescence image.

Type of presentation: Poster

IT-15-P-1509 At-wavelength observation of phase defect embedded in extreme ultraviolet lithography mask

Amano T.1, Terasawa T.1, Watanabe H.1, Toyoda M.2, Harada T.3, Watanabe T.3, Kinoshita H.3
1EUVL Infrastructure Development Center, Inc., Ibaraki, Japan, 2Tohoku University, Miyagi, Japan, 3University of Hyogo, Hyogo, Japan
tsuyoshi.amano@eidec.co.jp

Extreme ultraviolet (EUV) lithography is considered to be the most promising next-generation lithography after the point where 193-nm immersion lithography would cease to deliver smaller features. However, the path to establish the EUV lithography is not without technical difficulties. Issues with insufficient light-source power, defect-free mask fabrication, and resist material development are to be resolved. Regarding the types of mask defects, the nature of the pattern defects in the EUV mask is mostly same as in the case of optical masks except for those defects that are classified as reflective multilayer defects, such as bump or pit phase defects that propagate through the multilayer during its deposition on the substrate surface and it is hard to repair. Therefore, to reduce the effect of the phase defect on wafer printing image, two methods are suggested. One method is to cover the phase defects beneath the absorber pattern by shifting the location of device pattern during mask patterning. The other is to eliminate the influence of the phase error by removing the absorber away from the close proximity of the phase defects after fabricating the device pattern. To make these methods success, it would be necessary to be able to pinpoint the location of the phase defects and the affected areas.
In this presentation, influence of the phase defect structures on EUV microscope images were examined to predict the inclination angle dependency of the phase defect impact on wafers since the phase defect does not always propagate in a vertical direction from the substrate surface through the multilayer. Figures 1(a) and 1(b) show photograph of the EUV microscope and illustration of the imaging optics, developed by Tohoku Univ. that was utilized in this study. The EUV light was sourced from a beam line BL3 of the New SUBARU synchrotron facility at the Univ. of Hyogo. A programmed phase defect EUV mask was prepared. Figure 2(a) shows the cross-sectional transmission electron microscope (TEM) images of the vertical and inclined grown phase defects. The calculated inclination angles of the phase defects were 0 and 4 degrees. Figure 2(b) represents the scanning probe microscope (SPM) images of the phase defects with half-pitch 88 nm lines-and-spaces (L/S). The L/S with the phase defects were observed using the EUV microscope. Figure 2(c) show the EUV microscope images and their intensity profiles. The impacts of the inclination angles on EUV microscope images were significant even though the positions of the phase defect relative to the absorber line, as measured by scanning prove microscope, were same. As a result, the EUV microscope could identify the positional shift of the effective defect position caused by the inclined propagation through the multilayer.


This work was supported by New Energy and Industrial Technology Development Organization (NEDO).

Fig. 1: (a) Photograph of the EUV microscope. (b) Schematic model of the EUV microscope optics.

Fig. 2: (a) Cross-sectional TEM images of the vertical- and inclined-grown phase defects. (b) SPM images of the phase defects in half-pitch 88 nm L/S. (c) EUV microscope images and intensity profiles.

Type of presentation: Poster

IT-15-P-1561 Online Tools for Microscopy and Microanalysis Facilities.

Apperley M. H.1, Whiting J.1, Cribb B.2, Frost C.2, Ceguerra A.3, Liddicoat P.3
1Australian Microscopy and Microanalysis Research Facility, The University of Sydney, NSW, 2006, Australia, 2Center for Microscopy and Microanalysis, The University of Queensland, QLD, 4072, Australia, 3Australian Center for Microscopy and Microanalysis and the School of Aerospace, Mechanical and Mechatronic Engineering, The University of Sydney, NSW, 2006, Australia
miles.apperley@sydney.edu.au

Australian Microscopy and Microanalysis Research Facility (AMMRF) is a national grid of equipment, instrumentation and expertise in microscopy and microanalysis that provides nanostructural characterisation capability and services, from widely used optical, electron, X-ray and ion-beam techniques to world-leading flagship platforms. One of the benefits of the network of core facilities is the ability to collaborate to develop online tools that are then accessible by all the laboratories in the network. These tools assist researchers to identify the techniques they need to use, facilitate training and enable data analysis & management.
The Technique Finder (TF) is a web application that enables prospective facility users to identify the techniques most suited to their research, based on a researcher-centric approach and terminology as opposed to instrument-oriented jargon.
MyScope: Training for Advanced Research, is an online suite of education tools for teaching and learning in the area of microscopy and microanalysis. The modules in MyScope contain a number of components including: an interactive questionnaire to allow the user to assess their knowledge, guide choices and tailor the learning environment for flexible learning; self guided tutorials with videos, animations and glossary to prepare students with knowledge and specialist language; virtual instrument platforms to practice use of instrumentation; and online competency testing to demonstrate readiness for hands-on experience.
A Data Management System (DMS) addresses the needs of an increasing number of AMMRF users who are using high-end instruments to produce large datasets. Those users are facing the demands of a new wave of data-intensive instruments and software that enable: higher spatial resolution; higher chemical resolution; 3D and 4D+ approaches; more rapid dynamic processes; and multi-dimensional analyses.
A specific analysis platform being developed is the Atom Probe Workbench. This tool is a component of a larger national eResearch project in Australia, that is aiming to integrate existing tools and techniques with a network of specialised cloud-based computing systems and data-storage facilities. This integration will enable the atom probe research community to access and create valuable tools, accelerating the research process.


The authors acknowledge support from the National Collaborative Research Infrastructure Strategy; National eResearch Architecture Taskforce; Office of Learning and Teaching; and National eResearch Collaboration Tools and Resources (NeCTAR) project.

Type of presentation: Poster

IT-15-P-1564 MyScope: On-Line Microscopy and Microanalysis Training and Education in Core Facilities

Apperley M. H.1, Munroe P. R.1, White T.2, Shapter J.1, Muhling J.1, Soon L.1, Ringer S. P.1, Grinan E.1, Frost C.1, Cribb B.1
1Australian Microscopy and Microanalysis Research Facility, Sydney, Australia, 2School of Materials Science and Engineering, Nanyang Technological University, Singapore
miles.apperley@sydney.edu.au

The Australian Microscopy and Microanalysis Research Facility (AMMRF) is a national grid of equipment, instrumentation and expertise in microscopy and microanalysis that provides nanostructural characterisation capability and services, from widely used optical, electron, X-ray and ion-beam techniques to world-leading flagship platforms. One of the principal activities of the AMMRF is to provide research training in microscopy and microanalysis. Until recently, much of this training was provided either in the classroom or through one-to-one training at the instrument itself, however, these approaches faced limitations.

Firstly, the large number of researchers requiring training places pressures on the core facilities to balance the need to maximise the beam-time of expensive and complex instrumentation for research purposes with that for training new users, who will ultimately perform the research. Prioritising instrument time for research reduces time available for training and vice versa.

Another common challenge in such facilities is diversity of the student body. In the case of our project the cohort requiring training had a variety of backgrounds and goals: undergraduate students with different educational backgrounds seeking an overview of topic; final year students requiring specific techniques for project work; future career or postgraduate students; and also professional researchers, educators and managers. A more flexible approach to training and education was needed.

To address these challenges and improve the training outcomes of researchers, the AMMRF developed MyScope: Training for Advanced Research. MyScope is an online suite of education tools for teaching and learning in the area of microscopy and microanalysis. The modules in MyScope provide a novel advancement in online training. They contain a number of components including: an interactive questionnaire to allow the user to assess their knowledge, guide choices and tailor the learning environment for flexible learning; also, tailoring capability for academics and trainers; self guided tutorials with videos, animations and glossary to prepare students with knowledge and specialist language; virtual instrument platforms to practice use of instrumentation; and online competency testing to demonstrate readiness for hands-on experience.


The authors acknowledge funding from the Office for Learning and Teaching, Australian Government Department of Education, CG10-1490.

Type of presentation: Poster

IT-15-P-2171 Spectroscopic imaging with a new pnCCD camera: improved dynamic range, position resolution, anti-blooming and compactness

Ihle S.1, Eckhardt R.2, Hartmann R.1, Holl P.1, Huth M.1, Kalok D.1, Ryll H.1, Schmidt J.1, Simson M.2, Soltau H.2, Soltau J.2, Steigenhöfer D.1, Thamm C.2, Strüder L.1
1PNSensor GmbH, Munich, Germany, 2PNDetector GmbH, Munich, Germany
sebastian.ihle@pnsensor.de

pnCCDs are well known as radiation detectors for spectroscopic imaging for X-rays in many fields of science: X-Ray Fluorescence analysis (XRF), X-ray astronomy, X-ray Free Electron Laser science and at synchrotrons. pnCCDs are radiation detectors on high resistivity 450 µm fully sensitive silicon. They are back-illuminated, with a thin, homogeneous radiation entrance window, enabling the detection of X-rays from 30 eV up to 30 keV with high quantum efficiency. As all pnCCDs are equipped with a fully column parallel readout, frame rates of more than 1.000 frames per second are achieved, keeping the read noise level at 3 electrons. Some of the key performance figures are e.g. a quantum efficiency above 90% from 1 keV up to 10 keV, extreme radiation hardness, operation at temperatures around -20 °C or warmer, energy resolution of less than 130 eV (FWHM) at 6 keV and 37 eV (FWHM) at 90 eV. These properties have enabled a variety of spectacular measurements. The following improvements were made recently:

(a) New Colour X-ray camera module: The new CXC module (see Fig. 1) was designed to fit in small and tight surroundings. The new camera can be operated in vacuum without any entrance filter or with a Be filter in normal environments. Capillary optics can be coupled to the entrance window.

(b) High dynamic range mode: The previously applied standard operating modes were able to handle about 300.000 electrons in a single pixel. The new settings allow confining and transferring more than 2.5 million electrons in the CCD.

(c) Controlled charge extraction: If the amount of signal charge overcomes the charge handling limit the surplus charges can be taken out in a controlled way to avoid overflowing electrons to spoil the information content of the neighbouring pixels. A direct electrical access to the pixels allows to define a saturation level of the pixels (anti-blooming). We have tested this mechanism experimentally with a charge load of 2 billion electrons per pixel. An example is shown in Fig 2.

(d) Subpixel resolution: The low noise of the pnCCD system enables to centroid the signal charge cloud with a position precision of 2.5 µm (rms) (see Fig 3). This is achieved by increasing the charge cloud diameter by reducing the electric field during the charge collection process and therefore increasing the charge diffusion process.

All the above improvements are delivering new qualities to the compact X-ray camera stimulating new methods for spectroscopic imaging measurements.


The authors would like to thank the technical staff of PNSensor and PNDetector for their outstanding support.

Fig. 1: Photo of the new Colour X-ray Camera (CXC) module.

Fig. 2: Demonstration of anti-blooming operation. The pnCCD is illuminated with a bright spot. In normal operation (left) the signals spills into the neighboring pixels. In anti-blooming mode the excess charge is removed in a controlled way.

Fig. 3: Measured position precision as a function of operating parameters.

Type of presentation: Poster

IT-15-P-2231 Multilayer Laue Lenses as High Resolution hard X-ray Optics

Kubec A.1, 2, Niese S.1, Melzer K.1, 2, Braun S.2, Patommel J.1, Leson A.2, Zschech E.3
1Technische Universität Dresden, Dresden, Germany, 2Fraunhofer IWS Dresden, Dresden, Germany, 3Fraunhofer IKTS-MD, Dresden, Germany
adam.kubec@tu-dresden.de

X-ray Microscopy provides a high resolution sample analysis method while demands on sample preparation are typically less restrictive than for transmission electron microscopy. Resolutions are still limited by the available optics but improvements in this field can easily be implemented in existing setups and will thus directly influence and improve measurements. Multilayer Laue lenses (MLL) are promising x-ray optics to achieve high efficiency focusing with small spot sizes down to sub-10 nm with current and down to sub-1 nm with improved geometries.
We have deposited multilayer stacks with layer thicknesses according to the zone plate law and a total deposition thickness of more than 50 micrometer. Deposition processes using magnetron sputtering took up to 83 hours and the stacks containing up to 6500 individual layers with thicknesses down to 5 nm have been obtained.
These parameters require long term system process stability on the one hand and a high precision in zone deposition on the other hand. The actual one dimensionally focusing lens is subsequently fabricated with mechanical preparation and focused ion beam milling. Two of these lenses have to be placed perpendicularly in a distance of about 30 micrometers from each other to obtain point focusing [Fig. 1]. Two of these lenses were crossed perpendicularly at a distance of 30 µm to obtain a point focus [Fig.1].
We have successfully demonstrated several focusing and imaging experiments using crossed MLLs. In synchrotron beam times at ID13, ESRF in Grenoble, France and P06, PETRA III in Hamburg, Germany spot sizes down to 39x49 nm2 at 20 keV x-ray energy have been shown using the coherent diffraction imaging method of Ptychography [Fig. 2 and 3]. Furthermore we have demonstrated measurements with a wedged MLL realizing the improved geometry. Global diffraction efficiency was enhanced by more than 50% on average over the entire aperture of the lens. In addition, full field imaging was shown for the first time using multilayer Laue lenses at a laboratory microscope with a rotating copper anode.

References:
[1] A. Kubec, S. Braun, S. Niese, P. Krüger, J. Patommel, M. Hecker, A. Leson and C. Schroer: Ptychography with Multilayer Laue Lenses and their Initial Characterization with a Laboratory Based X-ray Microscope, to be published


The work has been supported by the BMBF within the cool silicon project and is partly funded by the European Regional Development Fund and the Free State of Saxony via the ESF project 100087859.
Portions of this research were carried out at the light source PETRA III at DESY and on the ID13 beamline at the European Synchrotron Radiation Facility (ESRF), Grenoble, France.

Fig. 1: A pair of crossed multilayer Laue lenses with the designated beam direction.

Fig. 2: A phase reconstruction of the test sample used for the ptychography measurements. [1]

Fig. 3: Amplitude reconstruction of the focal plane. [1]

Type of presentation: Poster

IT-15-P-2249 Multilayer Laue Lenses as High Resolution High Efficiency X-ray Focusing Optics

Kubec A.1, 2, Niese S.1, Melzer K.1, 2, Braun S.2, Patommel J.1, Leson A.2
1Technische Universität Dresden, Dresden, Germany, 2Fraunhofer IWS, Dresden, Germany
adam.kubec@tu-dresden.de

Multilayer Laue Lenses (MLL) are a promising approach based on diffraction to focusing hard x-rays and promise to open the path to nanometer spot sizes [1]. Limitations implied by the fabrication process of zone plates regarding possible zone widths and aspect ratios are circumvented. Using thin film deposition techniques alternating zones of two different materials are deposited onto a flat substrate with thicknesses according to the zone plate law. A lamella then cut out of the coating using Focused Ion Beam milling. This segment is the actual lens and produces a focal line. Combined with a second perpendicularly aligned lens a point focal is achieved. The structures accommodating the lenses are glue-bonded directly onto each other. The distance between the lenses is approximately 30 µm (fig. 1). The pair of lenses is then fixed onto a single mount. Only two precise tilting stages are necessary as well as two stages for coarse position alignment. At the ESRF beamline ID13 and the PETRA III beamline P06 we have shown such setups of pairs of crossed MLLs. The lenses were characterized using Ptychography [2]. According to the reconstructions of the complex wave field focal spots with a FWHM of about 50 x 50 nm2 and less have been achieved. In addition the local diffraction efficiency of a wedged MLL was compared to a regular tilted geometry lens. The results show an increase of intensity in the first focusing order of more than 30%. Particularly the local diffraction efficiency of the zones with less than 10 nm zone width increased noticeably.

References
[1] J. Maser et al., Optical Science and Technology, the SPIE 49 Annual Meeting (2004)
[2] S. Hönig et al., Optics Express, Vol. 19, Issue 17, pp. 16324-16329 (2011)

 


The work has been supported by the BMBF within the cool silicon project and is partly funded by the European Regional Development Fund and the Free State of Saxony via the ESF project 100087859.
Portions of this research were carried out at the light source PETRA III at DESY and on the ID13 beamline at the European Synchrotron Radiation Facility (ESRF), Grenoble, France.

Fig. 1: Process photograph of the Magnetron Sputter Deposition.

Fig. 2: SEM image of a pair of crossed MLL with a distance of approximately 30 µm.

Type of presentation: Poster

IT-15-P-2471 THERMODYNAMİCS OF ZİRCON AND SOLİD SOLUTİON REACTIONS

Kılıç A. D.1, Ateş C.2
1Fırat University Geology Department, Elazığ, Turkey
didem7399@hotmail.com


Zircon (ZrSiO4) from granitic gneisse and ortho gneisse in the Pütürge metamorphites were mineralogically characterized by inductively couples plasma mass spectrometry (ICP-MS), X-ray powder diffraction and Cathodolüminescence (CL) analyses show that the zircon grains have developed isostructural solid solutions with coffinite (USiO4), throite (ThSiO4). These zircons have different thermodynamic and processing. This processes are chemical reequilibration of crystalline zircon solid solutions. Zircons have textural and chemical variety which presents characteristic of metamict and partly metamict. Metamict zircons are high REE, U, Th content. Whereas partly zircons are more low REE, U, Th. The chemical characteristic of zircons can produce both aqueous fluids and melts. İn the zircons are observed by porous structure and inclusion rich spaces. The inclusion rich and porous zircon grains are featured by lower concentrations of trace elements. This were interpretationed dissoution-reprecipitation process intrastructure of zircons. This show that its reacts with an aqueous fluid, dissolution-reprecipitation process will produce more less trace elements than the other zircons.
İn order to constrain the timing of metamorphism, 39Ar/40Ar dating were performed on four biotite. The four samples is 83.21±0.069 Ma. Accordingly, greenschist and amphibolite facies metamorphism occured at Santonien. U-Pb cristallization age from zircon probably correspond to the timing of fluid influx or anatexis rather than to the age of peak metamorphic conditions.
Keywords: Solid solution inclusions, zircon, high metamorphism, CL analyses, 39Ar/40Ar biotite dating, U-Pb isotope


Type of presentation: Poster

IT-15-P-2661 Evaluating the detection of cracks in underwater wet welds by x-ray microtomography

Paciornik S.1, Silva L. F.1, Santos V. R.1, Bernthaler T.2
1DEMa PUC-Rio - Brazil, 2Institut für Materialforschung - Hochschule Aalen - Germany
sidnei@puc-rio.br

Despite efforts to improve the mechanical properties of wet welds, welding in direct contact with water still presents critical problems. High cooling rates and the presence of hydrogen derived from water dissociation leads to the formation of defects, such as pores and cracks in the weld metal (WM) and in the heat affected zone (HAZ) which adversely affect mechanical properties. The occurrence of hydrogen assisted cold cracking is considered as one of the most important factors for the usual low ductility in the WM [1, 2].
During cooling, weld beads contract both in transverse and longitudinal directions. It is well established that longitudinal contractions are responsible for higher residual stress after welding. As a consequence, the low WM toughness associated with hydrogen embrittlement can lead to nucleation of cracks with a predominant orientation transverse to the weld axis.
In previous works, both Optical Microscopy (OM) [3] and x-ray microCT [4,5], linked to Image Analysis (IA) have been used to characterize crack size, density, shape and orientation. Cracks are challenging to image and measure from microCT due to their strong anisotropy and the small dimension of the tip, which can go beyond the resolution of microCT.
In this work, a weld sample with varying cross-section (Fig. 1) was imaged by microCT before and after being submitted to a tensile test up to failure. The variation in strain due to varying thickness led to changes in crack tip opening and crack length. Fig. 2 shows a typical reconstructed layer of the weld with cracks. The presence of noise and limited contrast hinder the detection of cracks. After noise filtering and segmentation the cracks were rendered, as shown in Fig. 3, and measured, allowing the estimation of detection limits for this kind of defect by microCT.

REFERENCES

1. A. Q. BRACARENSE et al., “Comparative study of commercial electrodes for underwater wet welding”, (In International Congress of the International Institute of Welding, São Paulo, 2008).
2. V. R. SANTOS et al., “Recent Evaluation and Development of Electrodes for Wet Welding of Structural Ship Steels” (In 29th International Conference on Ocean, Offshore and Arctic Engineering, Shangai, 2010).
3. M. H. P. MAURICIO et al., “Quantitative Hydrogen Cracking Evaluation by Digital Optical Microscopy” (In IMC17, Rio de Janeiro, 2010).
4 . S. PACIORNIK ET al. “Characterization of Pores and Cracks in Underwater Welds by ct and Digital Optical Microscopy”. Proc. 1st International Conference on 3D Materials Science, 2012. p. 177-182.
5. PADILLA, E. et al . Image analysis of cracks in the weld metal of a wet welded steel joint by three dimensional (3D) X-ray microtomography. Mat Charac, 83, 139-144, 2013.


The financial support of CNPq, CAPES and FINEP, Brazilian agencies is gratefully acknowledged. 

Fig. 1: Tensile test specimen with varying cross section (dimensions in mm).

Fig. 2: Reconstructed microCT image of part of the sample in Fig. 1, revealing cracks.

Fig. 3: 3D rendering of cracks, after noise filtering and segmentation.

Type of presentation: Poster

IT-15-P-5855 Nanoscale characterization of hierarchical biological materials using synchrotron quantitative scanning-SAXS imaging

Gourrier A.1,2,3, Burghammer M.3, Reiche I.4, Boivin G.5
1Univ. Grenoble Alpes, LIPHY, F-38000 Grenoble, France, 2CNRS, LIPHY, F-38000 Grenoble, France, 3European Synchrotron Radiation Facility (ESRF), Grenoble, France, 4Laboratoire d’Archéologie Moléculaire et Structurale, UMR 8220 CNRS Université Pierre et Marie Curie, Sorbonne Universités, Ivry-sur-Seine, France., 5INSERM U1033, Université de Lyon, France
aurelien.gourrier@ujf-grenoble.fr

The characterization of biological materials often proves challenging due to their high degree of structural hierarchy and their composite nature at the nanoscale. Bone is a typical example which presents an additional level of complexity because of the variety of morphologies encountered from the nanometer to the centimeter scale. This stems from the physiological processes associated with the synthesis, mechanical adaptation to external loads and self-healing. As a result, there is a growing consensus in the biomedical field over the necessity of multiscale approaches for the evaluation of the effects of bone pathologies. The molecular and supra-molecular levels, in particular, are currently receiving a lot of attention. At these scales, bone consists of complex arrangements of collagen microfibrils mineralized with calcium phosphate nanoparticles. Precisely how this nanoscale organization affects the mechanical properties of the higher hierarchical levels is still poorly understood.

In this paper, we will highlight the potential of quantitative scanning-SAXS imaging1,2,5 for such studies. This technique relies on scanning a sample with a monochromatic X-ray beam much smaller than the sample dimensions (typically 100 nm- 10 μm), and recording the scattered intensity in forward geometry. The images acquired at small scattering angles (SAXS) provide atomic to nanoscale resolution. They are reduced to scalar values by various algorithms based on the theory of SAXS and mapped as a function of scan coordinates to produce the final images. Using state-of-the-art X-ray optics and detectors with synchrotron sources, nanoscale fluctuations in density within a size range of ~1-100 nm can be mapped with very high spatial resolution over sample regions comparable to histology (cm2). This new method is therefore highly competitive and bridges the gap between TEM or AFM and high resolution microscopies.

Various results will be presented from fundamental, biomedical and archaeological studies to demonstrate the potential of this method. In particular, the size, organization and orientation of the mineral nanoparticules in bone will be described in various healthy and pathological/altered conditions.

Various results will be presented from fundamental, biomedical2,4 and archaeological3 studies to demonstrate the potential of this method. In particular, the size, organization and orientation of the mineral nanoparticules in bone will be described in various healthy and pathological/altered conditions.


Fig. 1: qsSAXS Image of the particle size (nm) of a thin section (6 (H) x 10 (V) mm2 x 50 μm) illiac crest biopsy of a sheep model.

Type of presentation: Poster

IT-15-P-5878 A Hard X-ray Nanoprobe at Diamond Light Source

Cacho-Nerin F.1, Parker J. E.1, Peach A.1, Wilkin G.1, Quinn P.1
1Diamond Light Source, Harwell Science and Innovation Campus, Didcot, Oxon. OX11 0DE, UK
fernando.cacho-nerin@diamond.ac.uk

Beamline I14 is the hard X-ray nanoprobe beamline currently under construction at Diamond Light Source in Oxfordshire, UK. It is scheduled to come into operation in 2017. The beamline will be a dedicated facility for nanoscale microscopy and micro-nano SAXS, serving two endstations housed in a new external building approximately 175m from the main synchrotron ring. The nanoprobe endstation aims to achieve the smallest possible focus (initial aim 50nm) with the capability to exploit future optics developments. The optical design is optimised for scanning X-ray fluorescence, X-ray spectroscopy and diffraction. The mesoprobe endstation will be optimised to carry out simultaneous small and wide angle X-ray scattering studies as well as scanning fluorescence mapping, with a variable focus beam in the range 5µm – 100 nm. The beamline will complement electron and optical microscopy and enable new science in a number of areas spanning materials science, biology, engineering and earth science.

The I14 beamline will be housed in the same building as the new UK national electron microscopy facility, which provides 4 state-of-the-art electron microscopy suites covering the physical and life sciences. This facility combines staff and expertise from a number of different areas which we believe will allow us to make exciting progress in sample preparation techniques and correlative x-ray and electron microscopy studies. Here we present the design and key specifications of Beamline I14, and highlight potential applications.


Type of presentation: Poster

IT-15-P-6022 Microscopy in the extreme ultraviolet and soft X-ray spectral region and its applications

Herbert S.1, Danylyuk S.1, Loosen P.1, Maryasov A.2, Wilson D.2, Bußmann J.2, Rudolf D.2, Juschkin L.2, Küpper L.3, von Wezyk A.3, Bergmann K.3, Lebert R.4
1Chair for the Technology of Optical Systems, RWTH Aachen University and JARA - Fundamentals of Future Information Technology, 2Chair for the Experimental Physics of EUV, RWTH Aachen University and JARA – Fundamentals of Future Information Technology, 3Fraunhofer Institute for Laser Technology, 4Bruker ASC
stefan.herbert@ilt.fraunhofer.de

Since the invention of first optical microscopes, applications pushed the further development of microscopes. Constant push for higher resolution leads to continuous decrease of the working wavelength of microscopes from visible light to UV, VUV and finally to EUV and X-ray spectral regions. Numerous applications of such short-wavelength microscopes are being investigated in Aachen. In the past decade we have built several industrially relevant microscopes, based on laboratory plasma sources, for different applications and demonstrated their potentials. In this contribution a transmission microscope for 13.5nm wavelength (Fig.1), a reflection dark field microscope for 13.5nm (Fig.2), a water window microscope for 1-5nm (Fig.3) and a lensless Microscope for 17.3nm will be presented and discussed in detail.

Hereby the microscopes for 13.5nm have been developed for tasks, connected to deployment of the upcoming extreme ultraviolet (EUV) lithography and therefore are designed for investigations of defective multilayer mirrors or thin films. The use of multilayer mirrors in EUV lithography as imaging optics requires an actinic (at wavelength) inspection, as defects inside the multilayer stack could not be detected with surface techniques. Moreover, the independence of such an inspection tool from a synchrotron and a flexible table top design is crucial for industrial realization of EUV lithography.

Another application in the short wavelength region, which becomes more and more applicable as the power of computers increases is lensless imaging. The benefit of not having manufacturing limited optics in the beam path and with that having a diffraction limited microscope strikes the malus of required high computational power. The application of lensless microscopy with incoherent plasma sources has to our knowledge not been realized elsewhere.

For investigations of organic samples a microscope in the water window has been developed. In this spectral region of 1-5nm carbon and phosphor are absorbing and oxygen is transmitting the light, which gives excellent contrast for water-based samples. Extension of the microscope from 2D to tomographic measurements will be discussed.

Furthermore we present concepts and first experiments of a time resolved microscope for 17.3nm and a microscope, which can measure the magneto-optical contrast of materials at absorption edges, e.g. cobalt 3p at 20.3nm (Fig.4). The time resolved microscope targets sub-100nm spatial resolution and a time resolution of around 4ns, achieved by a triggered micro-channel plate. The microscope for magneto-optical investigations is for the first time realized using a laboratory plasma source, which can enable broad application of the technique, not limited to high brilliance synchrotrons.


Fig. 1: Extreme ultraviolet Schwarzschild-objective based transmission microscope for operation in brightfield or darkfield mode with an optional second magnification step, achieving a spatial resolution of around 100 nm.

Fig. 2: Extreme ultraviolet Schwarzschild-objective based darkfield reflection microscope for defect inspection of mask blanks with high troughput and moderate spatial resolution

Fig. 3: Laboratory zone plate based water window soft x-ray microscope, achieving a spatial resolution of around 40 nm

Fig. 4: Extreme ultraviolet microscopic setup for experiments on magneto-optical contrast of elements

IT-16. Electron microscopy theory and simulations

Type of presentation: Invited

IT-16-IN-1783 Calculation and Simulation in Determining Site-specific Magnetic Structure by Dynamical Electron Diffraction-EMCD

Zhu J.1, Wang Z. Q.1, Song D. S.1, Zhong X. Y.1, Yu R.1, Cheng Z. Y.1
1National Center for Electron Microscopy in Beijing, School of Material Science & Engineering, Tsinghua University, Beijing 100084, China
jzhu@mail.tsinghua.edu.cn

Quantitatively determining the magnetic structure of material on a nanometer scale is a potential task for future transmission electron microscope (TEM). Site-specific electron energy-loss magnetic chiral dichroism (site-specific EMCD) method is come up with to get the crystallographic site-specific magnetic information of nanostructures.[1-2]
This presentation will briefly introduce how we process calculations, simulations and experiments for determining the site-specific magnetic structure of a nanostructure of NiFe2O4. By constructively combining using the dynamical electron diffraction and EMCD methods, to calculate the coefficients of Bloch waves and draw out the relative EMCD intensity mappings in momentum space, in which the effect of asymmetry of the dynamical electron diffraction needs to be considered;[3] then to select the optimum experimental parameters in EMCD experiments, by adjusting dynamical diffraction conditions to enhance site-specific EMCD signals; with sample’s site-specific magnetic circular dichroism spectra, and the site-specific spin/orbital magnetic moments extracted.
Compared with X-ray magnetic circular dichroism, the site-specific EMCD method shows its unique capability for solving the crystallographic site-specific magnetic structure on nano-scale.
Reference:
[1] Schattschneider P, Rubino S, Hebert C, et al. Detection of magnetic circular dichroism using a transmission electron microscope. Nature 441, 486–488 (2006).
[2] Wang ZQ, Zhong XY, Yu R, Cheng ZY, Zhu J. Quantitative experimental determination of site-specific magnetic structures by transmitted electrons. Nature Communications, 4, 1395 (2013).
[3] Dongsheng Song, Ziqiang Wang and Jing Zhu, Effect of the asymmetry of dynamical electron diffraction on intensity of acquired EMCD signals, unpublished.


This work is financially supported by National 973 Project of China and Chinese National Nature Science Foundation. This work made use of the resources of the Beijing National Center for Electron Microscopy.

Type of presentation: Invited

IT-16-IN-1882 Thermal magnetic field noise and electron optics - more experiments and calculations

Uhlemann S.1, Müller H.1, Zach J.1, Berger C.1, Haider M.1
1CEOS Corrected Electron Optical Systems, Heidelberg, Germany
uhlemann@ceos-gmbh.de

The simultaneous correction of the spherical (Cs) and the chromatic aberration (Cc) in transmission electron microscopy (TEM) has been implemented for a broad range of beam voltages: 20-300kV [1-3]. In such an instrument the effects of the lateral and temporal incoherence of the illuminating electron beam are largely suppressed. The measured remaining focus spread for instance is by far small enough to allow information transfer beyond g=20/nm - even at 300kV where Cc-correction is most challenging [1]. However, during the development of the corrector hardware we recognized, that an additional incoherence mechanism deteriorates the contrast of the recorded images. By careful measurements of the contrast transfer we found that an envelope function of the form exp[-2(πσ|g|)2] perfectly matches the observations. It turned out, that an isotropic image spread σ reduces the image contrast. Image spread can be understood as a stochastic, high-frequency image displacement during acquisition. Recently, it could be proven experimentally that the origin of this image spread is magnetic field noise (Johnson-Nyquist noise) emitted from the conducting parts around the electron beam. The thermodynamic nature of this noise was clearly demonstrated by cooling beam tubes (made from stainless steel or permalloy) from room temperature down to liquid nitrogen temperature [4].

In the experiments we measure the standard deviation σ of this image shift. Its variance σ2 is proportional to the product of the field correlation length ξ along the path and the variance <B2> of the transversal magnetic field [4]. Here, we report on the progress we made to understand the experimental results theoretically.

Surprisingly, magnetic materials (μr>1) introduce more integral noise than non-magnetic materials like stainless steel. Hence, we were very much interested in the common situation were magnetic material is placed outside a liner tube made from stainless steel, see Figure 1. The question arose, if the thin stainless steel tube is transparent for the stronger noise emitted from the magnetic components. Here we also report on the experiment “tube-in-tube”: A thin-walled (0.15mm) stainless steel liner tube with 3mm outer diameter is placed in a stack of permalloy tubes, see Figure 1. Beside numerical strategies, a semi-analytical approach to understand the compound system is presented, see Figures 2+3.

After all, electron optical design - especially the design of extended corrector optics - has to take into account the existence of magnetic field noise emitted from the conducting parts. We discuss scaling rules and why hexapole-type aberration correctors are collecting less image spread from thermal magnetic field noise than quadrupole-octupole-type Cc-correctors.


References:

[1] M. Haider et al, Microsc.Microanal. 16 (2010) p. 393.
[2] H. Sawada et al, AIEP 168 (2011) 297.
[3] http://www.salve-project.de
[4] S. Uhlemann et al, Phys.Rev.Lett. 111 (2013) 046101.

Fig. 1: Experiment to compare the thermal magnetic field noise of a thin stainless steel liner tube with the compound system. Permalloy tube fitted over a 3mm stainles steel liner tube (a), tube stack and outer holder tube (b), end view of the compound sample (c), dimensions (d), copper cooler and two samples: bare liner tube and the compound sample (e).

Fig. 2: Theoretical treatment by means of the fluctuation-dissipation theorem. The power-loss induced within a conducting structure by a fluctuating magnetic dipole is calculated. In rare cases with high symmetry Maxwell's equations can be solved directly by separating variables: A sheet of non-magnetic material in contact with a magnetic half-space.

Fig. 3: For thin conducting sheets (thickness t, resistivity ρ) a boundary-element method for a triangular mesh covered with t/ρ is the preferable way to calculate the frequency spectrum of the magnetic field noise.

Type of presentation: Oral

IT-16-O-1476 Numerical Treatment of the Full, Non-Approximated Schrödinger Equation at Low Energies

Wacker C.1, Schröder R. R.1
1CryoEM, CellNetworks, Universitätsklinikum Heidelberg, Germany
christian.wacker@bioquant.uni-heidelberg.de

The imaging of samples at energies as low as 20 kV has attracted a lot of attention, recently [1]. With decreasing accelerating voltage even chemical elements with low atomic numbers (e.g. carbon, oxygen, nitrogen) must be treated as strong scatterers. Hence, simulations of the image formation process of organic molecules must include a correct description of high-angle scattering and inelastic scattering processes. In this work we present a new approach to deal with the first problem.

The conventional multislice (CMS) algorithm is derived from the Schrödinger equation using the high-energy or paraxial approximation. This approximation neglects the second derivative of the wave function along the optical axis and thereby replaces the Ewald sphere by a parabola. However, a careful mathematical analysis of the full Schrödinger equation allows developing a rigorous multislice (RMS) scheme without resorting to the high-energy approximation [2]. Recent implementations and numerical analyses of the RMS scheme demonstrated, that the CMS method is not accurate for accelerating voltages below 100 kV [3,4].

Therefore, it is interesting to note, that a numerical solution of the full Schrödinger equation without sophisticated mathematical treatment is also possible. Essentially, the Schrödinger equation can be regarded as a second-order ordinary differential equation. This enables us to use well-known numerical algorithms like the classical Runge-Kutta method. Similar to the multislice algorithms the wave is incrementally propagated through the sample. But instead of using integrated atomic potentials, the Runge-Kutta method integrates the potentials on-the-fly. Fig. 1 shows the amplitude-diffraction patterns of SmBa2Cu3O7-x calculated with this method at three different accelerating voltages. For reference, fig. 2 depicts the same sample but calculated with the CMS method implemented in real space. Besides, the Runge-Kutta method can also be applied to the high-energy approximation of the Schrödinger equation yielding comparable results to the CMS method (fig. 3).

<span>The computational effort increases with decreasing accelerating voltage, as the step size must be reduced to compensate for the higher scattering angles. Thus, we are currently investigating different implementations and parallelization approaches in order to reduce the required wall time and to allow for more complex samples.

[1] U. Kaiser, J. Biskupek, J.C. Meyer, J. Leschner, L. Lechner, H. Rose, M. Stöger-Pollach, A.N. Khlobystov, P. Hartel, H. Müller, M. Haider, S. Eyhusen, G. Benner, Ultramicroscopy, 111 (2011) 1239-1246

[2] J.H. Chen, D. van Dyck, Ultramicroscopy 70 (1997) 29-44

[3] C.Y. Cai, J.H. Chen, Micron 43 (2012) 374-379

[4] W.Q. Ming, J.H. Chen, Ultramicroscopy 134 (2013) 135-143


CW gratefully acknowledges the Studienstiftung des Deutschen Volkes for a PhD scholarship.

Fig. 1: Amplitude-diffraction patterns for SmBa2Cu3O7-x calculated by applying the Runge-Kutta method to the full Schrödinger equation.

Fig. 2: Amplitude-diffraction patterns for SmBa2Cu3O7-x simulated using the conventional multislice algorithm (CMS) in real space.

Fig. 3: Amplitude-diffraction patterns for SmBa2Cu3O7-x calculated by applying the Runge-Kutta method to the Schrödinger equation in the high-energy approximation.

Type of presentation: Oral

IT-16-O-1624 Possible tuned laser boosting of spatial and energy resolution in EELS

Howie A.1
1University of Cambridge, Cambridge CB3 0HE , UK
ah30@cam.ac.uk

The potential advantage of combining the spatial resolution of electron energy loss spectroscopy (EELS) with the spectral resolution of photons has long been apparent [1] but, apart from cathodoluminescence studies [2,3], has made little progress.  More striking has been photon-induced electron microscopy (PINEM) pioneered primarily as a pump-probe technique for time-resolved imaging capability [4].  Fast electrons passing close to a carbon nanotube in coincidence with an intense pulse of laser illumination at frequency ω, experienced energy losses and gains nhω (-5 < n < 5).  These results were explained through the e-beam interaction in the near field region of the wave emitted by the nanotube in response to the laser pulse [5].

Sacrificing the time resolution of pulsed operation, more systematic exploration of the dielectric resonances of nanostructures could be provided by combining continuous tuned laser illumination with EELS (fig. 1). The z-dipole as well as the x-dipole shown here could be used. Initial computations [6] suggest, at not too high laser power, a substantial boosting of EELS signals with its own dipolar angular dependence (fig.2). Non-linearity could limit laser pumping of object boson oscillator modes (eg plasmons and phonons) but could result in stimulated emission more than proportional to laser intensity.  Studies of the loss and gain intensities as a function of laser power and frequency could usefully clarify the basic physics involved here and show how the near field mechanism combines with the usual excitation theory of EELS [7].  Significant boosting of EELS losses could obviously be useful for very weak losses and also for probing otherwise inaccessible Raman or catalytic hot spots (fig. 3).  More generally it could counter the severe loss of intensity experienced when the minimum momentum transfer in low loss EELS is increased in order to improve spatial resolution [8].  A minimum in lateral momentum tranfer hqmin can be set by off-axis spectroscopy or better by use of STEM hollow cone illumination with pre-spectrometer lens tuning to match aperture gaps precisely (fig. 4).

[1] Howie A, (1999) Inst. of Physics Conf. Series 161, 311.

[2] Yamamoto Y, Araya K and Garcia de Abajo FJ, (2007) Phys. Rev. B 64, 205219.

[3] Tizei LHG and Kociak M, (2013) Phys. Rev. Lett. 110, 153604.

[4] Barwick B, Flannigan DJ and Zewail AH, (2009) Nature 409, 902.

[5] Garcia de Abajo FJ and Kociak M, (2008) New J. Phys. 10, 073035.

[6] Adenjo-Garcia A and Garcia de Abajo FJ, (2013) New J. Phys. 15, 103021.

[7] Talebi N, Sigle W, Vogelgesang R and van Aken P, (2013) New J. Phys. 15, 053013.

[8] Muller DA and Silcox J, (1995) Ultramicroscopy 59, 195. 


I thank Professor Javier Garcia de Abajo for several illuminating discussions.

Fig. 1: Schematic representation of laser excitation of the x-polarised dipole mode of a nano particle.  The emitted dipole radiation interacts with an electron travelling in the z direction. 

Fig. 2: Contour plot (one quadrant only) of the laser stimulated x-dipole EELS loss intensity in the x-y plane.  The angular dependence determined by the laser E-field is not observed for the usual e-beam stimulated losses.

Fig. 3: With appropriate choice of laser beam direction and polarisation it may be possible to generate sufficient excitation for e-beam energy losses and gains to be detected at hot spots where normal EELS excitation is impossible or very weak.

Fig. 4: The use of an annular aperture for STEM hollow cone illumination and a non-overlapping EELS collection aperture defines a minimum lateral momentum transfer hqmin.  The usual loss excitation profile for the z-polarised dipole (blue) is thereby more localised (red) for qmin = 3ω/v. 

Type of presentation: Oral

IT-16-O-1997 Validities of three multislice algorithms for quantitative low-energy transmission electron microscopy

Ming W. Q.1, Chen J. H.1
1Hunan University, Changsha, Hunan, China
suokesi_ming@163.com

Three different types of multislice algorithms, namely the conventional multislice (CMS) algorithm, the propagator-corrected multislice (PCMS) algorithm and the fully-corrected multislice (FCMS) algorithm, have been evaluated in comparison with respect to the accelerating voltages in transmission electron microscopy. Detailed numerical calculations have been performed to test their validities. The results show that the three algorithms are equivalent for the accelerating voltage above 100 kV. However, below 100 kV, the CMS algorithm will introduce significant errors, not only for higher-order Laue zone (HOLZ) reflections but also for zero-order Laue zone (ZOLZ) reflections. The differences between the PCMS and FCMS algorithms are negligible and mainly appear in HOLZ reflections. Nonetheless, when the accelerating voltage is further lowered to 20 kV or below, the PCMS algorithm will also yield results deviating from the FCMS results. The calculation efficiency of the PCMS is illustrated to be much higher than the FCMS, while the accuracy of numerical calculation can be compared to the CMS. Therefore, the PCMS can be an alternative for fast simulation of HRTEM images. The present study demonstrates that the propagation of the electron wave from one slice to the next slice is actually cross-correlated with the crystal potential in a complex manner, such that when the accelerating voltage is lowered to 10 kV, the accuracy of the algorithms is dependent of the scattering power of the specimen.


This work is supported by the National Basic Research (973) Program of China (No. 2009CB623704); the National Natural Science Foundation of China (No. 51171063, 51071064); Instrumental Innovation Foundation of Hunan Province (No. 2011TT1003); PhD Programs Foundation (20120161110036).

Type of presentation: Oral

IT-16-O-2356 Measuring structure parameters from electron microscopy images : what are the limits?

Van Aert S.1, Gonnissen J.1, De Backer A.1, Sijbers J.2, den Dekker A.2,3
1Electron Microscopy for Materials Science (EMAT), University of Antwerp, Antwerp, Belgium, 2iMinds-Vision Lab, University of Antwerp, Antwerp, Belgium, 3Delft Center for Systems and Control, Delft University of Technology, Delft, The Netherlands
sandra.vanaert@uantwerpen.be

State-of-the-art electron microscopy combined with advanced model-based methods can provide reliable numbers for unknown structure parameters. Aberration correction greatly improves the quality of experimental images, new STEM data collection geometries allow one to visualise light atoms, and detectors start to behave as ideal quantum detectors. In combination with statistical parameter estimation theory to analyse experimental data, electron microscopy then performs at its ultimate limits. For example, atomic column positions can be measured down to picometer scale precision [1], differences in averaged atomic number of only 3 can be detected from HAADF STEM images [2], and the number of atoms in an atomic column can be counted with single atom sensitivity [3,4] (see fig. 1). The question then arises: how far can we go?

Ultimately, the attainable precision with which unknown structure parameters can be estimated is set by the unavoidable presence of electron counting noise. For continuous parameters such as atom positions, this limit is expressed by means of the Cramér–Rao lower bound (CRLB). For discrete parameters such as the number of atoms, the probability of error can be derived (defining e.g. the probability of reporting an atom when there is none or reporting no atom when there is one) [5].

Using these expressions, we show that the precision of the 3D atom positions estimated from depth sectioning data is poor under realistic exposure times. However, when simplifying the problem to the estimation of the vertical position of each atomic column with known number of atoms, picometer precision can be reached. The performance of depth sectioning can now be compared with HAADF STEM tomography. Furthermore, evaluating the probability of error helps us to determine STEM detector settings resulting into the highest detectability of light atoms (see fig. 2). Even so, we can define the minimally required electron dose in order to attain a maximum allowable error for miscounting atoms. This is of great importance when studying beam sensitive structures.

In conclusion, statistical parameter estimation theory is used to explore fundamental limits with which structure parameters can be estimated. The CRLB and the probability of error not only outperform classical performance criteria (including resolution, contrast or SNR), they also allow us to predict attainable limits under given experimental conditions and to explore the optimal experimental settings.

[1] S. Van Aert et al., Advanced Materials 24 (2012), p.523

[2] S. Van Aert et al., Ultramicroscopy 109 (2009), p.1236

[3] S. Van Aert et al., Nature 470 (2011), p.374

[4] S. Van Aert et al., Physical Review B 87 (2013), 064107

[5] A.J. den Dekker et al., Ultramicroscopy 134 (2013), p.34


The authors kindly acknowledge funding from the Fund for Scientific Research, Flanders (FWO).

Fig. 1: Examples of quantitative analysis using statistical parameter estimation theory. (a)Measurement of Ti displacements of ca. 5 pm inside a twin boundary in CaTiO3 [1]. (b)Quantitative characterisation of a La0.7Sr0.3MnO3-SrTiO3 interface from an HAADF STEM image [2]. Counting results of Ag (c) and Au atoms (d) with single atom sensitivity [3,4].

Fig. 2: Probability of error as a function of the inner radius of an annular detector for an aberration corrected microscope (300kV acceleration voltage, 21.7mrad probe forming angle, 100mrad outer detector radius, 12000 incident electrons/Å2). Green, red, and blue correspond to the problem of detecting Li, H, and Al/Ti in LiV2O4, YH2, and SrTiO3/LaAlO3.

Type of presentation: Poster

IT-16-P-1553 Rutherford scattering of electron vortices

Van Boxem R.1, Partoens B.2, Verbeeck J.1
1EMAT, University of Antwerp, Belgium, 2CMT, University of Antwerp, Belgium
ruben.vanboxem@uantwerpen.be

Vortex beams have been met with great interest in various fields like optics, telecommunication, acoustics, and more recently in electron microscopy. Recently, techniques for the manipulation of the electron wave’s phase have received a boost1,2. Electron vortex beams show promise as a new tool to detect material properties at the nanoscale in a novel way. One feature which might contribute is the additional magnetic moment induced by a vortex electron’s orbital angular momentum (OAM).

Theoretically, various studies have been done describing free space electron vortices3 and how they behave in electromagnetic fields4,5. Relativistic aspects and the electron spin coupling to the OAM have also been considered6,7. On the other hand, basic scattering theory with electron vortex beams has not yet been fully understood, and that is why elastic scattering of an electron vortex beam on a screened Coulomb potential is considered here. This work8 introduces the incoming beam’s OAM (and associated transverse momentum) into the first Born approximation of quantum scattering theory. The influence of a beam’s OAM and corresponding spatial shape on the elastic scattering amplitude has been analyzed using the derived analytical formula.

Using the results here, we can propose scattering experiments in which high values of transverse momentum of the initial beam expose these novel features, proving the treatment here lives up to its intent: generalize plane wave scattering to cylindrically symmetric beams, including those with OAM. With this result established, more complicated scattering amplitudes can be calculated, leading to a complete electron vortex scattering theory.

1 A. Béché, R. Van Boxem, G. Van Tendeloo, and J. Verbeeck, Nature Physics, vol. 10, pp. 26–29, Jan 2014.
2 L. Clark, A. Béché, G. Guzzinati, A. Lubk, M. Mazilu, R. Van Boxem, and J. Verbeeck, Phys. Rev. Lett., vol. 111, p. 064801, Aug 2013.
3 P. Schattschneider and J. Verbeeck, Ultramicroscopy, vol. 111, no. 9-10, pp. 1461–1468, 2011.
4 K. Y. Bliokh, P. Schattschneider, J. Verbeeck, and F. Nori, Phys. Rev. X, vol. 2, p. 041011, Nov 2012.
5 K. Y. Bliokh, Y. P. Bliokh, S. Savel’ev, and F. Nori, Phys. Rev. Lett., vol. 99, p. 190404, Nov 2007.
6 K. Y. Bliokh, M. R. Dennis, and F. Nori, Phys. Rev. Lett., vol. 107, p. 174802, Oct 2011.
7 R. V. Boxem, J. Verbeeck, and B. Partoens, EPL (Europhysics Letters), vol. 102, no. 4, p. 40010, 2013.
8 R. V. Boxem, J. Verbeeck, and B. Partoens, to be published in PRA.


RVB acknowledges support from an FWO PhD fellowship grant (Aspirant Fonds Wetenschappelijk Onderzoek Vlaanderen).

JV acknowledges support from the ERC Starting Grant 278510 VORTEX.

Fig. 1: Schematic of the convergent beam scattering experiment. The relation of the transverse momenta and the aperture dimensions is shown, and an on-axis pinhole detector is shown.

Fig. 2: Transverse wave functions for the Bessel, aperture far field, and Laguerre-Gaussian beams. They all contain a first order vortex. The Laguerre-Gaussian has two intensity lobes (n=2), clearly showing the strongest localization of the three in the transverse plane.

Fig. 3: Elastic Coulomb scattering amplitude for fixed transverse momentum, and several values of OAM.

Fig. 4: Zeroeth order elastic scattering amplitude for several values of the transverse momentum. The limit to the plane wave result is clearly visible.

Type of presentation: Poster

IT-16-P-1689 A new approach to determine excess free volume at high-angle grain boundaries – a proof of concept

Buranova Y. S.1, Rösner H.1, Divinski S. V.1, Wilde G.1
1Institute of Materials Physics, University of Münster, Wilhelm-Klemm-Str. 10, D-48149 Münster, Germany
buranova@uni-muenster.de

Grain boundaries (GBs) have a significant impact on the physical, especially mechanical properties of polycrystals. The GBs are typically characterized by structure units which are different compared to crystalline unit cells. It is assumed that grain boundary excess free volume plays an important role since it can be related to the grain boundary energy and has a significant influence on the transport and thermodynamic properties (diffusion/segregation). It has been reported that different GBs exhibit rather different excess free volumes, i.e. different specific mass densities from the crystalline bulk [1,2].
In this study we describe a new approach to determine the excess free volume from high-resolution transmission electron microscopy (HRTEM) images. For this purpose, an image analysis tool has been designed that allows determination of the local mass density from HRTEM images. Thereto the intensities of the GB regions have been compared with that of the grain interiors and the difference identified to be proportional to the density change, Δρ =(I-Igb)/I, where I is the intensity of the grain interior and Igb is the intensity of the GB.
In order to prove this concept, symmetrical tilt GBs with zone axes along the [100], [110] and [111] directions have been generated using molecular dynamics simulation (applying the LAMMPS software [3]) and subsequently taken as input for the simulation of HRTEM images using the Kirkland code [4,5]. The maximal density change of the GB has been estimated to be around -6% for the analyzed GBs. Calculations show that this approach works for pure samples with thicknesses up to 15 nm including aluminum oxide layers. The reliability of this approach is evaluated for different artificially chosen configurations including chains of vacancies and solute atoms.
Experimentally, well-defined aluminum bi-crystals have been investigated using aberration-corrected HRTEM. The results are discussed with respect to the relation between local structure, excess free volume and specific excess energy density.
[1] HB Aaron, GF Bolling, Surf. Sci. 31 (1972) p. 27.
[2] D Wolf, Scripta Metall. 23 (1989) p. 1913.
[3] S Plimpton, J Comp Phys. 117 (1995) p.1
[4] DL Olmsted, SM Foiles, EA Holm Acta Mater. 57 (2009) p. 3694.
[5] EJ Kirkland: Advanced computing in electron microscopy. 2nd ed. New York, Springer (2010).


Fig. 1: Example of a symmetrical [100] tilt GB (left) and the work flow of the approach (right).

Fig. 2: Characteristic plot of the intensity across a symmetrical [100] tilt GB shown in Fig.1.

Fig. 3: Intensity change versus misorientation angle of simulated symmetrical [100] tilt GBs. Solid lines represent least squares fits to the data points.

Type of presentation: Poster

IT-16-P-1767 A Multislice Theory of Electron Scattering in Crystals including Backscattering and Inelastic Effects

Spiegelberg J.1, Rusz J.1
1Department of Physics and Astronomy, Uppsala University, Box 516, S-751 20 Uppsala, Sweden
jakob-spiegelberg@gmx.de

In order to interpret diffraction patterns obtained in transimission electron microscopy, scattering processes in the crystal have to be understood theoretically. Among all existent approaches to describe electron scattering, the multislice method has proven to be among the most versatile ones. To the best of our knowledge, existing multislice formalisms can not treat backscattering and inelastic processes simultaneously, what is of particular importance in electron microscopes working at reduced acceleration voltages. Furthermore, if the aim is to describe experiments using inclined illumination, e.g. reflection high-energy electron diffraction (RHEED), the conventional multislice method is inapplicable.

By combining Yoshioka's theory of inelastic scattering [H. Yoshioka, J. Phys. Soc. Jpn. 12, 6 (1957)] and van Dyck's approach for backscattering [J. H. Chen and D. Van Dyck, Ultramicroscopy 70, 29-44 (1997)], a general multislice formalism incorporating these phenomena can be derived. The new method is based on the slice transition operator technique defining an operator S(j-1) linking the wavefunctions of two adjacent slices, the jth and (j-1)th slice. In this case, S(j-1) is a 2n dimensional matrix where n is the number of inelastic excitations considered. Due to the complexity of the elements of S(j-1), however, a self-consistent solution of the matrix equation

          Φ(j)S(j-1) · Φ(j-1)

is computationally very costly. Making the single inelastic scattering approximation, we propose a computational scheme allowing control over the number of backscattering events considered. In a nutshell, one identifies the contributions to forward and backward scattering in S(j-1) and propagates the wavefunction forward using the corresponding operators from S(j-1). During the propagation, contributions to the backward scattered beam arise - the backward propagation can be started using the corresponding parts of S(j-1). Since the backward propagation creates contributions to the twice backscattered, forward traveling beam, this cycle can be repeated until convergence is reached.

We present a more detailed description of the computational scheme and the derivation of the slice transition operator matrix. Moreover, we compare predicted elastic diffraction patterns of our new approach with those predicted by conventional multislice [J. M. Cowley and A. F. Moodie, Acta Cryst. 10, 609 (1957)] and the recent development [C. Cai and J. Cheng, Micron 43, 374 (2012)]. Special interest is taken into the single backscattering approximation and higher order backscattering.


Fig. 1: Left Column: Wavefunction of the electron beam after passing a 10 nm thick sample of bcc iron as predicted by the conventional multislice method and the differences between the CMS and RSMS wavefunction or the wavefunction predicted by the new formalism presented here. Right Column: Corresponding Fourier transforms.

Type of presentation: Poster

IT-16-P-1962 Slice-by-slice simulations of absorption potential for high-angular resolution electron channeled X-ray spectroscopy

Ohtsuka M.1, Muto S.2
1Graduate School of Engineering, Nagoya University, Nagoya 464-8603, Japan, 2EcoTopia Science Institute, Nagoya University, Nagoya 464-8603, Japan
m-ohtsuka@nucl.nagoya-u.ac.jp

Characteristic X-ray spectra vary with the change in the symmetries of the Bloch waves propagating preferentially along particular atomic sites, depending on the incident beam direction. The technique, high-angular resolution electron channeled X-ray spectroscopy, is accessible to atomic site-selective elemental analysis by beam rocking (i.e. scanning reciprocal space), as a good counterpart of the atomic column-by-column analysis using the state-of-the-art scanning transmission electron microscopy (i.e. scanning real space).

The incoherent channeling pattern (ICP), a 2D intensity distribution obtained by a beam-rocking EDX technique, can be quantitatively analyzed by comparing with the inelastic scattering cross-section using e.g., the Bloch-wave method [1], which is, however, time-consuming for calculating 2D ICPs, particularly in calculating the cases where the sample contains an extended defects such as a surface, interface or planar defect etc. We have thus developed a computationally more efficient algorithm where two kinds of absorption potentials are introduced [2]: the absorption potential Uall incorporating the inelastic contributions such as phonons, plasmons and core excitations, and Upartial = Uall − UEDX, where UEDX is the core excitation of a particular atom to calculate its subsequent X-ray emission. The wavefunctions Ψall and Ψpartial are calculated by solving the Schrodinger equations with Uall and Upartial, respectively. The difference in the total intensities between these two wavefunctions should be related to the characteristic X-ray intensity. However, Ψpartial tends to be overestimated for larger thickness and the attenuation of the electron densities propagating through the particular atoms is underestimated. In order to avoid this problem, we introduce a slice-by-slice method which divides the specimen into many thin slices. The X-ray intensity is evaluated by taking the slice-by-slice difference between the corresponding wave functions. The wavefunctions are connected at the entrance surface of each slice.

Figure 1 shows the calculated thickness dependence of the relative Sr-L line intensities of SrTiO3 at the exact [001] incident. Figure 2 shows calculated ICPs, tilted around the [001] axis. These results show the calculated ICP with the sufficient number of slices is nearly identical to that of the conventional method. The present scheme is easily extended to the multislice method which is particularly suitable for the cases where the system of interest contains lattice defects.

References
[1] M. P. Oxley, and L. J. Allen, J. Appl. Cryst., 36, 940 (2003).
[2] K. Watanabe et al., Phys. Rev. B, 63, 085315 (2001).


A part of this work was supported by a Grant-in-Aid on Innovative Areas "Nano Informatics" (Grant number 25106004) from the Japan Society of the Promotion of Science.

Fig. 1: Thickness dependence of Sr-L line intensities of SrTiO3 at [001] incident, calculated by the present method with single slice (black line), and many thin slices by dividing 150 nm thick into 769 slices (red line), and conventional inelastic cross-section method (open circles).

Fig. 2: Two-dimensional Sr-L ICPs around [001] of 150 nm thick SrTiO3, calculated by the conventional inelastic cross-section method (a), and present method with single slice (b), and 769 slices (c), respectively.

Type of presentation: Poster

IT-16-P-2001 Modified Random Walk Algorithm for Monte Carlo Modeling of EBIC and Cathodoluminescence

Priesol J.1, Šatka A.1
1Institute of Electronics and Photonics, Slovak University of Technology, Ilkovičova 3, 812 19 Bratislava, Slovakia
juraj.priesol@stuba.sk

The aim of this contribution is to discuss the results obtained by Monte Carlo (MC) modeling and simulation of the interaction of accelerated primary electrons with semiconductors and semiconductor structures. In electron microscopy practice, MC method is routinely used for quantitative assessment of spatial distribution of energy deposited into the semiconducting material by primary electrons [1], [2], but there is a lack of papers dealing with the MC simulations of consequent diffusion and recombination processes of generated charge carriers [3], [4]. Diffusion of minority carriers has a radical impact on diffusion sensitive methods like electron beam induced current (EBIC) and cathodoluminescence (CL) and therefore it is important to pay a proper attention to it. Due to the complexity of MC simulations, diffusion is usually considered as a random motion of particles according to the random walk algorithm [5], i.e. each generated carrier passes constant distance Δs in random direction constant number of times k. Based on this simple model, three dimensional MC simulations of random diffusion from point source with initial number of N0 generated carriers were executed. MC simulations reveal non-exponential decrease of the carriers from the point source, which is not in agreement with analytical approximation and it has its origin in erroneous assumption of equal lifetime for each simulated carrier. It has been found out that this discrepancy has only a little effect on CL accuracy whereas it is significant for simulation of EBIC line profiles. To overcome this, a modification of random walk algorithm was proposed, where the value of k was determined using probability density function according to normal statistical distribution. The application of adapted model and its influence on the results of MC simulations of EBIC (Fig. 1) and CL, as well as the comparison of simulation and experiment (Fig. 2) performed on III-N semiconductor structures will be presented and discussed.

 

References

[1] Joy, D. C.: Monte Carlo modeling for electron microscopy and microanalysis. Oxford University Press, Inc., 1995, 224 pp., ISBN: 0-19-508874-3.
[2] Demers, H. et al, Scanning 33 (2011), p.135–146.
[3] Ledra, M. - Tabet, N., Superlattices and Microstructures 45 (2009), p.444–450.
[4] Doan, Q. T. - El Hdiy, A. - Troyon, M., J. Appl. Phys. 110 (2011), p. 124515.
[5] Pearson, K., Nature, no. 1865, vol. 72, (1905) p. 294.


This work has been supported by the Slovak Research and Development Agency (contract No. APVV-0367-11) and by Slovak Grant Agency (project VEGA No. 1/0921/13).

Fig. 1: Spatial distributions of 105 carriers recombining in GaN sample around the circular Shottky contact with radius rc = 2500 nm (left) and 150 nm (right) and  simulated for electron beam energy Epe = 5 keV, diffusion length of minority charge carriers L = 196 nm and surface recombination velocity vs = 0.

Fig. 2: Simulated EBIC line profiles at the circular Schottky contact with radius rc for different diffusion lengths L of minority carriers at Epe = 5 keV (left), and EBIC line profiles extracted from EBIC maps measured at reverse polarized Schottky contact formed by tungsten needle set onto the undoped GaN layer (right); fit shows diffusion length ~260nm.

Type of presentation: Poster

IT-16-P-2085 Calculations of elastic and inelastic scattering processes of relativistic electrons in oriented crystals

Hinderks D.1, Kohl H.1
1Physikalisches Institut der Universität Münster, Münster, Germany
dieter.hinderks@uni-muenster.de

Many modern electron microscopes operate at acceleration voltages of several hundred kV. The accelerated electrons thus reach velocities approaching the speed of light. Therefore the scattering processes have to be treated relativistically. We focus on inelastic scattering in crystals.

In a non-relativistic treatment the movement of the electrons inside the crystal is described using Bloch waves. Before the electrons enter into the crystal they are described by simple plane waves. This view is used in non-relativistic calculations in many cases. The periodic potential of a crystal provides Bloch waves as solutions of the Schrödinger equation. To ensure the boundary conditions at the interface of crystal and vacuum, the transmitted electrons are described using a superposition of plane waves. To obtain a reliable result for the scattering process, many excited Bloch waves have to be considered. The scattering process is mathematically described using matrix elements [1]. The computational complexity depends strongly on the number of Bloch waves considered. The general solution for the wave function of the incident electrons in the crystal is a superposition of many Bloch waves, which are excited at the same time. The excitation of every single Bloch wave is weighted with a excitation coefficient. The number of excited Bloch waves which have to be taken into account depends on the geometry of the crystal.

In this work we focus on an extension of this treatment for relativistic electrons. In contrast to the non-relativistic case the wave functions of the fast incident electrons and the atomic electrons have to be calculated using the Dirac equation. Therefore the incident electrons are described by relativistic four-component Bloch waves (Fig. 1). In our approach we use the relativistic propagator theory where the atomic electrons are seen under influence of a scalar and a vector potential generated by the fast incident electrons via their charge and current (Fig. 2). Furthermore retardation is considered in this relativistic treatment. This approach has previously been used for relativistic plane waves [2]. To consider crystalline materials the incident electrons are described by the relativistic Bloch waves. Consequently the matrix elements contain different sums over reciprocal space and the different single relativistic Bloch waves. The fourier coefficients of these Bloch waves depend on the crystal structure and can be calculated analoguously to the non-relatvistic treatment.

[1] A. Weickenmeier and H. Kohl, Phil. Mag. B60 (1989) 467.

[2] R. Knippelmeyer et al.,Ultramicroscopy 68 (1997) 25-41.


Fig. 1: Relativtistic Bloch wave

Fig. 2: Scattering Matrix

Type of presentation: Poster

IT-16-P-2124 The dependence of SNR, contrast and resolution on electron dose and sampling

Lee Z.1, Rose H.1, Lehtinen O.1, Biskupek J.1, Kaiser U.1
1Universität Ulm, Materialwissenschaftliche Elektronenmikroskopie
zhongbo.lee@uni-ulm.de

Due to the practical application of hardware aberration correctors, the instrumental resolution of transmission electron microscopes has been remarkably improved. In order to achieve the highest resolution in aberration-corrected (AC) high-resolution transmission electron microscopy (HRTEM) images, high electron doses are required. In the case of high accelerating voltages, materials can be damaged predominantly via the knock-on damage mechanism, where atoms are displaced by direct impacts of the energetic incident electrons. However, when reducing the accelerating voltage, ionization can become the dominating damage mechanism, as the inelastic scattering cross section increases [1]. Effective ways of reducing ionization damage may be cooling of the specimen [2], or conductive coating [3]. However, such approaches are not always feasible. In both, high and low accelerating voltages, images need to be acquired with limited electron doses.

In this work we have performed dose-dependent AC-HRTEM image calculations (Fig. 1), and the dose related noise is treated as stochastic fluctuations around the ideal electron counts on each image pixel, instead of the additive noise. We have studied the dependence of the signal-to-noise ratio (SNR), atom contrast and resolution on electron dose and sampling. Graphene is used as the example material due to the simplicity of its structure, as it is the thinnest and lowest Z-number crystalline material, which allows most straight forward interpretation of the results. We have introduced a dose-dependent contrast definition, which can be used to evaluate the visibility of objects under different dose conditions. Based on our calculations, we have determined optimum samplings for high and for low electron dose imaging conditions.

Our calculation shows: SNR, atom contrast and resolution, all improve with increasing electron dose, converging towards their values obtained at infinite dose. As the sampling increases, the SNR increases and the resolution decreases; the atom contrast improves as long as the damping of MTF is negligible. We have determined optimum sampling under high-dose and low-dose conditions. Under high-dose conditions, the optimum sampling depends mainly on the required specimen resolution; under low-dose conditions, the best sampling is determined by our criteria that the required specimen resolution should be achieved with the minimal electron dose.

[1] R. Egerton, Microsc. Res Tech. 75(2012)1550-1556.

[2] Y. Chen and W. Sibley, Physical Review, 154(1967) 842.

[3] L. Reimer and H. Kohl, Transmission Electron Microscopy: Physics of Imaging Formation. Imperial College Pr. 2008.

[4] Z. Lee et al., Ultramicroscopy (2014), http://dx.doi.org/10.1016/j.ultramic.2014.01.010.


This work was supported by the DFG (German Research Foundation) and the Ministry of Science, Research and the Arts (MWK) of Baden-Württemberg in the frame of the (Sub-Angstrom Low-Voltage Electron microscopy) (SALVE) project.

Fig. 1: Calculated HRTEM images of graphene for different doses and samplings with a usable aperture of 50 mrad under 80 kV. The last row shows the CTF (purple) for different samplings. The PCTF function (blue), focus spread envelope (red) and image spread envelope (yellow) are the same for each column. Reproduced from the reference [4].

Type of presentation: Poster

IT-16-P-2225 Inelastic scattering of electron vortex beams: mechanism and optimal conditions for EMCD measurements

Rusz J.1, Bhowmick S.2
1Department of Physics and Astronomy, Uppsala University, Uppsala, Sweden, 2Department of Materials Science and Engineering, Indian Institute of Technology, Kanpur, India
jan.rusz@physics.uu.se

Electron vortex beams (EVBs) have attracted a lot of attention since their introduction to the transmission electron microscopy [1,2,3]. In [2] it was reported that EVBs should allow measurement of electron magnetic circular dichroism (EMCD; [4]). Our recent simulations of scanning transmission electron microscopy (STEM) seem to rule out utility of EVBs for measurement of EMCD, unless performed in atomic resolution [5].

The distribution of the EMCD signal in an intrinsic EMCD experiment [4] is anti-symmetric with respect to the mirror symmetry axes. As a consequence, a detector centred on a transmitted beam will not detect any net EMCD signal due to cancellation of positive and negative contributions. In contrast, vortex-induced EMCD can be measured at the transmitted beam. At the level of scalar-relativistic theory and assuming dipole-allowed transitions, we show that this only happens, when the discs in convergent beam electron diffraction (CBED) pattern overlap. Simulations in Fig. 1 illustrate the principle. The top row shows real-space probe wavefunction after passing through 10nm slab of bcc iron oriented in (001) zone axis. The beam diameter (FWHM is indicated in the top) is gradually reduced from left to right. Corresponding elastic CBED pattern shows discs, which eventually start overlapping, as the beam diameter is reduced. In the third row there is energy-filtered Fe-L3 diffraction pattern, which acquires chirality once the CBED discs start overlapping. Finally, bottom row is the distribution of EMCD signal in the diffraction plane. Note the symmetric component of EMCD developing in the middle, once CBED discs start overlapping.

There are several parameters, which influence the inelastic scattering of EVBs: acceleration voltage, beam diameter, EVB angular momentum and a distance of the beam from an atomic column. Optimum may vary as a function of crystal thickness, structure and orientation. We fixed the latter two to (001) zone axis of bcc iron. All the other parameters were independently varied. The results of the optimization are summarized in a condensed form in Fig. 2. The best conditions are predicted for 10nm slab of bcc iron using an EVB of FWHM diameter 1.6Å with an angular momentum 1ħ at acceleration voltage 200keV, using an annular detector of inner (outer) diameters of 1G (5G), respectively, G=(100). These values appear to be reachable with state-of-the-art STEM instruments available today [6].

[1] M. Uchida and A. Tonomura, Nature 464 (2010), 737.
[2] J. Verbeeck, H. Tian, and P. Schattschneider, Nature 467 (2010), 301.
[3] B. J. McMorran et al., Science 331 (2011), 192.
[4] P. Schattschneider et al., Nature 441 (2006), 486.
[5] J. Rusz and S. Bhowmick, Phys. Rev. Lett. 111 (2013), 105504.
[6] O. Krivanek et al., submitted.


We acknowledge Swedish Research Council and Swedish National Infrastructure for Computing.

Fig. 1: Scattering of EVB with angular momentum 1ħ on a 10nm thick bcc Fe crystal oriented along (001) zone axis. Beam diameter is shown by a violet label. Top rows shows scattered EVB wavefunction, elastic diffraction pattern (2nd row), energy filtered diffraction pattern at Fe-L3 edge (3rd row) and the EMCD distribution in the diffraction plane (bottom).

Fig. 2: Optimization of vortex-beam EMCD as a function of acceleration voltage Vacc and angular momentum <Lz>. Optimal beam diameter and inner aperture diameter Rin are shown in 4th and 3rd row. Resulting absolute and relative strength of EMCD are shown in the 1st and 2nd row, respectively. The overall optimum is the white square in the top left panel.

Type of presentation: Poster

IT-16-P-2289 An accurate parameterization for the scattering factors for neutral atoms that obey all physical constraints

Lobato I.1, Van Dyck D.1
1Emat, University of Antwerp , Groenenborgerlaan 171, B-2020, Department of Physics, Antwerp, Belgium
Ivan.Lobato@uantwerpen.be

In electron microscopy and electron diffraction the high energy electrons interact with the atoms of the sample trough their electrostatic Coulomb potential. This interaction is described by the high-energy Schrödinger equation. The simplest approach to describe the electrostatic specimen potential is by linear superposition of the spherically symmetric electrostatic potentials of each atom in the specimen. The atomic potential is related to the atomic charge distribution via Poisson's equation. The electron charge distribution can be computed from the knowledge of the atomic wave function, which can be obtained by numerically solving the Dirac equation. From the electron charge distribution, one can then calculate the X-ray scattering factor. By accounting for the contribution of the nucleus to X-ray scattering factor using the Mott-Bethe formula, one can then calculate the electron scattering factor.

In the simulation programs, one can use the numerical values for the scattering factors and interpolate them to obtain the atomic potential or scattering factor at the required points. But in order to reduce the data, one parameterizes the scattering factors using basic functions such as Gaussians and Lorentzians [1-3]. The drawback of all these parameterizations is that they are obtained by fitting with a discrete set of numerical values from which the asymptotic behavior is then extrapolated.

However, it becomes increasingly clear that electron scattering and simulation programs should also include scattering at very large angles. Indeed, high-angle annular dark-field (HAADF) STEM calculations are done using a frozen lattice model in which the atoms remain sharp and the scattering factors are not dampened by a Debye-Waller factor. The same argument holds for accurate computations of scattering into Higher Order Laue Zones (HOLZ). This means that reliable quantitative simulations for TEM or STEM can only be done using an accurate parameterization for the electron scattering factors that are valid up to very large angles.

We developed a new parameterization of the electron scattering factor using the analytic non-relativistic hydrogen electron scattering factors as a basis functions. The inclusion of the correct physical constraint in the electron scattering factor and its derived quantities allow using the new parameterization in different fields. A comparison between the new parameterization [4] and the previous analytical fittings are shown through figures 1-3.
References
1. A. Weickenmeier, H. Kohl, Acta Cryst. A24(1991), 390.
2. L.M. Peng, G. Ren, S.L. Duvared, J. Whelan, Acta Cryst. A52(1996), 257.
3. E.J. Kirkland, Advanced Computing in Electron Microscopy, Plenum Press, New York, 1998.
4. I. Lobato, D. Van Dyck to be published.


Fig. 1: The root mean square values of the deviation e between the numerical and fitted electron scattering factors vs. atomic number Z using four different parameterizations.

Fig. 2: Comparison among different fittings of the electron and X-ray scattering factors vs. scattering angle for copper. The green markers are the tabulated values of the scattering factors of Kirkland.

Fig. 3: Comparison among different calculated atomic potentials and electron densities vs. the three dimensional radius for copper.

Type of presentation: Poster

IT-16-P-2511 Recent improvements in the STEM-CELL software

Grillo V.1,2, Rotunno E.2, Campanini M.2, Spadaro M. C.2,3, D'addato S.2,3
1CNR-Istituto Nanoscienze, Centro S3, Via G Campi 213/a, I-41125 Modena, Italy, 2CNR-IMEM, Parco delle Scienze 37a, I-43100 Parma, Italy, 3Università di Modena e Reggio Emilia, via G Campi 213/a, I-41125 Modena, Italy
vincenzo.grillo@cnr.it

TEM and STEM analytical studies are going in toward a stronger involvement of computing for the image interpretation. Geometric Phase Analysis [1], for example, has proved to be a useful tool to evaluate the strain in different structures. TEM Image simulations are another very important tool to be compared with complicated structures to obtain quantitative interpretation of the contrast. Moreover the simulation of STEM images is a fundamental step to perform quantitative HAADF measurements [2]. Finally phase retrieval methods in TEM and probe deconvolution in STEM are useful tools to improve the image information [3]. Many simulation software are already available, however it is difficult to find a free and graphical tool that permits to perform both simulation and analysis on the same platform. For this reason we created the STEM CELL project [3][4]. The proposal of STEM CELL is

1) to facilitate multislice simulations by creating, manipulating complicated cells, facilitating the selection of the simulation parameters and interface with simulation routines (the work was based on Kirkland routines [5])

2) to implement analysis methods on both simulated and experimental images so that simulation can be more directly used as a benchmark for experiments.

3) To implement new simulation/analysis methods.

Among the most recent new features it is worth mentioning the probe deconvolution in STEM HAADF images, the phase reconstruction by means of the transport of intensity equation, the simulation of diffraction patterns for any unit cells and the column by column quantitative analysis of the HAADF contrast.

We show in fig. 1 a simulation of two GaAs tetrapods and a Ni particle. A layer of amorphous carbon has been also added to provide a more realistic effect. The sample has been constructed within STEM_CELL and simulated using the embedded Kirkland’s software. Fig. 2 is an example of experimental analysis of the contrast of a CeO2 nanoparticle. The reported experimental image has been obtained by deconvolution of the original image (not shown). The contrast of each column is interpreted approximately in terms of thickness (i.e. number of Ce atoms per column) using a “quick” calibration based on the analysis of the intensity histogram and of the minimal intensity step. Fig. 3 is an example of the transport of intensity equation application on magnetic particles [6].

[1] F M Hytch, et al Ultramicroscopy 74, (1998) 131.

[2] E. Carlino et al. Physical Review B 71 (2005) 235303.

[3] V. Grillo et al. Ultramicroscopy 125 (2013) 112

[4] V. Grillo et al. Ultramicroscopy 125 (2013)97

[5] E.J. Kirkland, Advanced Computing in Electron Microscopy, Plenum Press, NY 1998.
[6] V.V. Volkov et al. Ultramicroscopy 98 (2004) 271


Fig. 1: Model and HREM Simulation for 2 GaAs tetrapods and a Ni particle

Fig. 2: a) Deconvoluted STEM image of a CeO2 nanoparticles and b) quantitative column by column analysis of intensity

Fig. 3: Example of TIE (transport of intensity equation) analysis and graphical representation of magnetic field of two different groups of magnetic nanoparticles. The hue and brightness refer to different B direction/modulus. The line of the field are also shown.

Type of presentation: Poster

IT-16-P-2612 A Perturbation Theory Study of Electron Vortices in Electromagnetic Fields: the Case of Infinitely Long Line Charge and Magnetic Dipole

Xie L.1, Wang P.1, Pan Q. X.1
1National Laboratory of Solid State Microstructures and College of Engineering and Applied sciences, Nanjing University, Nanjing 210093, People’s Republic of China
wangpeng@nju.edu.cn

Electron vortex beam with a quantized orbital angular momentum l and half-integer spin angular momentum has been intensively studied since its introduction intro transmission electron microscopy (TEM) and scanning-TEM because it leads to appreciable number of potential applications in electron microscopy and nanomanipulation [1-3]. Until now, the fundamental physics of electron vortices is still in its infancy and only a few work concerning the basic interactions between electron vortex beams and matter were reported [4][5]. In this work, we study three fundamental interactions: the electron-electric potential interaction, the electron-magnetic potential interaction and the spin-orbit-coupling (SOC) of electron vortex beams in electromagnetic fields based on the relativistically corrected Pauli-Schrӧdinger equation and the perturbation theory.
We first study the interactions between the vortex beam and an infinitely long line charge with a line charge density δE. According to our calculation with an accelerating voltage of 300 kV and a convergence angle of 20 mrad electron vortex beam, both the electron-electric potential and SOC interactions are proportional to δE and decrease monotonically as a function of the topological charge l, as shown in Fig 1. Compared with the electron-electric potential interaction, which is several eV, the SOC interactions are much weaker and are in the range of 10-5~10-3 eV. Next we investigate the interactions between the vortex beam with an infinitely long magnetic dipole with a magnetic moment density δM polarized along z-direction. Fig. 1 shows the electron-magnetic interaction is also in the range of 10-5~10-3 eV. Our calculations indicate that the SOC and electron-magnetic potential interactions are too weak to be observed practically. Nevertheless, it is theoretically predicted that these weak interactions can be raised if a large convergence angle is used. In Fig. 2, we show the calculated results with an accelerating voltage of 300 kV and a convergence angle of 100 mrad electron vortex beam and it is evidently that both the SOC and electron-magnetic potential interactions are dramatically increased by one order, which is favorable for the observation of their effect in future aberration-corrected electron microscopy.

References:
[1] Bliokh, K. Y., et al, Phys. Rev. Lett. 99, (2007), 190404.
[2] Verbeeck, J., et al, Nature (London) 467, (2010), 301-304.
[3] McMorran, B. J., et al, Science 331, (2011), 192-195.
[4] Bliokh, K. Y., et al, Phys. Rev. Lett. 107, (2011),174802.
[5] Lloyd, S. M., et al, Phys. Rev. Lett. 109, (2011), 254801.
[6] Xie, L., et al, Micron, Accepted, (2014).


The authors would like to thank 1000 young talent plan program of China..

Fig. 1: Eel, Esoc and Emag as a function of topological charge l for 300 kV, 20 mrad electron vortex beam. Note that Esoc and Emag are ×10000 scaled.

Fig. 2: Eel, Esoc and Emag as a function of topological charge l for 300 kV, 100 mrad electron vortex beam. Note that Esoc and Emag are ×10000 scaled.

Type of presentation: Poster

IT-16-P-2621 Realistic amorphous carbon model for high resolution microscopy and electron diffraction simulations

RICOLLEAU C.1, LE BOUAR Y.2, AMARA H.2, LANDON-CARDINAL O.2, ALLOYEAY D.1
1Laboratoire Matériaux et Phénomènes Quantiques, CNRS-UMR 7162, Université Paris Diderot-Paris 7, Case 7021, 75205 Paris Cedex 13, France, 2Laboratoire d’Etude des Microstructures, UMR CNRS/ONERA 29, avenue de la Division Leclerc, 92322 Châtillon, France
christian.ricolleau@univ-paris-diderot.fr

Amorphous carbon and amorphous materials in general are of particular importance for high resolution electron microscopy, either for bulk materials, generally covered with an amorphous layer when prepared by ion milling techniques, or for nanoscale objects deposited on amorphous substrates. In order to quantify the information of the high resolution images at the atomic scale, a structural modeling of the sample is necessary prior to the calculation of the electron wave function propagation. It is thus essential to be able to reproduce the carbon structure as close as possible to the real one. The approach we propose here is to simulate a realistic carbon from an energetic model based on the tight-binding approximation in order to reproduce the important structural properties of amorphous carbon.
In this work, we propose a new method to model in a more realistic way amorphous carbon (a-C) that accurately accounts for its 3D structure. It is based on an energetic approach with a tight-binding (TB) potential in which the electronic band structure of the material is calculated with the recursion technique. The main advantage of this model is that it gives a very good description of the sp, sp2, and sp3 hybrid bonds and their competition [1].
At first, the model and the main structural properties of the generated carbon will be presented and compared with a simple model of carbon, where the atom positions are generated randomly. We have shown that the limit thickness for the wave phase approximation if 30% overestimated if we consider the random carbon model (Fig. 1). In a second step, we have studied the influence of the carbon model on the contrast of single Cu, Ag and Au atoms deposited on amorphous carbon substrate. Our work does not indicate any significant influence of the carbon structure on single-atom contrast when statistically relevant measurements are performed. Finally, we have compared both model structures for the determination of the long-range order parameter in a small CoPt nanoparticles deposited on a-C layer. We have clearly shown the importance to use realistic amorphous carbon model to obtain quantitative values for the diffracted intensities. As it can be observed on Figure 2, diffuse scattering intensity due to the 3D atomic arrangements of the realistic carbon has a significant contribution to the scattered information especially at low spatial frequencies [1].
This work emphasizes the necessity to use realistic carbon model for TEM image and diffraction simulation in order to extract very sensitive quantitative information, particularly in diffraction experiments.

[1] Ricolleau C., Le Bouar Y., Amara H., Landon-Cardinal O. and Alloyeau D., J. Appl. Phys., 114, 213504 (2013).


*Currently address: Département de Physique, Université de Sherbrooke, Sherbrooke, Québec J1K 2R1, Canada.

Fig. 1: Comparison of the (a) phase and (b) intensity of the transmitted beam (i.e., the unscattered part of the electron wave) calculated using Random and tight-binding carbon structures as a function of the thickness layer. The slice thickness is 0.25 nm. The simulation box size in the layer plane is 5nmx5 nm with a 1024x1024 sampling.

Fig. 2: Electron diffraction calculation of a CoPt NP with a LRO of 0.4 deposited on a 10 nm thick amorphous carbon layer simulated by (a) the random model and (b) the tight-binding model. (c) Radially integrated intensity profile as a function of the scattering vector for both diffraction patterns (a) and (b) in black and red, respectively.

Type of presentation: Poster

IT-16-P-2686 Remove the CCD influence from high-resolution electron microscopy images

Lin F.1, Jin C.2, Yang Y.1
1College of Science, South China Agricultural University, Guangzhou, Guangdong 510642, PR China, 2State Key Laboratory of Silicon Materials, Department of Materials Science and Engineering, Zhejiang University, Hangzhou, Zhejiang 310027, PR China
flin_163@163.com

Slow-scan charge-coupled device (CCD) camera is now an important recording medium for high-resolution transmission electron microscopy (HRTEM). However, the pixel of CCD chips has a certain size, and then the point spread effect cannot be neglected in image. In frequency domain, signal transfer is described by a modulation transfer function (MTF). In addition, Possion noise is inherent in electron images. The noise transfer is described by noise transfer function (NTF), which gives the attenuation of noise power relying on its frequencies [1].
To remove the MTF from HRTEM image, the MTF should be calculated firstly from experiments. All experimental images were recorded on a JEM-2010F TEM equipped with a Gatan-894 CCD camera. The beam-stopper image is shown in Fig. 1(a). Through averaging the image intensities of positions equally distancing from the edge, we could get the edge profile shown in Fig. 2(b). Deconvoluted from the filtered profile of edge, the PSF of CCD was resolved. We show it in Fig. 1(c) and use the Fourier transform to get the MTF of CCD.
To measure the NTF, 16 uniform-illumination images without beam-stoppers and 8 dark-current images were recorded. The ‘NTF 1’ and ‘NTF 2’ in Fig. 2 are estimated from the 16 uniform illumination images of totally ~2190 and ~4760 counts, respectively. Although the electron doses are different, the noise characteristics of NTF are almost the same. The NTF is mainly affected by the accelerated voltage and exposure time.
Based on an improved Wiener deconvolution filter, the restored image I’(u,v) in frequency domain of (u,v) is calculated as [2],
I'(u,v)=I(u,v)MTF*(u,v)/{|MTF(u,v)|2+Pn(u,v)/[PD(u,v)-Pn(u,v)]},
in which, I(u,v) is a HRTEM image, MTF(u,v) is the MTF of CCD in 2-dimensional space, and PD(u,v) and Pn(u,v) are the mean power spectral densities of detected image I(u, v) and additive noise, respectively. PD(u,v) is actually the image diffraction pattern multiplied by its conjugate, and Pn(u,v) is estimated from NTF. The NTF provides the power distributions of noise at various frequencies and is measured from uniform-illumination images. For weak scattering objects, such as few-layer graphene or boron nitride, the image is considerably “uniform” because of the weak scattering of atoms. Fig 3(a) gives a raw HRTEM image of graphene. After deconvolution, the lattices in center region are resolved from noise.


Authors acknowledged supports from the National Science Foundation of China (61172011 and 51222202), National Basic Research Program of China (2014CB932500), the Program for Innovative Research Team in University of Ministry of Education of China (IRT13037), the Fundamental Research Funds for the Central Universities (2014XZZX003-07) and Guangdong Natural Science Foundation (10151064201000006).

Fig. 1: (a) The beam-stopper image. The white line indicates one of the beam-stopper edges. (b) The blue line is the edge profiles averaged from one edge in (a). The red line gives a filtered edge profile. (c) The CCD’s PSF calculated from (b). Ideal step function convoluting with this PSF is the red line in (b).

Fig. 2: The NTF of CCD. ‘NTF 1’ and ‘NTF 2’ are calculated from the images uniformly illuminated under different electron doses after normalization.

Fig. 3: (a) A raw HRTEM image of graphene. (b) The image restored with the improved Wiener deconvolution filter.

Type of presentation: Poster

IT-16-P-2731 Methods for Scanning Transmission Electron Microscopy High Angle Annular Dark Field based for three dimensional analysis of the local composition in solid alloys

Rotunno E.1, Grillo V.1,2, Markurt T.3, Remmele T.3, Albrecht M.3
1CNR-IMEM, Parco delle Scienze 37a, I-43100 Parma, Italy, 2CNR-Istituto Nanoscienze, Centro S3, Via G Campi 213/a, I-41125 Modena, Italy , 3Leibniz Institute for Crystal Growth, Max-Born-Strasse 2, 12489 Berlin, Germany
vincenzo.grillo@cnr.it

We report on a novel approach to quantitatively reconstruct the number and the three-dimensional distribution of guest atoms inside a host matrix by STEM HAADF technique. Since the position of the guest atoms in the column strongly affects the chemical quantification for each column [1][2] this technique allows in addition for an improved quantification of the composition.

Our method is based on the joint analysis of a set of experimental data gained with variable beam convergence and/or defocus. It allows to invert the intensities into an atomic distribution along the columns for any dependence of the probe intensity on the thickness. It is therefore well suited to incorporate channeling effects that are usually neglected in other approaches [3]. We focus here on the systematic variation of the beam convergence that permits to set the maximum of the channeling oscillations at different depth [4].

From Fig 1, showing the dependence of the probe intensity with depth inside an InGaN sample for convergence angles of 9, 15 and 21 mrad, it is evident that with changing the convergence angle we are probing different parts of the sample.

To extract detailed information we represent the guest atom distribution in a given column mathematically as a continuous profile and develop it in terms of harmonic components. The components can be determined from the experimental contrast at variable probe conditions by an inversion of the functional dependencies between these parameters.

To test the application of the method we used simulations with the multislice Frozen Phonon algorithm. The results are shown in fig 2 a,b,c. Using the inversion algorithm we were able to retrieve the map of the number of atoms per column. In fig 2 d,e the real composition per column and that retrieved by the algorithm are compared, showing a substantial agreement.

While the use of a single image permitted to obtain the quantification for a single column with a confidence level less than 60%, the new paradigm permits to obtain a confidence level of 95%. .

We are also able to retrieve an estimation of the guest atoms distribution along z. Fig 2f shows the superposition of the estimated distribution and the actual position. We propose the method as a general framework to perform 3D quantitative analysis to be used in all multiple STEM-HAADF experiments including through-focal experiments.

[1] P.M. Voyles, et al. Ultramicroscopy 96 (2003) 251–273

[2] E. Carlino V. Grillo Physical Review B 71 (2005) 235303.
[3] K.v. Benthem, et al. Appl. Phys. Lett. 87, 034104 (2005)

[4] Y. Peng, et al. Journal of Electron Microscopy 53 (2004) 257


Fig. 1: Probe intensity vs depth inside an InGaN sample for convergences of 9 (a), 15(b) and 21 mrad (c). In fig d the profile in the center of the probe are plotted as a function of z.

Fig. 2: Simulation of a STEM-HAADF image of InGaN with convergence of of 9 (a), 15(b) and 21 mrad (c). Comparison of actual d) and retrieved e) column by column composition map. f) Retrieved “continuous” distribution of In atoms in a column compared with the actual distribution

Type of presentation: Poster

IT-16-P-2748 The forward dynamical electron scattering (FDES) software; a graphics-processing-unit accelerated multislice algorithm

Van den Broek W.1, Koch C. T.1
1Institute for Experimental Physics, Ulm University, Albert-Einstein-Allee 11, 89081 Ulm, Germany
wouter.vandenbroek@uni-ulm.de

The forward dynamical electron scattering (FDES) software is a multislice algorithm taking advantage of the graphics processing unit (GPU) to speed up computation. Further acceleration is attained through efficient computation of the projected atom potential and an approximation of Poisson noise by the inverse Anscombe transform. Thermal diffuse scattering can be taken into account by the frozen phonon approximation. In the following, the three most time consuming processes are labelled A, B and C and their runtimes are compared.

FDES was written in the CUDA programming language [1] (with the CUFFT, CUBLAS and CURAND libraries) and run on a Tesla K20c GPU (NVIDIA).

Calculating the projected potential in real space is not optimal since the potential’s divergence in the origin requires special treatment; furthermore, one faces a trade-off between accuracy and computation speed in deciding on the projected potential‘s cut-off radius. Both problems are overcome in FDES. The four pixels nearest to the atoms are assigned the value of their intersection with a pixel-sized square centered on that atom (A), thus ensuring sub-pixel accuracy of the atom positions, see Fig. 1. Next, the result is convolved with the projected potential and deconvolved with the pixel-sized square (B). Since this (de)convolution is implemented as a multiplication in Fourier space where the scattering factors and the sinc-function are well-behaved, no numerical difficulties are encountered.

The multislice algorithm proper, i.e. the computation of the propagation of the electron wave through the sample (C), is sped-up by computing the necessary Fourier transforms on the GPU.

Since normally distributed values require no pre-processing on the central processing unit, FDES applies the inverse Anscombe transform p = round( 0.25 n2 – 3/8 ) [2] to variables n drawn from a normal distribution with mean 2 ( μ + 3/8 )1/2 − 0.25 μ1/2 and variance 1 − exp( − μ / 0.78 ) to yield approximate Poisson distributed variables p with expectation value μ. This is dubbed Anscombe noise. See Fig. 2.

The runtime was measured on a simulation of a 25x25x15nm3 Al crystal with an FCC lattice (no advantage was taken of this symmetry) with a lattice constant of 0.405nm. The pixels were 0.025nm wide and the slice thickness was 0.1nm. The processes A, B and C took 0.022s, 0.191s and 0.488s respectively, between them accounting for 97% of the total runtime of 0.723s.

[1] NVIDIA CUDA C Programming Guide Version 5.0, 2012.
[2] F.J. Anscombe, Biometrika 35 (1948), 246–254.


The authors acknowledge the Carl Zeiss Foundation as well as the German Research Foundation (DFG, Grant No. KO 2911/7-1).

Fig. 1: Each entry in the list of atoms is assigned to a different processor on the GPU that calculates to which slice the atom belongs and what its four nearest neighboring pixels are in that slice. These pixels values are set to the intersection with a pixel-sized square centered on the atom.

Fig. 2: Left: Histograms of Anscombe distributed values with mean 5.0 and of Gaussian distributed values of mean 5.0 and variance 5.0, contrasted to a Poisson distribution of mean 5.0. (N = 107). Right: Absolute values of the difference of the histograms with the Poisson distribution; the error is nearly always lower for the Anscombe noise.

Type of presentation: Poster

IT-16-P-2768 Influence of the delocalization of inner-shell excitations on atomic-resolution elemental maps

Park M.1, Majert S.1, Kohl H.1
1Physikalisches Institut und Interdisziplinäres Centrum für Elektronenmikroskopie und Mikroanalyse (ICEM), Westfälische Wilhelms-Universität Münster, Wilhelm-Klemm-Straße 10, 48149 Münster, Germany
parkmi@uni-muenster.de

Using elemental maps with atomic resolution obtained by energy dispersive x-ray spectroscopy (EDX) in a scanning transmission electron microscopy (STEM), it is possible to investigate the elemental distribution in a specimen. To analyse the experimental results it is necessary to compare them with calculations in order to distiguish specimen properties from imaging artifacts. We investigate the influence of channeling and of the delocalized excitation of inner-shell electrons on the elemental map.

For our calculations, we use the multislice method [1], which describes the behaviour of electrons passing through a thick specimen. In this theorythe specimen is divided into thin slices treated as a pure phase object.

In a first approximation, we assume that the number of x-ray quanta emitted by an atom is proportional to the electron intensity at its position. This means that the atoms are approximated as being point-shaped. We then replace the point-shaped description of atoms with the delocalized excitation function [2].

To estimate the influence of an atom exicitation probability on elemental maps we also compare the results of both approximations with each other and these with experimental results.

An interface between a strontium titanate crystal (SrTiO3) and a lead titanate crystal (PbTiO3) serves as an example of our calculation.

The results of the simulation with the localized approximation of Ti, Sr and Pb signals in an elemental map of this sample are shown in Figure 1. For this calculation we used an acceleration voltage of 200 kV, assumed an objective lens free of astigmatism, a spherical aberration coefficient Cs = 0,5 mm, and a 15 mrad aperture semiangle [3]. These results are compared with the experimental data from L. J. Allen et al. [4].

The major difference between both results is that the position of the boundary is clear in the simulation but not in the experimental results. This could be the result of the unevenness of interface in the experiment and the use of the localized approximation in the calculation. For a better interpretation of the results we currently investigate how the radius of the atomic column and the boundary changes when considering a delocalized excitation function in the simulation.

[1] Earl J. Kirkland “Advanced Computing in Electron Microscopy” (Plenum Press, New York, 1998) p. 157

[2] D. Von Hugo, H. Kohl, and H. Rose, Optik 79 (1988) p. 19

[3] S. Majert “Simulation atomar aufgelöster Elementverteilungsbilder mit der Multislice Methode”, BSc thesis (2012)

[4] L. J. Allen et al. “Chemical mapping at atomic resolution using energy-dispersive x-ray spectroscopy” (MRS BULLETIN, Volume 37, January 2012) p. 47


Fig. 1: Elemental map of 25 slices (≈10 nm) of an interface between a strontium titanate crystal (SrTiO3) and a lead titanate crystal (PbTiO3) in [001]-orientation simulated with the localized approximation

Type of presentation: Poster

IT-16-P-2950 The partial spatial coherence function and the distribution of scattered electrons

Nguyen D. T.1, Findlay S. D.2, Etheridge J.1,3
1Department of Materials Engineering, Monash University, VIC 3800, Australia, 2School of Physics, Monash University, VIC 3800, Australia, 3Monash Centre for Electron Microscopy, Monash University, VIC 3800, Australia
dan.nguyen@monash.edu

The development of aberration correctors has made images with sub-Ångstrom spatial resolution in scanning transmission electron microscopy (STEM) routinely possible, greatly advancing our knowledge of the relationships between a material's atomic structure, composition, and its properties. It has been noted, however, that possessing atomically-resolved features does not guarantee an image can be quantitatively analysed column-by-column, especially for inhomogeneous materials [1,2]. The mechanism limiting our ability to attribute the spatial origin of the signal to specific atomic sites is the spreading of the electron probe as it travels through the specimen. The finite effective source profile, also referred to as the partial spatial coherence function, has been reported to have a large impact on the intensity distribution in STEM images [3,4]. However, the consequences of the effective source distribution for the spreading of the electron probe have not been much explored.

We investigate the implications of the partial spatial coherence function for quantitative analysis in STEM, especially for interpreting the spatial origin of imaging and spectroscopy signals. In particular, we examine how the shape of the source distribution, especially the length of its “tails”, influences the degree to which the electron probe scatters onto adjacent atomic columns. This was explored via the use of three different source distribution models applied to a GaAs crystal case study. The shape of the effective source distribution was found to have a large influence not only on the STEM image contrast, but also on the distribution of the scattered probe through the specimen and hence on the spatial origin of the detected electron intensities.

This has implications for our ability to extract column-by-column information via annular dark field, X-ray and electron energy loss STEM imaging, as precise knowledge of the spatial origin of the measured signal is a prerequisite for any high precision determination of structure, bonding and chemical composition at the atomic scale.

 

References
1. C. Dwyer and J. Etheridge, Ultramicroscopy, 96 (2003) 343-60.
2. P. M. Voyles, J. L. Grazul and D. A. Muller, Ultramicroscopy, 96 (2003) 251-73.
3. J.M. LeBeau, S.D. Findlay, L.J. Allen, S. Stemmer, Phys. Rev. Lett., 107 (2008) 206101
4. C. Dwyer, C. Maunders, C.L. Zheng, M. Weyland, P.C. Tiemeijer, J. Etheridge, Appl. Phys. Lett., 100 (2012) 191915

 


This research was supported under the Discovery Projects funding scheme of the Australian Research Council (Project Nos. DP120101573 and DP110101570).

Fig. 1: Comparison between a Gaussian (G), Gaussian + Lorentzian (G+L) and Lorentzian (L) source profile, with the FWHM chosen to give a 60% reduction in contrast with respect to a coherent probe. The resultant ADF-STEM images are shown below it (from left to right: Gaussian, Gaussian + Lorentzian, Lorentzian)

Fig. 2: The proportion of ADF signal generated from a single Ga atomic column in GaAs <001> recorded in a Voronoi cell around the column (top) and in an empty cell immediately adjacent to it (bottom). P refers to a diffraction-limited probe. G, G+L and L correspond to a Gaussian, Gaussian + Lorentzian and Lorentzian effective source, respectively.

Fig. 3: Cross-sectional intensity map of the scattered probe (using three different source profiles) through a GaAs crystal of <110> and <112> orientation, and the corresponding plot of the normalised probe intensity on the Ga atomic column (integrated within a radius of 0.5 Å). The arrows point to the centre Ga column and the closest As column.

Type of presentation: Poster

IT-16-P-2955 Removing the effects of elastic and thermal scattering from spectrum images in scanning transmission electron microscopy

Lugg N. R.1,2, Neish M. J.1, Haruta M.3, Kothleitner G.4, Grogger W.4, Hofer F.4, Kimoto K.3, Mizoguchi T.5, Findlay S. D.6, Allen L. J.1
1School of Physics, The University of Melbourne, Melbourne, Australia, 2Institute of Engineering Innovation, The University of Tokyo, Tokyo, Japan, 3National Institute for Materials Science, Tsukuba, Japan, 4FELMI, Graz University of Technology, Graz, Austria, 5Institute of Industrial Science, The University of Tokyo, Tokyo, Japan, 6School of Physics, Monash University, Melbourne, Australia
nrlugg@sigma.t.u-tokyo.ac.jp

Scanning transmission electron microscopy (STEM) has proven to be a powerful tool for the acquisition of atomic-column-resolved elemental maps using electron energy-loss spectroscopy (EELS) and, more recently, energy-dispersive x-ray (EDX) spectroscopy. Furthermore, in EELS, given the ability to study how spectra change on the atomic scale, there has also been much interest in mapping not only the positions of atoms, but also their bonding states by studying changes in the fine structure of EELS data [1-3]. Specifically, atomic-resolution oxygen EELS data in transition metal oxides can potentially provide information about the entire electronic structure of a material since oxygen atoms bonded to different transition metals will display distinct spectra, reflecting the local bonding state. In atomic-resolution EDX, given the localized nature of the interaction potential involved and that such maps are not further complicated by subsequent elastic and thermal scattering after ionization, there is huge promise in quantitatively measuring elemental densities at atomic resolution.

However, due to the complex elastic and inelastic scattering of the electron probe, direct qualitative and quantitative interpretation of both elemental and bond maps is difficult. In bond mapping, the scattering of the electron probe causes the spectra from inequivalently bonded atoms to become mixed, and the features that distinguish them to become ambiguous. In elemental mapping, the highly non-linear response to atomic-column densities due to electron scattering denies any direct correspondence between signal and atomic-column densities [4].

Recently, a method has been developed to remove the effects of elastic and thermal scattering from spectrum images [5]. Using the quantum excitation of phonons model [6], the cross-section expression for inelastic scattering in STEM may be expressed as an inverse problem, and the elastic and thermal scattering deconvolved from experimental data. Here we show applications of this method in both EELS [7] and EDX [8] data of SrTiO3.

[1] DA Muller et al, Science 319 (2008) 1073
[2] H Tan et al, PRL 107 (2011) 107602
[3] M Haruta et al, APL 100 (2012) 163107
[4] BD Forbes et al, PRB 86 (2012) 024108
[5] NR Lugg et al, APL 101 (2012) 183112
[6] BD Forbes et al, PRB 82 (2010) 104103
[7] MJ Neish et al, PRB 88 (2013) 115120
[8] G Kothleitner et al, PRL 112 (2014) 085501


Fig. 1: (a) HAADF map of SrTiO3 [001] (projected structure inset). (b) EELS map using the O K edge (potential obtained from inversion inset). (c) SrTiO3 structure showing the two inequivalent O atoms in the [001] projection. Spectra from inequivalent O columns obtained from (d) background-subtracted data, (e) after inversion and (f) Wien2k calculation.

Fig. 2: SrTiO3 [001] experimental and simulated (inset) (a) HAADF (structure overlayed) and EDX (b) Sr K, (c) Ti K (d) O K edge maps. Numbers inset show atomic densities (atom/nm3) obtained from: the ideal structure (averaging the entire EDX map) [averaging specific columns] {inversion}. (e) EDX colour composite with masks used for column specific average.

Type of presentation: Poster

IT-16-P-3028 Phase mapping at the interface retrieved by FFT based transport of intensity equation

Zhang X.1, 2, Oshima Y.2, 3
1Tokyo Institute of Technology, Tokyo, Japan, 2JST-CREST, Tokyo, Japan, 3Osaka University, Osaka, Japan
zhang.x.am@m.titech.ac.jp

   Phase retrieval using transport of intensity equation (TIE) is convenient and powerful to obtain electrostatic potential of materials, because it needs only three transmission electron microscope (TEM) images taken at different foci. In our previous results, potential map of gold nanoparticles adsorbed on a thin amorphous carbon (a-C) film was successfully obtained with high spatial resolution (1nm). However, it is difficult to obtain potential map of two-phase interface. In this study, we investigate how to obtain the phase map of the boundary between a-C and vacuum regions quantitatively.
  TEM observation was taken by 50pm-resolved R005 equipped with cold field emission gun and double aberration-correctors. Figure 1(a-c) shows three TEM images of under-focus 1000 nm, just-focus and over-focus 1000 nm. The TIE retrieved phase map using these three images is given in (d). Figure 1(e) and (f) are the line profiles along the blue lines indicted in (a) and (d). We notice that the mean intensity of vacuum, thin C-film and thick C-film regions are different in the original TEM image and the intensity profile does not reflect the expected phase shift among these regions. Moreover, a strong low-frequency contrast appeared in the phase map.
  We found the reasons of such discrepancy. Firstly, the periodic boundary condition required for FFT process of solving the TIE equation influences the retrieved phase map. Secondly, brightness difference exists in the focal-series TEM images, which is caused by the variation in objective lens current. Since the intensity difference applied to TIE is obtained by subtracting the over-focus TEM image from the under-focus one, the difference in image brightness causes deviation in the intensity differences. For example, all the values of ΔI are negative in the vacuum area. However, they should only oscillate around zero according to the wave theory.
  We consider that applying padding or mirror methods to the process is effective to satisfy the boundary condition, and aligning the current center rather than the voltage center of objective lens can eliminate the inhomogeneous illumination when taking a focal-series TEM images.


Fig. 1: (a-c) Three TEM images taken at under-focus 1000 nm, just-focus and over-focus 1000 nm. (d) The TIE retrieved phase map obtained using (a-c). (e,f) The line profiles along the blue lines indicted in (a) and (d), respectively.

Type of presentation: Poster

IT-16-P-3032 Computer vision in the service of Crystallography: Automated analysis of atomic-resolution images

Klinger M.1
1Institute of Physics ASCR, Prague, Czech Republic
miloslav.klinger@seznam.cz

An automated tool for a crystallographic analysis of HRTEM (High resolution Transmission electron microscopy), HRSTEM (High resolution Scanning Transmission electron microscopy) and diffraction images is proposed. Algorithms of artificial intelligence and computer vision are employed to detect features carrying the information and to process them in order to determine or estimate crystallographic quantities. This shall result in an expedited analysis, possibly higher precision and little to no human effort compared to manual analysis.

In the case of SAD (Selected area diffraction) images, diffraction spots or disks are detected in the widest possible area of the pattern and the zone axis is calculated. If the observed material is not known, the tool can choose the most probable candidate from a list candidates. HREM (High resolution electron microscopy) images can be segmented to separate individual grains depicted. If the image contrast features directly correspond to positions of atomic columns, the zone axis is determined and the crystallographic planes and direction in the image are identified. Dislocation detection and quantification can be performed as well as a grain misorientation estimation, reconstruction of positions of individual atoms and so on. If the image contrast features do not correspond directly to the atomic column positions, the atomic columns can be found using HREM simulations.

The proposed tool (implemented in MATLAB) has been successfully tested on number of real world images and diffraction patterns. It has proven its ability to autonomously provide correct results.


I would like to thank Professor Michael Mills and the National Center for Electron Microscopy for providing HRTEM images. Financial support offered by GACR GBP108/12/G043 and MEYS LM2011026 is appreciated.

Fig. 1: Segmented grains in HRTEM image of alluminium. Original image acquired in the National Center for Electron Microscopy.

Fig. 2: Three dimensional reconstruction of grains depicted on Fig. 1.

Fig. 3: Dislocation detected in alluminium. Burgers circuit can be seen on the left and visualization of inserted plane on the right. Original image acquired by Professor Michael Mills.

Type of presentation: Poster

IT-16-P-3168 Simultaneous thickness and orientation mapping by dark-field transmission electron microscopy

Tyutyunnikov D.1, Koch C. T.1
11. Institute for Experimental Physics, Ulm University, Albert-Einstein-Allee 11, 89081 Ulm, Germany
dmitry.tyutyunnikov@uni-ulm.de

Transmission electron microscopy (TEM) is a powerful tool for investigating the atomic structure and morphology of nano- and micro-objects. To reveal the structure of a material the specimen prepared for TEM should be reasonably thin in order to be transparent for the electron beam. The thickness of the sample is an important parameter one should account for when analyzing images acquired in TEM. The probability for electrons to scatter multiple times increases with the specimen thickness. This effect has, for example, a strong influence on the contrast of high-resolution TEM images. There are several techniques available to estimate the thickness by TEM [1]. These methods can be divided into several categories, e.g. imaging methods, Convergent Beam Electron Diffraction (CBED) method, electron energy-loss spectroscopy (EELS) - based methods, and methods based on X-ray spectroscopy (EDXS). The most popular technique among them is based on EELS [2], which can be used for a variety of samples and is easy to implement computationally. However if one has thin crystalline samples one often has to deal with bending due to lattice relaxation. Bending can strongly influence the thickness values obtained by EELS. Here we report about a new technique which simultaneously delivers thickness and specimen surface orientation maps. Our approach is based on analysis of rocking curves extracted from experimental dark-field (DF) images acquired at different specimen tilts. We fit the parameters that affect dynamical electron diffraction rocking curves to experimental DF images. In its simplest version, our approach uses 2-beam theory, for which the intensity of the diffracted beam is given according to C.Humhreys 1979 review [3]. To determine the thickness we fit the power spectra to a sum of Gaussians. Since the rocking curves usually exhibit oscillatory behavior reflecting the thickness of the specimen one should observe peaks in power spectra of rocking curve. Fitting was done in MATLAB by using the unconstrained nonlinear optimization routine fminsearch [4] and the Matlab Curve Fitting Toolbox. As an example the mapping was done for a commercial semiconductor device. Figure 1 shows a single slice from DF image stack of this device acquired for the {220} reflection. Figure 2 illustrates how a dark-field tilt series samples reciprocal space to demonstrate the oscillatory behavior of the extracted rocking curve (inset in Fig. 1).

[1] D.B. Williams and Carter C.B. Transmission Electron Microscopy. Springer Science+Business Media, 2009

[2] T. Malis, et al. J. Electron Microsc.Tech., 1988

[3] C. J. Humphreys. Rep. Prog. Phys., 42, 1979.

[4] J.C. Lagarias, et al. SIAM Journal of Optimization, pages 112-147, 1998


The authors acknowledge financial support by the Carl Zeiss Foundation as wellas the German Research Foundation (DFG, Grant No. KO 2911/7-1). The authors also acknowledge Stuttgart Center for Electron Microscopy (StEM) for sample preparation and possibility to carry out the experiment.

Fig. 1: Single slice from a DF image stack acquired for the {220} reflection of Si in a MOSFET structure. The red square shows the ROI used for thickness and orientation mapping. The sample was prepared by automatic tripod polishing and consequent low-voltage Ar ion milling at T of liquid N2

Fig. 2: The Dark-field tilt series intersects the diffraction signal with the Ewald sphere indicated as green arks at different positions along kz axis. The diffraction signal was calculated by FFT for a 20 nm thick slab.

Fig. 3: Thickness map in Å. The data was binned before fitting. The thickness fitted outside the crystalline area is meaningless, of course.

Fig. 4: Map of the misorientation a0.

Type of presentation: Poster

IT-16-P-3170 Structure factor refinement from electron diffraction for structures with arbitrarily large unit cells

Feng W.1, Kazzazi A.1, Koch C. T.1
1Institute for Experimental Physics, Ulm University, Albert-Einstein-Allee 11, 89081 Ulm, Germany
feng.wang@uni-ulm.de

Structure-factor refinement by quantitative convergent-beam electron diffraction (QCBED) [1] has been able to reveal the charge distribution responsible for the bonding between atoms [2]. However, in order to be able to fit a set of complex structure factors by comparing dynamical electron diffraction simulations to the contrast within CBED discs, the sample must typically be at least 100 nm thick, and the lattice parameters should not exceed 1 nm. Also, in order to extract the elastic scattering signal it must be assumed that the incoherent background (mainly thermal diffuse scattering (TDS)) varies only slowly, an assumption that is not generally correct. In contrast, large-angle rocking-beam electron diffraction (LARBED) [3] data (see, for example Fig. 1) can be acquired for arbitrarily large unit cell structures and reveals features that are clearly due to dynamical electron diffraction even at specimen thicknesses below 10 nm. This makes structure factor refinement from nanocrystals possible.

As is common in QCBED work we write the scattering matrix S in the form S=eiTA, by assuming previous knowledge of the excitation errors (diagonal of A), we reduce the problem to finding the factor T, which is acceleration voltage and specimen thickness related, and all the off-diagonal entries Ug-h in A, from a series of observed norms of entries in one column of S (the squared norms correspond to the diffraction intensities shown the example pattern in Fig. 1). To efficiently solve such a nonlinear programming problem up to several hundreds of variables, a gradient-based iterative method is critical.

In this work, we design a two-layer model to accelerate the calculation of the expensive Bloch-wave simulation and the cumbersome gradient approximation. We also compare and discuss the convergence and performance of different optimization algorithms applied to this problem.

[1] C. Deininger, G. Necker, J. Mayer, Ultramicroscopy 54 (1994) 15-30

[2] J.-M. Zuo, M. Kim, M. O’Keefe, J.C.H. Spence, Nature 401 (1999) 49

[3] C.T. Koch, Ultramicroscopy 111 (2011) 828 – 840


The authors acknowledge the Carl Zeiss Foundation as well as the German Research Foundation (DFG, Grant No. KO 2911/7-1)

Fig. 1: Bright-field and dark-field large-angle rocking convergent beam (LARBED) patterns of SrTiO3 in the [-110] zone axis, acquired on the Zeiss SESAM operated at 200 kV. This kind of data can be used to fit dynamical scattering equations to, even for very thin samples.

Type of presentation: Poster

IT-16-P-3277 BlochSim: A new, free to use, open source Bloch Wave Simulation program

Evans K. L.1, Roemer R. A.1,2, Beanland R.1
1Department of Physics, University of Warwick, Coventry, West Midland, CV4 7AL, 2Centre for Scientific Computing, University of Warwick, Coventry, West Midlands, CV4 7AL
keith.evans@warwick.ac.uk

We present a new software package developed at the University of Warwick for simulation of dynamical diffraction in transmission electron microscopes using the Bloch wave method [1]. While the Bloch wave method is well established [2], freely available software to scientists is limited, less still is open source [3]. This prevents further development and adoption by the general microscopy community. The X-ray crystallography community has benefited hugely from open source and freely available packages such as SHELX [4], and availability of similar platforms in electron crystallography are very desirable. Our new software is intended to help bridge this gap by being open source and free of charge in order to facilitate greater usage throughout the community and ease of development.

The software package allows for simulation of Convergent Beam Electron Diffraction (CBED) patterns as well as the recently developed Digital Large Angle CBED (D-LACBED) [5]. More advanced features include crystal thickness determination, structural refinement and electron density mapping. All of these features take can advantage of the superior data range of D-LACBED [5].

Historically, structure solution and refinement has been dominated by X-ray methods. However, X-rays lack the nanometre size resolution required for the study of nanostructured materials such as interfaces, grains, nanobeams etc. Electrons, due to their much smaller wavelengths and stronger interaction with matter, have the required resolution and hence there is a niche for structural refinement using electron diffraction. We will describe the use of BlochSim with structure refinement strategies, in particular identifying and taking advantage of regions within large electron diffraction datasets which have enhanced sensitivity to structural parameters.

BlochSim uses a set of base tools that are compatible with X-ray crystallography and a key aim is to make the software multiplatform and user friendly. It will be configured to accept many forms of input structure (.cif, .pdb, .xyz) in order to reach a large user base. We encourage use and further code development by new users.

1. Bloch, F., Über die Quantenmechanik der Elektronen in Kristallgittern. Zeitschrift für Physik, 1929. 52(7-8): p. 555-600.

2. Stadelmann, P.A., A Software Package for Electron-Diffraction Analysis and HREM Image Simulation in Materials Science. Ultramicroscopy, 1987. 21(2): p. 131-145.

3. Tsuda, K. 2013; Available from: www.tagen.tohoku.ac.jp/labo/terauchi/personal/tsuda/mbfit_win.zip.

4. Sheldrick, G.M., A short history of SHELX. Acta Crystallographica Section A, 2008. 64: p. 112-122.

5. Beanland, R., et al., Digital electron diffraction - seeing the whole picture. Acta Crystallographica Section A, 2013. 69: p. 427-434.


Type of presentation: Poster

IT-16-P-3287 Electron vortex beam diffraction via multislice solutions of the Pauli equation

Edström A.1, Rusz J.1
1Department of Physics and Astronomy, Uppsala University
alexander.edstrom@physics.uu.se

Electron magnetic circular dichroism (EMCD) has gained plenty of attention as a possible route to high resolution measurements of, for example, magnetic properties of matter via electron microscopy. However, certain issues, such as low signal-to-noise ratio, have been problematic to the applicability. In recent years, electron vortex beams\cite{Uchida2010,Verbeeck2010}, i.e. electron beams which carry orbital angular momentum and are described by wavefunctions with a phase winding, have attracted interest as potential alternative way of measuring EMCD signals. Recent work has shown that vortex beams can be produced with a large orbital moment in the order of l = 100 [6, 7]. Huge orbital moments might introduce new effects from magnetic interactions such as spin-orbit coupling.

The multislice method[2] provides a powerful computational tool for theoretical studies of electron microscopy. However, the method traditionally relies on the conventional Schrödinger equation which neglects relativistic effects such as spin-orbit coupling. Traditional multislice methods could therefore be inadequate in studying the diffraction of vortex beams with large orbital angular momentum. Relativistic multislice simulations have previously been done with a negligible difference to non-relativistic simulations[4], but vortex beams have not been considered in such work.

In this work, we derive a new multislice approach based on the Pauli equation, Eq. 1, where q = −e is the electron charge, m = γm0 is the relativistically corrected mass, p = −i is the momentum operator, B = ∇ × A is the magnetic flux density while A is the vector potential and σ = (σx , σy , σz ) contains the Pauli matrices. ψ±(r) represent wavefunctions for spin up (+) and down  (-) electrons. A two component fast electron equation of the form given in Eq. 2 is found. The solutions of this equation are studied computationally via a real-space multislice[1] approach. Results are presented for large orbital angular momentum vortex beams passing through model systems, such as bcc Fe, and are compared to results from the traditional Schrödinger equation  based calculations.


Fig. 1: Equations referred to in the text.

Fig. 2: References referred to in the text.

Type of presentation: Poster

IT-16-P-3331 Atomically Resolved Low-Loss Imaging of Graphene in the Aberration-Corrected STEM

Oxley M. P.1, Kapatenakis M. D.1, Prange M. P.3, Zhou W.2, Idrobo J. C.2, Pennycook S. J.4, Pantelides S. T.1
1Vanderbilt University, Nashville, TN, USA, 2Oak Ridge National Laboratory, Oak Ridge, TN, USA., 3Pacific Northwest National Laboratory, Richland, WA, USA., 4University of Tennessee, Knoxville, TN, USA.
oxleymp@gmail.com

Aberration correction in the scanning transmission electron microscope (STEM) has led not only to improved resolution but also increased contrast and sensitivity at lower accelerating voltages. Imaging of beam sensitive two dimensional materials at atomic resolution has been enabled by operating at energies below the knock on threshold. In this way, single atom impurities have been imaged in BN using annular dark field (ADF) imaging [1] and electron energy-loss spectra (EELS) obtained in graphene with high spatial resolution [2]. Improved spectroscopic sensitivity has even allowed the measurement of energy-loss near-edge spectra (ELNES) providing information about local bonding of single impurity atoms in graphene [3]. We extend recent theoretical developments allowing the simulation of core-shell ELNES as a function of probe position to examine inelastic image formation based on low-loss spectroscopy [4].

In Fig. 1 we show an ADF image and simultaneously acquired low loss spectrum of graphene obtained using ORNL’s Nion UltraSTEM100 operating at 60 kV. Spectrum images derived from the energy ranges of 20-26, 34-40 and 50-56 eV are also shown. The 34-40 eV image shows resolved atomic columns while the other images show no apparent contrast. It should be noted that the intensity variation is of the order of 5% and the image would be considered delocalized by most commonly used definitions. Such low levels of contrast are only visible due to the increased sensitivity of modern instruments.

Preliminary image simulations based on first-principles dynamical electron scattering and density functional theory are shown in Fig. 2. Statistical noise at a level of 5 % has been added to show the effects of experimental noise on such low contrast images. While the simulated lower energy images both show graphene-like contrast, the addition of noise significantly reduces the visibility of the graphene structure at the lowest energy, while the graphene ring structure is still evident at the higher energy. Graphene like contrast is not observed at the higher energy loss. We will discuss the mechanisms underlying image contrast for inelastic STEM imaging based on low-loss spectroscopy.

References

[1] O.L. Krivanek et al., Nature 464, (2010), 571.

[2] W. Zhou et al., Microsc. Microanal. 18, (2012), 1342

[3] W. Zhou et al., Phys. Rev. Lett. 109, (2012), 206803.

[4] M. P. Prange et al., Phys. Rev. Lett. 109, (2012), 246101.


This work was supported by DOE Grant No. DE-FG02-09R46554, by the DOE Office of Basic Energy Sciences, Materials Sciences and Engineering Division, by NSF grant No. DMR-0938330.

Fig. 1: Experimental ADF image and low-loss spectrum of graphene. Images derived from the indicated energy-loss regions of the spectrum are also shown. Scale bars are 1 Å.

Fig. 2: Theoretical spectrum and images derived from the indicated portions of the spectrum. Noise has been added to allow visual comparison with experiment.

Type of presentation: Poster

IT-16-P-3434 Super-Resolution applied to Magnetite boundaries images

Bárcena G.1, Guerrero M.1, Guerrero E.1, Kepaptsoglou2 D.2, Gilks D.3, Lari L.3, Lazarov V. K.3, Galindo P. L.1
1Department of Computer Science and Engineering, Universidad de Cádiz, 11510 Puerto Real, Spain, 2SuperSTEM Laboratory, STFC Daresbury Campus, Warrington, WA4 4AD, United Kingdom, 3Department of Physics, University of York, Heslington, York, United Kingdom
guillermo.barcena@uca.es

High-Angle Annular Dark-Field (HAADF) imaging is a useful tool to understand the nature of the interaction between different materials domain boundaries, but one must overcome the limitations imposed by the characteristics of the microscope that directly affects resolution. An approach to face this problem is to deal with super-resolution techniques. These techniques attempt to obtain high-resolution images from several observed low-resolution images captured from the same scene, thus the resolution of an image can be improved by bringing out details that might otherwise not be seen.
In this work we illustrate the application of super-resolution techniques to a series of 10 low resolution HAADF images of Magnetite (Fe3O4), oriented along the 001 direction.
Since classical super-resolution reconstruction programs running on a standard computer may take up to 6 hours to get the results, a specialized software suite running in GPUs [1] has been developed to speed up this process, and now results can be obtained just in 10 minutes.

Figure 1 show an experimental low resolution image of Magnetite where the presence of noise is noticeable. Three atoms of Fe have been marked in figure 1 and the corresponding intensity profile is plotted in figure 2, but just two intensity peaks can be appreciated.
In the super-resolution approach two steps are applied: alignment and reconstruction. In this work the alignment process is carried out by filtering the image with a Gaussian filter and then applying the Vandewalle’s modification [4]. Then, a variant of the Non Local Mean algorithm [2,3] is used in order to obtain a high-resolution image, where noise has been substantially reduced, so that the three atoms of Fe can be clearly identified, as shown in figure 3. This fact is made apparent in the corresponding intensity profile shown in figure 4.
These results indicate that super-resolution techniques can provide enhanced HAADF images in terms of resolution, quality and details definition.

[1] Bárcena-González, G. M. Sc. Thesis. Study and application of Superresolution's algorithms in electron microscopy images. October, 2013

[2] Binev, P., Blanco-Silva, et al (2012). High-quality image formation by nonlocal means applied to high-angle annular dark-field scanning transmission electron microscopy (HAADF–STEM). Modeling Nanoscale Imaging in Electron Microscopy, 127-145.

[3] Buades, A., Coll, B., & Morel, J. (2005). A non-local algorithm for image denoising. Computer Vision and Pattern Recognition, 2005. CVPR 2005. IEEE Computer Society Conference on, 2 60-65.

[4] Vandewalle, P. et al (2004). Double resolution from a set of aliased images. Proc. SPIE 5301, Sensors and Camera Systems for Scientific, Industrial, and Digital Photography Applications V, 374


Fig. 1: An experimental low resolution image of Magnetite, three atoms of Fe has been marked with a red circle. The presence of noise is noticeable, in fact, just two atoms can be clearly appreciated.

Fig. 2: Intensity profile corresponding to the marked area in figure 1, Two of the three intensity peaks can be observed.

Fig. 3: High-resolution image obtained by super-resolution techniques, noise has been substantially reduced, so that the three atoms of Fe marked with a red circle can be clearly observed.

Fig. 4: Intensity profile corresponding to the marked area in figure 3. The three atoms of Fe can be clearly identified.

Type of presentation: Poster

IT-16-P-3481 Chiral-dependent electron vortex energy loss spectroscopy

Yuan J.1, Lloyds S.1, Babiker M.1
1Department of Physics, University of York, Heslington, York, YO10 5DD, United Kingdom.
jun.yuan@york.ac.uk

Chiral electron-vortex beams, carrying a well-defined orbital angular momentum (OAM) about the propagation axis, are potentially useful as probes of magnetic and other chiral materials. In particular, it has been proven that, unlike the optical vortex beams, electron vortex beams can directly induce dipole transitions between states normally assessable only with circular polarised light absorption. This has lead to the expectation that electron vortex beams can be used to acquire chiral dichrotism spectroscopy which will be particular useful for characterization of magnetic materials. Experimentally, the situation is confusing despite initial result [2,3]. To understand this, we present a theory [4] based on an effective operator, expressible in a multi-polar form, describing the inelastic processes in which electron-vortex beams interact with atoms, including those present in Bose-Einstein condensates, involving exchange of OAM. We show clearly that the key properties of the processes are dependent on the dynamical state and location of the atoms involved as well as the vortex-beam characteristics. The later is due to the extrinsic nature of the orbital angular momentum and distinguish chiral electron energy vortex beam energy loss spectroscopy from circular magnetic dichrotic spectroscopy. Our results can be used to identify optimal experimental schemes in which chiral-specific electron-vortex spectroscopy can probe magnetic sublevel transitions normally studied using circularly polarized photon beams with the advantage of atomic-scale spatial resolution. One of the schemes is shown in Fig. 1 where localization of the energy-loss signal is achieved through confocal arrangement and dipole transition selected through orbital-angular-momentum analyzer. [1] S. Lloyds, M. Babiker and J. Yuan, Phys. Rev. Lett. 108 (2012), 074802 [2] J Verbeeck et al, Nature 467 (2010), p. 301. [3] P. Schattschneider et al, Ultramicroscopy, 136 (2014) p81-5 [4] J Yuan, S. Lloyds and M. Babiker, Phys. Rev. A88, 031801


The authors gratefully acknowledge funding from the UK EPSRC (Grant No. EP/J022098).

Fig. 1: The experimental arrangement optimized for spatially resolved chiral-dependent electron vortex electron energy loss spectroscopy.

Type of presentation: Poster

IT-16-P-5927 Simultaneous Imaging of O Atoms and relative heavier Sm atoms in Iron-based Superconductor SmFeAsO0.85F0.15 by HRTEM

Wang Y M Ge B H Che G C
Beijing National Laboratory for Condensed Matter Physics, Institute of Physics, Chinese Academy of sciences, Beijing, China
wangym@iphy.ac.cn

High-resolution imaging of light elements using electron microscopy has always been a challenge. Image resolution is mainly limited by lens aberrations, especially the spherical aberration of the objective lens. Image deconvolution, as a special kind of image processing in High-resolution transmission electron microscopy could remove the image distortion by lens aberrations and transform an arbitrary image intuitively into the structure map, the resolution of which is limited by the information limit of electron microscope. Electron diffraction, which is not restricted by lens aberrations could overcome the resolution limitation. Here we show a combination of electron diffraction and image deconvolution to study the crystal structure of iron-based superconductor SmFeAsO1.85F0.15. Usually it is unexpected to distinguish O atoms in HRTEM images taken with a conventional microscope, especially those adjacent to relative heavier atoms at a very close distance. For the iron-based superconductor SmFeAsO1.85F0.15, every atom, not only heavy atoms Fe and As with atomic space 1.36 Å but also light O and relative heavier Sm atoms with 1.17 Å can be resolved individually in the final image (Fig.3) by using the image deconvolution combined with the corrected electron diffraction data. This is for the first time to image the oxygen atoms which is adjacent to the heavier atoms with the distance of 1.17 Å using the conventional electron microscope by this approach. It allows us not only to determine the crystal structure at atomic level but also to simultaneously reveal light and heavy atoms with a relatively big difference in atomic number and a much smaller atomic distance than the microscope resolution for related compounds.

[1] F.H. Li, Phys. Status Solidi A 207 (2010) 2639-2665.

[2]H.F. Fan, Z.Y. Zhong, C.D. Zheng, F.H. Li, Acta Cryst. A41 (1985)163-165.

 


This project is supported by the National Natural Science Foundation of China (Grant No. 51102275) and by the National Basic Research Program of China (973 Program, No. 2011CBA00101).

Fig. 1:  Fourier filtering image

Fig. 2: Deconvoluted image  at defocus value -600 Å. Rectangles indicate the projections of unit cells.

Fig. 3: Projected potential map obtained from Fig. 2 after phase extension in combination with the diffraction intensity correction.

IT-17. Atom probe and non-traditional microanalytical tasks

Type of presentation: Invited

IT-17-IN-3294 Correlative microscopy applied to atom probe specimen preparation – application to selected metallurgical problems

DANOIX F.1, GOUNE M.4, CUVILLY F.1, CAZOTTES S.2, ALLAIN S.3, ZAPOLSKY H.1
1Université de Rouen, GPM, Avenue de l’université, 76801 St Etienne du Rouvray - France, 2INSA de Lyon, MATEIS Bât. B. Pascal, 7 Avenue Jean Capelle 69621 Villeurbanne Cedex - France, 3Matériaux-Métallurgie-Nanosciences-Plasmas-Surfaces, IJL (Institut Jean Lamour) Parc de Saurupt 54000 NANCY - France, 4Université de Bordeaux, Institut de Chimie de la Matière Condensée de Bordeaux, UPR CNRS 9048, 87 Avenue du Docteur Schweitzer, 33608 PESSAC cedex - France
frederic.danoix@univ-rouen.fr

For more than 10 years, site specific specimen preparation for atom probe tomography using FIB and lift out techniques has drastically widened the possible applications of the technique. In particular, it made it possible the preparation of semiconductor or insulator specimen that laser assisted instrument can now analyze almost routinely. FIBs are systemically implanted in SEM chambers, where other imaging devices are also present. In particular, EBSD is now a common equipment in many labs. It provides complementary crystallographic information to SEM surface micrographs, which can be used in different ways. However, very little has been done using these crystallographic information provided by EBSD for APT specimen preparation. In this presentation, we show the various advantages of combining lift out and EBSD techniques, illustrated by different applications from localization of low volume fraction phases (metallographic aspect of EBSD) to ageing under external load and internal interfaces selection, where the orientation of each specific grain is an important parameter. In addition, the preparation of specimens with specific crystal orientation is a new approach to study and control specific experimental artefacts, such as chromatic aberrations, local magnification effect and surface diffusion.
Selected examples in various material science fields will illustrate theses experimental aspects, from carbon redistribution between phases in modern steels, the effect of the crystal orientation on the phase separation under uniaxial load in FeCr model alloys, and the chemical segregation at internal interfaces and defects.


Fig. 1: Influence of the crystal orientation on the internal nanostructure in a FeCr model alloy aged at 500°C

Type of presentation: Invited

IT-17-IN-6090 Atomic scale studies of nanostructured and nanoscale materials with atom probe tomography

Cairney J. M.1, Felfer P. J.1, Scherrer B.1,2
1Australian Centre for Microscopy & Microanalysis, The University of Sydney, Sydney, NSW 2006 Australia, 2Nanometallurgy, ETH Zurich, 8093 Zurich, Switzerland
julie.cairney@sydney.edu.au

Atom probe tomography is a powerful microscopy technique that provides atomic-resolution 3D maps that show the precise location of the atoms within a volume of material [1]. It has seen widespread use for the characterisation of bulk metals and alloys, but new developments in specimen preparation and the use of lasers have now made it applicable to the study of a much wider range of material types. This presentation will provide an overview of a range of different studies involving ‘non-traditional’ functional and nanoscale materials. This includes bulk Pt/ZrO2-multilayers, which are model systems for applications like the newest generation of micro solid oxide fuel cells and sensors metallic glass nanowires. In these samples, which consist of nanocrystalline layers 10-40 nm thick, we are interested in the solubility and diffusion of oxygen in the noble metal layers. However, the study of such layers in atoms probe is subject to limitations due to the large difference in the field evaporation behaviour between the metal and the oxide. We will discuss how these issues have been addressed in our study, and the observed location of the oxygen atoms. We will also show recent results from the study of bimetallic Au@Ag core-shell nanoparticles. The catalytic performance of these particles greatly influenced by the distribution of the atoms of each element within the particle and on the particle’s surface. However, almost no quantitative, experimental data is currently available on the precise location of the individual atoms within particles less than 100 nm in size. We will demonstrate atom probe can be used to quantitatively determine the distribution of the individual chemical elements in 3D both inside and on the surface of nanoparticles extracted directly from a suspension. Other examples will include ion-irradiated bulk metallic glass nanowires, in which the distribution of embedded Ga ions is being investigated in order to understand their influence on the plasticity of the nanowires. In each case we will discuss how challenges around the specimen preparation have been overcome, any artifacts that are expected to arise in the data (and how these have been addressed), and the new analysis methods applied. References [1] B. Gault, S.P. Ringer, M.P. Moody, J.M. Cairney, Atom Probe Microscopy, Springer, 2012.


The authors acknowledge the facilities and the scientific and technical assistance of the Australian Microscopy & Microanalysis Research Facility at the Australian Centre of Microscopy & Microanalysis

Type of presentation: Oral

IT-17-O-2516 Glass analysis with Atom Probe Tomography utilizing a rapid nanotip preparation technique

Bell D. C.1, Magyar A. P.2, Graham A. C.2
1Harvard University, School of Enginnering and Applied Sciences, Cambridge MA USA, 2Harvard University, Center for Nanoscale Systems, Cambridge MA USA
dcb@seas.harvard.edu

The analysis of glass samples with high resolution electron microscopy is challenging, because these techniques are not ideal for imaging amorphous samples. Visualizing the nanoscale elemental distribution and aggregation within glasses can lead to better modeling and understanding of their thermal, electrical, and mechanical properties. Advances in glass technology have lead to “gorilla glass”, the practically indestructible material found in the screens of many mobile devices. Understanding the nanoscale properties of materials like “gorilla glass” can lead to the development of new even stronger materials.

In the local electrode atom probe (3D LEAP) traditionally amorphous glasses have been difficult to run, with a particularly high failure rate. Until now the predominant approach for the preparation of such samples has been FIB liftout from a bulk specimen, which is time consuming and costly. We have developed technique for drawing glass rods or capillaries into sub 100 nm atom probe tips (Fig.1). This rapid and low expense technique means that even though the possible failure rate in the atom probe may be high, a significant amount of time is not wasted with sample preparation.

We found that after coating the glass nanotips with metal we were able to use voltage mode on the atom probe to successfully characterize a glass specimen (Fig . 2) which shows a very well defined Boron concentration surface . Further development of this technique could lead to new understandings of the structure property relationships in glasses and provide new pathways for studying materials doped or encapsulated within glasses.


The authors greatfully Acknowledge the support of the National Science Foundation through  NSF MRI:1040243

D.C. Bell. gratefully acknowledges funding through the National Science Foundation award  (NSF DMR-1108382)

Fig. 1: SEM Image of the rapid preparation Glass APT nanotip, showing diameter less than 100nm  (Inset) Glass nanotip mounted  in the analysis chamber

Fig. 2: APT Reconstruction of Glass nanotip showing an Isosurface of Boron concentration in the glass sample

Type of presentation: Oral

IT-17-O-3072 An atom probe insight on corrosion of stainless steels

La Fontaine A. J.1, 2, Cairney J.1, 2
1School of Aerospace, Mechanical, Mechatronic Engineering, The University of Sydney, NSW 2006, Australia, 2Australian Centre for Microscopy and Microanalysis, The University of Sydney, NSW 2006, Australia
alex.lafontaine@sydney.edu.au

For high temperature applications, such as in new-generation energy technologies, high-performance stainless steels offer an attractive combination of economy and mechanical / corrosion properties. For example, concentrated solar power (CSP) requires cost-effective and corrosion resistant materials that can operate for extended periods at high temperatures and withstand thermal cycling between 900°C and room temperature.

Stainless steels develop a passive layer protecting the steel from detrimental corrosion. When used in extreme conditions at high temperature in ultra-corrosive atmosphere, with thermal cycling or high pressure, this passive layer can be destroyed and leave the steel exposed to corrosion resulting in material failure.

The analysis of the chemical composition and microstructure of oxide layers using traditional analytical spectroscopy techniques has some limitation due to either limited lateral resolution or mass resolution. Atom probe tomography (APT), however, is a unique 3D technique allowing for atomic scale chemical characterization.

In the last decade, local electrode atom probe (LEAP) was applied to study oxide layers. For example, surface oxide layers in stainless steels or nickel-based alloys [1-3] were characterized by using laser-pulsed LEAP. The current work demonstrates our efforts in studying by laser-pulsed LEAP the complexity of oxides layers in intergranular corrosion cracks formed in a commercial austenitic stainless steel during thermal cycling.

References

[1] Lozano-Perez S, Saxey DW, Yamada T, Terachi T. Scripta Materialia 2010;62:855.
[2] Kruska K, Lozano-Perez S, Saxey DW, Terachi T, Yamada T, Smith GDW. Corrosion Science 2012;63:225.
[3] Baik S-I, Yin X, Seidman DN. Scripta Materialia 2013;68:909.


This work was made possible by an ARENA PhD scholarship. The authors acknowledge scientific and technical input and support from the Australian Microscopy & Microanalysis Research Facility (AMMRF) at The University of Sydney as well as CSIRO Energy Centre.

Fig. 1: (a) Photo of the concentrated solar power (CSP) plant used to heat treat the samples (b) Z-contrast back-scattered electron image of cross section of the sample (c) EBSD pattern quality map of IGC area with b.c.c. phase highlited in red, EDS O-K map - EDS Cr-K maps, atom probe volume of grain boundaries in chromite (Cr,O and Fe)

Type of presentation: Oral

IT-17-O-3061 3D Field Ion Microscopy for Characterization of Atomic Scale Radiation Damage in Fusion Related Materials

Dagan M.1, Roberts S. G.1, Gault B.1, Bagot P.1, Moody M. P.1
1Department of Materials, University of Oxford, Parks Road, Oxford OX1 3PH, United Kingdom
michal.dagan@materials.ox.ac.uk

Prospective materials for plasma facing components of fusion reactors must withstand conditions of high temperature and radiation dose. Tungsten is a leading candidate for this role. Hence, an extensive atomic-scale investigation of tungsten behavior under extreme conditions is critical. However the detection of nanoscale damage to the material’s atomic lattice is particularly challenging even for advanced microscopy techniques. The formation of dislocations, nano-voids, self-interstitials and clustering effects requires the ability to directly view the arrangements of individual atoms on the crystal lattice.
Here we present a technique for the atomic scale study of tungsten utilizing field ion microscopy (FIM). Conventionally, FIM is a 2D imaging technique: only the surface of the needle-shaped specimen is imaged by evaporating image-gas atoms from atomic sites on the surface. However, by increasing the voltage applied on the tip during imaging, it is possible to field evaporate constituent surface atoms from the specimen, in a similar manner to atom probe tomography (APT). The result is a series of highly resolved 2D FIM images that can be tomographically stacked to retrieve a 3D reconstruction of the sample. FIM-based 3D reconstruction has the potential to significantly exceed the resolution of APT analysis as it is not subject to the same detection efficiency limitations.
In this work a new 3D atom-by-atom reconstruction procedure for FIM is described. A significant challenge is automating the identification of atoms within an image. Furthermore, every atom can potentially appear across hundreds of FIM images before it is evaporated. Hence, the progress of each atom must also be automatically tracked throughout the sequence. An automated identification approach has been developed based on clustering together high intensity pixels within each image and across the sequence. The results of this clustering algorithm, applied to identification of atoms within two consecutive (244) planes in a tungsten sample are presented in Figure 1. All atoms are successfully identified and represented by different colors. A smearing effect in the direction along the depth of the sample is the result of different evaporation rates of individual atoms.
In complementary work to study radiation damage, tungsten samples have been ion-irradiated. Figure 2 demonstrates the increased spatial resolution offered by FIM in comparison to APT. A radiation induced dislocation is clearly imaged using FIM while the APT reconstruction of tungsten atoms does not directly indicate its presence. Interestingly, carbon impurities segregate to this region and indirectly reveal the dislocation in the APT data. However, to image single vacancies only FIM has the necessary atomic resolution.


Fig. 1: a: 150 stacked FIM images of the evaporation of two (244) planes in a tungsten sample. The colors represent intensity ranges of the pixels. b: Atoms are identified automatically by clustering together high intensity pixels. Each atom appears in a different color and is smeared in proportion to its evaporation rate.

Fig. 2: a-c: FIM and APT of irradiated tungsten. a: FIM image of a dislocation. b(c): APT analysis of tungsten(carbon) atoms. Tungsten reconstruction shows no sign of dislocations however carbon segregation reveals their presence. d-f:lower dose irradiated tungsten. d:FIM of vacancies. e(f): tungsten(carbon) APT reconstructions, no indication of vacancies.

Type of presentation: Oral

IT-17-O-3384 Nucleation and lateral growth of NiSi phase

El Kousseifi M.1, Hoummada K.1, Epicier T.2, Panciera F.1, Mangelinck D.1
1Aix-Marseille Université, CNRS, IM2NP, Case 142, 13397 Marseille Cedex 20, France, 2Université de Lyon, MATEIS, umr 5510, Bât. B. Pascal, INSA de Lyon, F-69621
mike.elkousseifi@im2np.fr

The Ni-based self-aligned silicide is widely used as contacts and interconnections in very large-scale integrated circuits [1]. They are obtained by solid state reaction between Ni thin film and Si substrate. Therefore, the fundamental mechanisms related to their formation, including the first stages of the nucleation, phase formation sequence, the growth kinetics, and the microstructures of the silicide, are of great interest for applications. NiSi is the desired phase in the Ni silicide sequence as the contact material in advanced integrated circuits [2]. However, a major disadvantage of NiSi is its degradation at high temperature. The addition of percent-wise Pt to Ni film increases significantly the stability of NiSi on Si substrates [3].
In this work, in situ-XRD annealing followed by atom probe tomography (APT) and in-situ transmission electron microscopy (TEM) analysis were used to study the reaction between 10 nm Ni (10% Pt) alloy film and Si(100) substrate. Isothermal annealing in in-situ XRD at different temperatures (200°C, 215°C and 230°C) have shown a nucleation-controlled behavior for NiSi growth at the epitaxial θ-Ni2Si/Si interface in contrast to the diffusion-controlled growth usually reported for NiSi [4]. TEM measurements have provided information about the NiSi nuclei shape and their distribution in the sample (Figure 1(a)) and additional information about the growth kinetics, while APT analyses were used to determine the 3D distribution of Ni, Si, and Pt atoms (Figure 1(b)). The Pt distribution was obtained in two cases: (1) in the θ-Ni2Si phase without the presence of NiSi nuclei and (2) inside the NiSi nuclei. A model for nucleation and lateral growth of NiSi at θ-Ni2Si/Si interface is proposed.

1. R. W. Mann, et al. IBM J. Res.Dev. 39, 403 (1995).
2. R. Mukai, et al., Thin Solid Films 270, 567 (1995).
3. D. Mangelinck, et al. Appl. Phys. Lett. 75 1736 (1999).
4. F. d’Heurle, et al. J. Appl. Phys. 55 4208(1984).


Thanks are due to the french METSA (www.metsa.fr) network for access to TEM at the CLYM platform (www.clym.fr), and to B. Van De Moortele (LGL, ENS-Lyon) for his assistance in sample preparation

Fig. 1: Figure 1: a) TEM image showing the NiSi nuclei inside the θ-Ni2Si phase. The cylinder represents approximately the region of APT analysis. b) APT concentration profile across θ-Ni2Si and NiSi phase.

Type of presentation: Poster

IT-17-P-2954 Atom Probe Workbench for Materials Science & Engineering

Ceguerra A. V.1, Stephenson L. T.1, Apperley M.2, Goscinski W. J.3, Ringer S. P.1
1ACMM, and School of AMME, The University of Sydney, NSW Australia, 2AMMRF, The University of Sydney, NSW Australia, 3Monash eResearch Centre, Monash University, VIC Australia
anna.ceguerra@sydney.edu.au

The Atom Probe Workbench (APW)[1] is a tool for analysing Atom Probe Microscopy data[2] for materials science. Developed under the National eResearch Collaboration Tools and Resources (NeCTAR)[3] Characterisation Virtual Laboratory (CVL)[4] project, the APW collates existing atom probe tools[5] within an Australian NeCTAR cloud infrastructure. APW is deployed on top of the CVL fabric, a software infrastructure common to all CVL application drivers. Features of the APW include those that are inherent to open source programs utilised as part of the workbench, which are Galaxy workflow engine[6] and MyTardis[7]. Additional developed features of the APW include usage tracking & reports, citation reporting[8]. Initial reports from users during development show the potential of APW to be used by researchers around Australia and worldwide.

1 A. Ceguerra, P. Liddicoat, M. Apperley, W. Goscinski, and S. Ringer, Atom Probe Workbench version 1.0.0: An Australian cloud-based platform for the computational analysis of data from an Atom Probe Microscope (APM), used for chemical and 3D structural materials characterisation at the atomic scale. (2014) 10.4227/11/53014684A67AC.

2 B. Gault, M.P. Moody, J.M. Cairney, and S.P. Ringer, Atom probe microscopy (Springer, 2012) 9781461434351.

3 NeCTAR, National eResearch Collaboration Tools and Resources, http://www.nectar.org.au (18 March 2014).

4 CVL, NeCTAR Characterisation Virtual Laboratory, http://www.massive.org.au/cvl (18 March 2014).

5 A.V. Ceguerra, A.J. Breen, L.T. Stephenson, P.J. Felfer, V.J. Araullo-Peters, P.V. Liddicoat, X. Cui, L. Yao, D. Haley, M.P. Moody, B. Gault, J.M. Cairney, and S.P. Ringer, The rise of computational techniques in atom probe microscopy, Current Opinion in Solid State and Materials Science 17, 224 (2013) 10.1016/j.cossms.2013.09.006.

6 D. Blankenberg, G. Kuster, N. Coraor, G. Ananda, R. Lazarus, M. Mangan, A. Nekrutenko, and J. Taylor, Galaxy: A web-based genome analysis tool for experimentalists, Current Protocols in Molecular Biology Chapter 19 (2001) 10.1002/0471142727.mb1910s89.

7 S. Androulakis, J. Schmidberger, M.A. Bate, R. DeGori, A. Beitz, C. Keong, B. Cameron, S. McGowan, C.J. Porter, A. Harrison, J. Hunter, J.L. Martin, B. Kobe, R.C.J. Dobson, M.W. Parker, J.C. Whisstock, J. Gray, A. Treloar, D. Groenewegen, N. Dickson, and A.M. Buckle, Federated repositories of X-ray diffraction images., Acta crystallographica. Section D, Biological crystallography D64, 810 (2008) 10.1107/S0907444908015540.

8 A.V. Ceguerra, P.V. Liddicoat, W.J. Goscinski, S. Androulakis, and S.P. Ringer, A Tool for Scientific Provenance of Data and Software (2013) Computer and Information Technology Conference, Sydney, Australia 10.1109/CSE.2013.89.


We acknowledge the support of NeCTAR, AMMRF, The University of Sydney, Monash University, Intersect Australia.

For the full text of the acknowledgements, please see: http://dx.doi.org/10.4227/11/53014684A67AC

Type of presentation: Poster

IT-17-P-2966 Crystallographic calibration of Si-based atom probe reconstructions for enhanced short-range ordering information.

Breen A. J.1, 2, Ceguerra A. V.1, 2, Araullo-Peters V. J.1, 2, Moody M. P.3, Ringer S. P.1, 2
1Australian Centre for Microscopy and Microanalysis, The University of Sydney, NSW 2006, Australia, 2School of Aerospace, Mechanical and Mechatronic Engineering, The University of Sydney, NSW 2006, Australia, 3Department of materials, The University of Oxford, Parks Road, OX13PH, Oxford, UK
andrew.breen@sydney.edu.au

Atom probe tomography (APT) enables the position and chemical identity of millions of atoms to be reconstructed in 3D with sub-nm precision. However, the spatial resolution is still generally not high enough to unequivocally determine the lattice positioning of atoms in crystalline materials. This presents a significant roadblock in understanding structure-property relationships at the atomic level. An important example is the need to more accurately characterise the dopant positioning within the ultra-shallow junctions of the latest generation of transistor devices that are now usually only several nm deep. Subtle differences in dopant positioning can have significant influence on device performance and this must be controlled more accurately if continual size reductions are to occur.  

In this study, we have developed methods to study the short-range ordering of dopants within silicon in unprecedented detail using APT. Latent crystallographic structure can be detected within the reconstructions with the help of newly developed crystallographic mapping tools. After doing this, it is apparent that the detected crystal structure is slightly different to the theoretical structure in size and shape, even after careful calibration using the method described by Gault et al. (Gault, et al., 2009). This is most likely due to assumptions used in the initial reconstruction algorithm. Further crystallographic calibration, using linear transformations of the reconstructed atom co-ordinates, is then used to perfectly restore the latent crystal structure. However, we have found that the spatial noise present is still high enough to be detrimental to the measured short-range ordering. Lattice rectification approaches (Moody, et al., 2011), where atoms are repositioned to the closest lattice site, are then used to restore this short-range order and the results are compared and discussed. This work provides a significant contribution to lattice based atom probe studies of doped silicon.

Gault, B., Moody, M.P., de Geuser, F., Tsafnat, G., La Fontaine, A., Stephenson, L.T., Haley, D. & Ringer, S.P. (2009). Advances in the calibration of atom probe tomographic reconstruction. Journal of Applied Physics 105(3).

Moody, M.P., Gault, B., Stephenson, L.T., Marceau, R.K.W., Powles, R.C., Ceguerra, A.V., Breen, A.J. & Ringer, S.P. (2011). Lattice Rectification in Atom Probe Tomography: Toward True Three-Dimensional Atomic Microscopy. Microscopy and Microanalysis 17(2), 226-239.


The authors acknowledge scientific and technical input from the AMMRF node at The University of Sydney, particularly Baptiste Gault, Leigh Stephensen, Takanori Sato and Julie Cairney. The authors are also grateful to the ARC for providing funding. We also thank the ANFF at the University of NSW, particularly Joanna Szymanska, for Bosch processing of Si wafers. 

Fig. 1: The latent crystallographic information within a Sb doped Si reconstruction. (a) The tomographic atom probe reconstruction. A thin red slice has been cropped out for crystallographic analysis. (b) 2D density map of Si. (C) 1D spatial distribution maps (SDM) of detected planes within the reconstruction (d) 2D SDM of {001} planes.

Type of presentation: Poster

IT-17-P-3236 Improving Yield and Data Quality in Atom Probe Tomography

Larson D. J.1, Ulfig R. M.1, Prosa T. J.1, Lawrence D. F.1, Martin I. Y.1, Giddings A. D.1, Olson D. P.1, Kelly T. F.1
1CAMECA Instruments Inc., 5500 Nobel Drive, Madison, WI 53711 USA
rulfig@hotmail.com

The electric fields used in atom probe tomography (APT) generate stresses normal to the specimen axis proportional to the electric field squared [1]. As such, specimen fracture has long been a serious problem for APT [2-3]. There are many methods to improve yield, often with a trade-off required in experimental quality or convenience in some other area. Methods to improve yield include: 1) decrease ion detection rate, 2) increase specimen temperature, 3) use laser pulsing, 4) increase laser pulse energy, 5) increase background chamber pressure, 6) change feature orientation [4], 7) coat a sharpened specimen with a thin layer of material [3,6], and 8) slightly anneal the specimen. Figure 1 schematically illustrates several of these methods to improve APT specimen yield. The current work explores the use of thin coatings to modify the thermal and/or optical properties of 302 stainless steel and a Si/SiO2(10nm)/Si/Ni test structure in attempts to improve yield.

Although method number 3+4 above works well to improve yield, substantial increases in laser energy are often limited in materials with poor thermal diffusivity due to a degradation in data quality (e.g., long “thermal tails” in the mass spectrum, non-uniform detector hit-maps, etc.) [7,8]. An example of this effect is shown in the non-uniform detector hitmaps in a poor thermal diffusivity material (302 stainless steel). Figure 2 shows the effects of increasing laser pulse energy and suggests an increasing trend in azimuthal asymmetry (note laser incidence is lower left). Figure 3 presents the mass resolving power (MRP) at tenth maximum, uniformity in the form of XY hitmap signal correlation (lower values are desirable) [6], and spectral background as a function of laser energy.

In order to investigate the yield improvement of coatings, we have developed a test structure of uncoated Si/SiO2(10nm)/Si/Ni on which we can reach data collection conditions where this sample both does and does not yield. These data have been reproduced by successful data collection from the entire structure multiple times. Statistically, yield changes are difficult to prove unequivocally; however we have good evidence for an improvement for this oxide-based structure with thin metal coatings in Figure 4.

1. D. G. Brandon in, Field Ion Microscopy, eds. J. Hren, S. Ranganathan (Plenum, 1968) p.64.
2. T. J. Wilkes et al., J. Phys. D: Appl Phys. 5 (1972) p.2226.
3. S. Kölling and W. Vandervorst, Ultramicroscopy 109 (2009) p.486.
4. D. Lawrence et al., Microsc. Microanal. 14(S2) (2008) p.1004.
5. G. L. Kellogg, J. Appl. Phys. 53(9) (1982) p.6383.
6. D. J. Larson et al., Microsc. Microanal. 20(S2) (2014) in press.
7. J. H. Bunton et al., Microsc. Microanal. 13 (2007) p.418.
8. B. Gault et al., Ultramicroscopy 110 (2010) p.1215.


The authors would like to thank the team at CAMECA Instruments Inc. for continued software and hardware improvements to the LEAP®.

Fig. 1: 1)Methods to improve APT yield. 2)Effects of increasing laser pulse energy suggesting a trend in azimuthal asymmetry laser incidence is LL). 3)Mass resolving power at tenth maximum, uniformity in the form of XY hitmap signal correlation (lower=better). 4)Yield differences on a difficult to run specimen as a funciton of detection rate and coating.

Type of presentation: Poster

IT-17-P-3469 The effects of laser wavelength and pulse energy on the measured oxide compositions by atom-probe tomography

Kruska K.1,2, Lozano-Perez S.1, Schreiber D. K.3
1Department of Materials, University of Oxford, Parks Road, Oxford, OX1 3PH, United Kingdom, 2Fundamental and Computational Science Directorate, Pacific Northwest National Laboratory, Richland, WA, United States, 3Energy and Environment Directorate, Pacific Northwest National Laboratory, Richland, WA, United States
karen.kruska@pnnl.gov

Atom-probe tomography (APT) has been shown to be a useful tool to study environmental degradation, and in particular oxidation and corrosion, of steels and Ni-base alloys frequently used in the high temperature corrosive environments. In these corroded samples, APT is uniquely able to capture highly localized changes in composition with both high chemical sensitivity and near-atomic spatial resolution in 3D. In principle, it is possible to correlate the local O concentration with specific oxide phases. In practice, however, the measured O concentration is erroneously O deficient. This “loss” of oxygen atoms in relevant oxides is neither well documented nor understood.

In this study bulk Fe and Ni oxides – NiO, FeO and NiFe2O4 – have been systematically analyzed in two complementary APT systems with a green (λ=532 nm) and a UV (λ=355 nm) laser, respectively, to better understand the measured O deficiency. These oxides were selected primarily for their relevance to corroded microstructures. The laser pulse energies were varied and repeated in increasing and decreasing series to eliminate additional effects of increasing tip radius during data collection. The measured composition at a given pulse energy was found to be consistent, repeatable, and independent of tip radius. The consistency of measured compositions was also remarkable between multiple tips.

In all cases the measured O concentration increased with decreasing pulse energies, which is consistent with previous studies on other oxide and nitride materials. Figures 1a and b show this with both the green and the UV wavelength lasers. The ratio of O:O2 ions within the measured O was also studied (see Figure 1b and 2b),and was found to increase with increasing O concentration, suggesting that more molecular O are lost at higher laser energies, likely through sublimation of neutral molecular O2.

Figure 2 shows data from the spinel oxide, NiFe2O4, acquired with the UV laser. Although the O content increased and the overall metal content decreased with decreasing laser pulse energy, irregularities were observed in the Fe:Ni ratio (see Figure 2c). While it remained constant at high laser pulse energies it starts to drop at a laser pulse energy of 1 pJ. It appears that Fe cations are lost at low laser pulse energies in this case.

These results show that the evaporation behavior in the studied oxides is strongly cation dependant. It appears that the O:O2 ratio cannot be used as an indicator for the accuracy of the composition. A more promising indicator may be the cation ratio, especially in mixed spinel oxides.


A portion of the research was performed using EMSL, a national scientific user facility sponsored by the Department of Energy's (DOE) Office of Biological and Environmental Research and located at Pacific Northwest National Laboratory. DKS was funded by the US DOE Office of Basic Energy Science.

Fig. 1: a) O concentration evolution depending on the laser power in NiO and FeO for both green and UV laser systems. b) O concentration in dependence of O:O2 ratio in both green and blue laser systems.

Fig. 2: a) O concentration evolution depending on the laser power in NiFe2O4 for UV laser system. b) O concentration in dependence of O:O2 ratio UV laser system. c) Cation ratio normalised to the nominal Fe and Ni contents in NiFe2O4 in dependence of laser pulse energy.

Type of presentation: Poster

IT-17-P-5928 Quantitative evaluation of C, Mn and Si in martensite steels by atom probe tomography, electron microscopy and X-ray diffraction

M. Kozuka 1 S. Otani 1 Y. Aruga 1
Materials Research Laboratory, Kobe Steel, Ltd., 1-5-5 Takatsukadai, Nishi-ku, Kobe, Japan 1
kozuka.masaya@kobelco.com

Martensite steels have been widely used due to high strength, which strongly depends on the microstructural factors such as solute carbon content in martensite phase. Due to the high spatial resolution and capability for analyses of light elements, atom probe tomography (APT) is one of the promising methods to quantitatively evaluate the solute carbon content of martensite phases. However, it is well known that the solute carbon content evaluated by APT apparently depends on the measurement parameters such as specimen temperature and peak identification of the mass spectrum [1]. Although Miyamoto et al. recently proposed that the dependence of apparent solute carbon content in Fe-C binary alloys on the specimen temperature is due to the detection loss of iron ions [2], the quantitative evaluation of the solute carbon content still has remained issues. For example, micro-alloying elements will affect the evaporation behavior and may change the APT analysis results. In this study, we conducted APT measurements of Fe-C, Fe-C-Mn and Fe-C-Si martensite steels with varied specimen temperature and pulse fraction. The effects of a micro-alloying element on evaluation of the solute carbon content were investigated in terms of measurement parameter and interpretation of mass spectrum. The steels were austenitized at 1203 K for 360 s followed by water quenching. Fig. 1 is a secondary electron image of the as-quenched Fe-1.0C-1.5Si steel etched by picral, showing that the alloy has martensitic structure and no mm-sized carbon segregation. Fig. 2 shows APT maps of carbon and silicon in the same steel. Although nm-sized segregation is confirmed in the carbon atom map (Fig. 2 (a)), the alloy seems to be suitable for the quantitativity check of the APT measurement because the carbon segregation looks uniform in the steel. In the presentation, quantitative evaluation results of electron probe microanalyser and X-ray diffraction are also shown to discuss their complementary use.

[1] J. Takahashi et al., Ultramicroscopy 111 (2011) 1233-1238.

[2] G. Miyamoto et al., Scripta 67 (2012) 999-1002.


We would like to thank Dr. T. Murakami (Kobe Steel, Ltd) for beneficial discussion.

Fig. 1: A secondary electron image of the Fe-1.0C-1.5Si steel.

Fig. 2: APT elemental maps of (a) C and (b) Si in the Fe-1.0C-1.5Si steel.

MS-1. Nanoobjects and engineered nanostructures, catalytic materials

Type of presentation: Invited

MS-1-IN-1523 Atomic Ordering and Phase Separation in Magnetic Alloy Nanoparticles

Sato K.1, Konno T. J.1
1Institute for Materials Research, Tohoku University, Sendai, Japan
ksato@imr.tohoku.ac.jp

Recent demands for ultrahigh density magnetic storage technology require the development of novel recording media with higher magnetocrystalline anisotropy energy (MAE), with the aim of increasing storage density and reducing a recording noise. For such a purpose, L10-type CoPt ordered alloy nanoparticles are one of the candidate materials: the hard magnetic properties of this alloy can be attributed to the tetragonal ordered structure with a high MAE. Therefore, the atomic ordering and the stability of the ordered phase are the key issues for the magnetic properties. We hence intend to examine kinetics of ordering in CoPt nanoparticles. On the other hand, structure and properties of alloy nanoparticles can be tuned as well by taking advantage of phase separation. For this purpose, we focused on immiscible Co-Au system. Bimetallic nanoparticles were fabricated by sequential electron-beam depositions of Pt (or Au), Co and Al2O3 onto NaCl(001) substrates kept at 520-653 K. For CoPt nanoparticles, we carried out post-deposition annealing for promoting atomic ordering with different cooling rates. The structure and morphology of the nanoparticles were characterized using TITAN80-300, JEM-3011, and ARM200F (S)TEM. Figures 1 (a)-(c) and (a’)-(c’) show CS-corrected HRTEM images of the CoPt nanoparticles with sizes of ~5 nm and ~3 nm in diameter, respectively. The annealing conditions are as follows: (a, a’) 873 K-1h, slow cooling (1.5 K/min), (b, b’) 873 K-1h, rapid cooling (110 K/min), and (c, c’) 973 K-1h, rapid cooling. For 5 nm-sized particles (left panel), clear (110) atomic planes of the L10-ordered structure can be seen. In contrast, size dependence of the atomic ordering was found in the specimens followed by rapid cooling: the disordered phase was observed in nanoparticles smaller than 3 nm in diameter. For example, a nanoparticle shown in Fig. 1(b’) is the disordered phase, characterized by crossed {200} atomic planes. The population of the disordered particles was found to be 14% (Fig.1(e)). Thus, this study offers a kinetic explanation for size-dependent atomic ordering, in addition to the thermodynamic explanation of a reduced stability for the ordered phase in nanoparticles [1]. Figure 2 (a) shows an SAED pattern of Co-Au nanoparticles. The cube-on-cube orientation epitaxy is seen between Au and fcc-Co. An example of a HAADF-STEM image of a Co-Au particle is shown in Fig. 2(b). Formation of core-shell structure (Au-core) is clearly seen due to the large atomic number difference. As the particle size reduces (<8 nm), nanoparticles with Au-shell are also formed as shown in Fig. 2(c). It is expected that smaller sized Au-shell particles have a potential to nanocatalyst applications. [1] K. Sato et al. Philos. Mag. Lett. 92 (2012) 408.


This study was partially supported by the Grant-in-Aid for Scientific Research from the Ministry of Education, Culture, Sports, Science, and Technology, Japan.

Fig. 1: CS-corrected HRTEM images and the corresponding FFT patterns of the CoPt nanoparticles (average alloy composition: Co-61at%Pt): (a-c) ~5 nm in diameter, (a’-c’) ~3 nm in diameter. (d) a structure model (truncated octahedron) and a simulated image, (e) particle size range, where the disordered nanoparticles are observed, is marked in the histogram.

Fig. 2: (a) SAED pattern of Co-Au nanoparticles (average alloy composition: Co-46at%Au). (b, c) Z-contrast images by HAADF-STEM for Co-Au nanoparticles with core-shell contrasts: (b) a Au-core particle (D~13 nm), (c) a Au-shell particle (D~7 nm). It is presumed that the lower surface tension of Au than that of Co is responsible for the Au-shell formation.

Type of presentation: Invited

MS-1-IN-5776 Morphology and inner structure of anisotropic nanoparticles studied by advanced TEM and FIB techniques

Spiecker E.1, Butz B.1, Winter B.1, Niekiel F.1, Vieweg B.1
1Center for Nanoanalysis and Electron Microscopy (CENEM), Universität Erlangen-Nürnberg, Erlangen, German
Erdmann.Spiecker@ww.uni-erlangen.de

Extensive research in nanoscience is currently devoted to the synthesis and characterization of complex and anisotropic nanoparticles (NP), with particular focus on low-symmetry NP, patchy structures as well as hybrid particle assemblies. Anisotropy and increased complexity provide additional degrees of freedom for tailoring the physical and chemical properties, making such NP ideal candidates for enhanced chemical, catalytic, and optical applications. In this contribution, microscopic studies of anisotropic and complex NP are presented which illustrate the importance of combining advanced TEM and FIB techniques for comprehensive characterization.

Ge/Si patchy NP formed by gas phase synthesis in a two-stage hot wall reactor [1,2] have been studied by combining HAADF-STEM, STEM-EELS and HRTEM, as illustrated in Fig. 1. The microscopic investigations focused on two key questions, namely the accommodation of misfit (~ 4 %) by defect formation and possible interdiffusion, addressed by HRTEM and STEM/EELS, respectively. Another experimental aspect is the oxidation of the particle surface which could be suppressed by transfer via glove box and vacuum transfer holder.

PbSe quantum dots grown on single wall carbon nanotube (SWCNT) bundles for optoelectronic applications [3] have been studied by combining electron tomography and HRTEM at 80 kV. During wet-chemical synthesis the PbSe NP grow around the SWCNT bundles as revealed by the tomogram in Fig. 2a. In combination with HRTEM contributing crystallographic information (Fig. 2b) a 3D atomistic model is derived (Fig. 2c). By comparison of experimental and simulated HRTEM images the 3D model of the PbSe NP can be further refined.

Ag nanorods formed by wet chemical synthesis [4] are used to illustrate a new FIB method [5] which enables site and orientation specific cross-sectioning of anisotropic NP (Fig. 3). By employing a shadow geometry (Fig. 3b) in which a Si substrate protects the NP from the Ga+ ion beam no protection layer has to be deposited on the NP. Inspection of the NP during milling enables precise positioning of the final cross section. In the case of the Ag nanorod, HRTEM of the final cross section nicely reveals the characteristic five-fold twin structure with additional defects indicating partial strain relaxation (Fig. 3c).

[1] Körmer et al., Crystal Growth & Design 12 (2012), 1330.
[2] Mehringer et al., Journal of Aerosol Science 67 (2013), 119.
[3] Schornbaum et al., Chemistry of Materials, 25 (2013), 2663.
[4] Damm et al., Small 7 (2011), 147.
[5] Vieweg et al., Ultramicroscopy, 113 (2012), 165.
[6] http://www.gmp.ww.uni-erlangen.de/nanoSCULPT.php


The authors gratefully acknowledge collaboration with the Institute for Particle Technology (Prof. Peukert) and the Nanomaterials for Optoelectronics Group (Prof. Zaumseil). They thank Prof. Bitzek for providing the software nanoSCULPT [6] used for constructing the atomistic model of the PbSe NP. Financial support by the DFG through the projects EXC315 and GRK1161 is gratefully acknowledged.

Fig. 1: TEM study of Ge/Si patchy NP produced by gas phase synthesis [2]. a) HAADF-STEM image of aggregated NP with Si cores (dark) and Ge islands (bright). b) Concentration profiles across an interface (line in a)) from EELS line scan. c) Map of (111) lattice plane spacing derived from a HRTEM image (not shown) by geometric phase analysis.

Fig. 2: Morphology and atomic structure of PbSe NP grown on SWCNT bundle [3]. a) Tomogram of PbSe NP. b) HRTEM image of the same NP in [011] zone axis. c) Atomistic model obtained by filling the 3D volume with the PbSe structure in correct crystallographic orientation derived from HRTEM. d) HRTEM image simulation based on atomistic model.

Fig. 3: FIB technique for site and orientation specific sectioning of anisotropic NP [5]. a) TEM image of Ag nanorod with desired cross section indicated by black rectangle. b) Shadow-FIB geometry used for cross-sectioning without protection layer. c) HRTEM image of nanorod cross section showing the characteristic five-fold twin structure.

Type of presentation: Invited

MS-1-IN-5815 STEM Investigation of Nanostructured, Ceria-based, Surface Phases: Novel Catalysts with Outstanding Redox Properties

Calvino J. J.1, Sánchez-Gil J. J.1, Tinoco M.1, Arias D. C.1, Yeste M. P.1, Muñoz M. A.1, Hungría A. B.1, Hernández-Garrido J. C.1, Pérez-Omil J. A.1, Cauqui M. A.1, Blanco G.1, Trasobares S.1, Bayle-Guillemaud P.2, López-Haro M.2, Carlsson A.3
1Departamento de Ciencia de los Materiales e Ingeniería Metalúrgica y Química Inorgánica, Facultad de Ciencias, Universidad de Cádiz, Campus Río San Pedro, 11510-Puerto Real, Cadiz, Spain., 2Univ. Grenoble Alpes, F-38000 Grenoble, France, CEA-INAC/UJF-Grenoble 1 UMR-E, SP2M, LEMMA, Minatec Grenoble, F-38054, France, 3FEI Company, FEI Europe B.V., Achtseweg Noord 5, 5651 GG Eindhoven, Netherlands
jose.calvino@uca.es

Access to Rare Earth Elements (REE) and Platinum Group Metals (PGMs) is nowadays considered a major limiting factor for the development of Green Technologies. Aiming to contribute to the improvement in the efficient use of such strategic materials in the field of Heterogeneous Catalysis, novel nanocatalysts, based on ceria (CeO2) are being investigated in our lab, featuring the following characteristics: low lanthanide contents and a noble-metal free formulation.

The strategy to synthesize the new materials consists in structuring the ceria component as nanometer thick surface layers coherently grown onto the surface of a carrier oxide (ZrO2, YSZ, MgO). The analysis of the Redox properties of this new type of catalysts, which play in fact a key role in their catalytic performance in a variety of reactions, indicates a large improvement with respect to materials based on bulk ceria [1]. A better performance is observed both in H2-reducibility at low temperatures as well as in the stability of the redox response against aging treatments at very high temperatures.

STEM analysis of these new materials has been key both to check the success in the nanostructuration targets proposed for their synthesis and, what´s more important, to understand the behavior observed at macroscopic level in their redox properties. Thus, by combining experimental HAADF-STEM and HREM images recorded on these materials, Figure 1, with simulated ones, Figure 2, it has been possible to detect the presence of a variety of nanostructures ranging from isolated, atom-like, Ce-species, up to 3D, well-faceted, nanoparticles, going through patch or raft-like 2D nano-objects. Such exotic, highly dispersed, nanostructures pose challenges not only to their detection but also to their ultimate 3D-characterization by Electron-tomography and to the analysis of their chemical nature by STEM-EELS or STEM-XEDS. All these aspects will be covered during the lecture.

References

[1] D.C. Arias et al., J. Mat. Chem. A., 1 (2013) 4836


Financial support from MICINN/FEDER IMAGINE (CSD2009-00013) and FP7-I3 ESTEEM 2 (Grant Agreement 312483) Projects is gratefully acknowledged. Authors would like to thank the nanocharacterisation platform (PFNC) of CEA-Grenoble for access to their FEI-TitanUltimate.

Fig. 1: (a) HAADF-STEM image of a 2-atomic planes thick CeOx surface structure recorded on a 4% mol. CeO2/MgO catalyst; (b) HREM view of the same catalyst showing at the same time both highly dispersed, atom-like, Ce-species and nanosized CeOx-rafts.

Fig. 2: Structural models (a,c) and HREM simulated images (b,d) of isolated Ce-species (top row) and 1-atom thick CeOx nano-rafts (lower row) supported on MgO.

Type of presentation: Oral

MS-1-O-1571 Influence of Strain State on the Formation of Short-Period InGaN/GaN Nanowire Superlattice by Electron Energy-Loss Spectroscopy

Woo S. Y.1, Gauquelin N.1,2, Kociak M.3, Nguyen H. P.4, Mi Z.4, Botton G. A.1
1Department of Materials Science & Engineering, Brockhouse Institute for Materials Research, and Canadian Centre for Electron Microscopy, McMaster University, Hamilton, ON, Canada, 2Now at EMAT, University of Antwerp, Antwerpen, Belgium, 3Laboratoire de Physique des Solides, Université Paris-Sud XI, Orsay, France, 4Department of Electrical & Computer Engineering, McGill University, Montreal, QC, Canada
woosy@mcmaster.ca

Ternary InGaN alloys have been sought-after for various optoelectronic device applications, including their prospect as highly efficient phosphor-free white light-emitting diodes (LEDs). The growth of high quality InGaN epilayers over the entire compositional range, however, faces a few obstacles impeding the realization of their full visible wavelength range tunability. The large InN/GaN lattice mismatch can induce a high density of threading dislocations, and the InGaN miscibility gap leads to inhomogeneity and difficult indium incorporation. Therefore the determination of composition, in particular quantitative elemental mapping at high spatial resolution, is imperative to further understanding the formation of III-N heterostructures.
The growth of high quality III-N heteroepitaxy in a nanowire (NW) geometry is a promising alternative, as shown in the recently developed InGaN/GaN quantum dot (QD) superlattice towards controlled light emission across the entire visible spectrum [1]. In this work, multiple InGaN/GaN dot-in-a-wire nanostructures grown on Si(111) substrates by molecular beam epitaxy were characterized by aberration-corrected scanning transmission electron microscopy (STEM) to correlate their structural to optical and electrical properties. High-angle annular dark-field (HAADF) Z-contrast imaging showed that the 10 InGaN QDs are centrally confined within the active region, embedded between n- and p-doped GaN in the NW LED structure. Core-loss EELS spectrum imaging is used to evaluate the elemental distribution in the NW heterostructures. The In-content is quantified using a multiple linear least squares (MLLS) fitting routine, with combined internal and external reference spectra to fit the N K (399 eV) and In M4,5 (451 eV, in close proximity to the N K) edges in the spectrum image. A surface plot of the thickness-corrected In-map clearly illustrates the non-uniformity of InxGa1-xN composition between the 10 dots (Fig. 1), which has occurred systematically despite the constant growth conditions of the QDs. Geometric phase analysis (GPA) of corresponding atomic-resolution STEM images was used to extract the local strain components at the nanoscale. Along the growth direction (Fig. 2(c)), a direct correlation between a GaN barrier’s strain state and the amount of In incorporated into the subsequent QD can be deduced. Examining the strain distribution of the QDs aids to elucidate their formation as governed by the incorporation of In. In addition, effects of the varying composition on emission wavelength in single NWs using nm-resolved STEM-cathodoluminescence will also be shown.
[1] Nguyen et al., Nano Lett., 12(3), 1317-1323 (2012)


This work was supported by the Natural Sciences and Engineering Research Council of Canada (NSERC).

Fig. 1: (a) HAADF-STEM image of a NW studied using STEM-EELS spectrum imaging, followed by subsequent MLLS fitting using combined internal and external reference spectra. (b) Surface plot of the thickness-corrected In-content map, generated from normalizing the MLLS-fitted In-map with the N-map, showing InxGa1-xN composition ranging between x=0.12–0.38.

Fig. 2: (a) MLLS-fitted relative In-content map from EELS with internal references. (b) Corresponding HAADF-STEM image with its strain maps along the growth (c) and in-plane (d) directions, and dilatation matrix (e). The maps show that there is a direct correlation between the In-content as highlighted in (a) and the strain along the growth direction (c).

Type of presentation: Oral

MS-1-O-1786 Catalytically active structures in Au nanoparticulate catalysts studied by quantitative environmental TEM

Kuwauchi Y.1, Yoshida H.1, Takeda S.1
1The Institute of Scientific and Industrial Research, Osaka University
takeda@sanken.osaka-u.ac.jp

To validate the usefulness of in-situ environmental TEM (ETEM) in catalyst chemistry, there remain several issues to be addressed such as electron irradiation effects, heterogeneity of real catalysts, temperature and pressure gaps [1]. It is recently shown that these issues in ETEM observation can be settled in supported gold nanoparticulate catalysts (AuNP catalysts) and others by quantitative data analyses [2-4]. These quantitative analyses can confirm that an area of interest under atomic resolution ETEM observation acts as catalyst. Based on the analyses, we show some recent results that can only be derived by quantitative atomic resolution ETEM.
We used a prototype ETEM [1] that is equipped with a corrector for the spherical aberration of the objective lens and was operated at 80, 200 and 300 kV. The basic part of the prototype ETEM is commercially available as FEI Titan ETEM G2. The sample was Au/CeO2 powder that has exhibited high catalytic activity for the oxidation of CO even below room temperature. Details of the samples were already described before [2].
We have studied dynamic structures of AuNP catalysts. The observation proved that catalytically active AuNPs move reversibly and stepwise by approximately 0.09 nm on CeO2 support surface at room temperature and in a reaction environment (Fig. 1). The lateral displacements and rotations indicate that AuNPs are loosely bound to oxygen-terminated CeO2. The AuNPs are likely anchored to oxygen-deficient sites [5]. Observations indicate that the most probable activation sites in AuNP catalysts, which are the perimeter interfaces between an AuNP and a support, are not structurally rigid. It is also shown that the surfaces of AuNPs were structurally reconstructed under reaction conditions, via interactions with CO molecules [6]. CO molecules were observed on the surfaces of catalysts under reaction conditions (Fig. 2). We present more details on in-situ ETEM observation of the catalysts and others.

References
[1] S. Takeda and H. Yoshida, Microscopy. 62 (2013) 193-203.
[2] T. Uchiyama, H. Yoshida, Y. Kuwauchi, S. Ichikawa, S. Shimada, M. Haruta, and S. Takeda, Angew. Chem. Int. Ed. 50 (2011) 10157-10160.
[3] Y. Kuwauchi, H. Yoshida, T. Akita, M. Haruta, and S. Takeda, Angew. Chem. Int. Ed. 51 (2012) 7729-7733.
[4] H. Yoshida, K. Matsuura, Y. Kuwauchi, H. Kohno, S. Shimada, M. Haruta, and S. Takeda, Appl. Phys. Express 4 (2011) 065001/1-3.
[5] Y. Kuwauchi et al., Nano Lett. 13 (2013) 3073-3077.
[6] H. Yoshida, Y. Kuwauchi, J. R. Jinschek, K. Sun, S. Tanaka, M. Kohyama, S. Shimada, M. Haruta, S. Takeda, Science 335 (2012) 317-319.


This study was partially supported by a Grant-in-Aid for Scientific Research (A) , Grant No. 25246003 from the Ministry of Education, Culture, Sports, Science and Technology, Japan. This research was partly supported by the Management Expenses Grants for National Universities Corporations from the Ministry of Education, Culture, Sports, Science and Technology of Japan (MEXT).

Fig. 1: Stepwise displacement and rotation of a AuNP supported on CeO2 [5]. (a) In-situ observation. Observation time is indicated. (b) Instantaneous structural models that are depicted in lateral (b) and top (c) views.

Fig. 2: Glimpse of gas molecules (CO) on the reconstructed surface of a AuNP [6]. In reaction gas in (b), the {100} facet, indicated by rectangle is structurally reconstructed, while in vacuum in (a) the facet is unreconstructed. Faint image contrast on the reconstructed facet in (b) can be accounted by adsorbates (CO molecules).

Type of presentation: Oral

MS-1-O-1802 3D elemental mapping of the atoms in bimetallic nanocrystals

Goris B.1, De Backer A.1, Van Aert S.1, Van Tendeloo G.1, Liz-Marzan L. M.2, Bals S.1
1EMAT, University of Antwerp, Groenenborgerlaan 171, 2020 Antwerp, Belgium, 2BioNanoPlasmonics Laboratory, CIC biomaGUNE, Paseo de Miram´on 182, 20009 San Sebastian, Spain
bart.goris@uantwerpen.be

A three dimensional (3D) characterization of complex heterostructures is required for an optimized understanding of their properties. For example, it is known that the optical properties of bimetallic Au@Ag nanostructures are largely determined by the presence of certain surface facets and interfaces. Furthermore, effects such as alloying or intermixing of the atoms at the interfaces results in a shift of the plasmon resonances, urging the need for a thorough 3D study at the atomic scale. [1] Electron tomography is a technique to obtain 3D reconstructions based on a series of 2D projection images and recently, different approaches enable a 3D investigation at the atomic scale as well. [2-4]
Here, we apply a tomography approach which is based on compressive sensing to reconstruct the atomic lattice of Au@Ag bimetallic nanoparticles having important applications in the field plasmonics. In order to obtain a reliable reconstruction, 5 high resolution HAADF-STEM projection images are acquired along different orientations of the nanorod and used as an input for a tomographic reconstruction. More detailed information about the experimental set-up can be found in [5]. 3D visualizations of the results are illustrated in figure 1, where a visual distinction can be made between the Au core (yellow) and the surrounding Ag shell (blue). Figures 1a-c correspond to visualizations where the sample was tilted by 0º, 45º and 90º around the [010] axis, respectively. The resulting Fourier transforms correspond to the expected symmetry for a fcc crystal structure and the atomic lattice can clearly be recognized from the visualizations themselves.
Since the reconstruction is based on HAADF-STEM projection images, the intensity of the reconstructed atoms scales with their atomic weight. Therefore, Ag and Au atoms can be identified by analysing intensity profiles through the reconstruction. An example of such an intensity profile is presented in figure 2 which enables the labelling of each atom in the cross sections presented in figures 2b and 2c to be either Ag or Au. The result is shown in figures 2d and 2e yielding a correct indexing of the facets composing the interfaces. We conclude that the interface between the Au core and the Ag shell is sharp, without intermixing and mainly composed of {520} facets where bevels are observed at the <100> and the <110> directions. This type of detailed information is crucial to understand and optimize the physical properties of these materials. [6]
[1] M.B. Cortie and A.M. McDonagh, Chemical Reviews 11 (2011) 3713-3735
[2] S. Van Aert et al., Nature 470 (2011) 374-377
[3] M.C. Scott et al., Nature 483 (2012) 444-447
[4] B. Goris et al., Nature Materials 11 (2012) 930-935
[5] B. Goris et al., Nano Letters 13 (2013) 4236-4241


The authors acknowledge support from the European Research Council (ERC Grants # 24691-COUNTATOMS and #335078-COLOURATOMS) and from the Flemish Fund for Scientific Research.

Fig. 1: (a-c) 3D renderings of the reconstruction viewed along different directions where the sample was tilted around the [010] axis for 0º, 45º and 90º. The atoms in the Au core are rendered yellow whereas the surrounding Ag shell is shown in blue. The Fourier transforms of these projected views correspond to a fcc crystal lattice.

Fig. 2: (a) Three orthogonal slices through the reconstruction show the structure of the nanorod. (b,c) Detailed view of the slices through the reconstruction. An intensity profile is acquired along the direction indicated by the white rectangle in (b). (d,e) Slices corresponding to (b) and (c), in which each Au atom is indicated by a yellow circle.

Type of presentation: Oral

MS-1-O-2106 Using electron beam to investigate metal-carbon catalyst

Su D.1, Zhang B.1
1Shenyang National Laboratory for Materials Science, Institute of Metal Research, Chinese Academy of Sciences, Shenyang, China
dssu@imr.ac.cn

Carbon (e.g. carbon nanotube and graphene) supported metal nano-structured catalysts with superior performance has drawn considerable attentions in heterogeneous catalysis.[1] However, “metal-carbon support interaction” is much less understood. For instance, it is not clear that how to tune metal-carbon support interaction for preparing the supported metal nanoparticles with optimal morphologies and structures, stabilizing the chemical environment in interfacial area of metal and supports, and keeping the catalytic performance in severe conditions. Advanced techniques in transmission electron microscopy (TEM) are powerful research tools for probing the metal-carbon support interaction. It can provide the fine surface/interface structures at atomic and sub-electron-volt level, such as defects, location of doped atoms, coordination state, functional group species and electronic structures.[2-3] Here, selected several examples will be demonstrated in this presentation. Figure 1 shows typical high angle annular dark field- scanning TEM (HAADF-STEM) image of Pd rings on oxygen functionalized carbon nanotube (O-CNT) and the schematic representation of the formation of Pd rings on O-CNT. Metal–support interactions between Pd nanoparticles (NPs) and functionalized CNTs were established for controlling the metal-size distribution. Thermal detrapping permits nanosized Pd to possess similar dynamics as generated carbonaceous species under electron irradiation and results in a carbon metal hybrid structure with Pd rings on the edges of the nanobead.[4] Figure 2 shows the surface/interface changes of Pd supported on L-CNT (a) and H-CNT (b) after one hour of Suzuki–Miyaura reactions. Corresponding to the commonly reported high reactivity in homogeneous catalysis, carbon–carbon couplings with high efficiency can be achieved on supported Pd NPs by improving surface functionalization and the dispersibility of the catalyst. Such a system offers opportunities for characterizing surface catalysis with atomic precision, which is crucial for detecting dynamic changes on catalytically active species and understanding catalysis pathways.[5]

Reference
[1] D.S. Su, S. Perathoner, G. Centi, Chem. Rev. 2013, 113, 5782-5816.
[2] B. Zhang, D.S. Su, C.R. Physique 2014, in press (DOI: 10.1016/j.crhy.2013.11.001).
[3] L. Shao, B. Zhang, W. Zhang, D. Teschner, F. Girgsdies, R. Schlögl, D.S. Su, Chem. Eur. J. 2012, 18, 14962-14966.
[4] B. Zhang, L. Shao, W. Zhang, D.S. Su, ChemCatChem 2013, 5, 2581-2585.
[5] L. Shao, B. Zhang, W. Zhang, S.Y. Hong, R. Schlögl, D.S. Su, Angew. Chem. Int. Ed. 2013, 52, 2114-2117.


We gratefully acknowledge the financial support provided by NSFC of China (21133010, 21203215, 51221264, 21261160487), MOST (2011CBA00504), Strategic Priority Research Program of the Chinese Academy of Sciences (No. XDA09030103) and the China Postdoctoral Science Foundation (2012M520652).

Fig. 1: a) HAADF-STEM image of Pd rings/O-CNT. Schematic representation of the formation of Pd rings on O-CNT: b) Morphology of Pd/O-CNT in the initial stage. c) Coupled with heating (red ribbons, bottom), the electron-beam induces a significant rearrangement in the Pd NPs. The Pd rings are formed at the borders of the irradiated areas.[4]

Fig. 2: HRTEM images after 1 hour catalysis of: a) a Pd NP on L-CNTs, b) a Pd NP on H-CNTs. CNTs annealed at 700 oC have more defects than CNTs annealed at 1500 oC, HNO3 treatment introduced a high functionalization (H-CNTs) on defective CNTs and a low functionalization (L-CNTs) on graphitized CNTs.[5]

Type of presentation: Oral

MS-1-O-2191 Visualising the Three-dimensional Morphology and Surface Structure of Metallic Nanoparticles at Atomic Resolution by Automated HAADF STEM Atom Counting

Jones L.1, Fauske V. T.2, MacArthur K. E.1, van Helvoort A. T.2, Nellist P. D.1
1Department of Materials, University of Oxford, Oxford, UK, 2Department of Physics, Norwegian University of Science and Technology, Trondheim, Norway
lewys.jones@materials.ox.ac.uk

Because of their large proportion of surface atoms and favourable chemical activity, metallic nanoparticles are used to catalyse a wide range of technologically important reactions. However, many utilise expensive or rare metals, leading to the desire to account for their content and efficiency at the atomic scale. Aberration-corrected high-angle annular dark-field scanning transmission electron microscopy (HAADF STEM) proves a powerful tool here, with readily interpretable mass-thickness, or Z, contrast images facilitating analysis on an atomic column by column basis. In this work, pure platinum nanoparticles were imaged, such that the image intensity depends on the sample thickness. High-resolution images were recorded using a JEOL-ARM200F whose ADF detector was also calibrated to allow the data to be expressed in units of ‘fractional beam-current’ [1] (E = 200kV, convergence and detector angles of 27 and 69–279mrad).
After magnification calibration, the raw data (Fig 1, left) and the detector efficiency scan were passed to the in-house ‘Absolute Integrator’ software. This software automatically identifies the image peak-positions, normalises the data to units of fractional beam-current, performs a locally adaptive background subtraction (to account for the amorphous carbon-black support), divides the image into Voronoi cells and integrates the signal at each atomic column to yield a map of the absolute scattering cross-sections [2] (Fig 1, right). Comparing these with simulation (multi-slice, 30 phonon runs), the number of atoms per column was identified and a provisional three-dimensional (3D) model was built. Owing to the beam-sensitivity of the particles and the desire for high-throughput analysis, tomography was not possible; instead to obtain the likely z-positions an energy minimisation was performed.
Columns in the starting model were positioned symmetrically about the mid-plane (z = 0) with x-y positions taken from the peak-finding results. The model contained 238 atomic columns with 1656 atoms in total, was 11 atoms high at it thickest, and was assumed to contain no vacancies. This was then energetically relaxed using a modified Lennard-Jones potential; Fig 2 (left) represents the result after around 3½ hours.
From this 3D model the number of nearest-neighbours were calculated and used to colour-code the visualisation. A histogram of these coordination numbers (Fig 2, right) then directly indicates the ratios of the various crystal facets. This ability to observe the relative areas of surface facets opens new possibilities for surface science on an individual particle level and for exploring this in relation to catalytic performance.
[1] Lebeau & Stemmer, Ultramicroscopy 108 (2008) 1653–8
[2] E et al., Ultramicroscopy 133 (2013) 109–19


The research leading to these results has received funding from the European Union Seventh Framework Programme under Grant Agreement 312483 - ESTEEM2 (Integrated Infrastructure Initiative–I3).

Fig. 1: HAADF image of [110] oriented pure platinum nanoparticle (left) and the associated scattering cross-section analysis for each resolvable atomic column (right).

Fig. 2: Relaxed three-dimensional structure of the nanoparticle with colour indicating the nearest-neighbour coordination (left) and accompanying histogram analysis indicating ratios of faceting types.

Type of presentation: Oral

MS-1-O-2203 Plasmon energy from strained GaN quantum wells

Sigle W.1, Benaissa M.2, Korytov M.3, Brault J.4, Vennéguès P.4, van Aken P. A.1
1Max Planck Institute for Intelligent Systems, Stuttgart Center for Electron Microscopy, Stuttgart, Germany, 2CNRST, Rabat, Morocco, 3Leibnitz Institute for Crystal Growth, Berlin, Germany, 4CNRS-CRHEA, Valbonne, France
sigle@is.mpg.de

Monochromated valence electron energy-loss spectroscopy (VEELS) has been used to study the plasmon energy (Ep) from strained GaN quantum wells (QWs) embedded in AlN matrix [1]. The QWs were grown by molecular beam epitaxy. The width of the studied wells was 4, 3, and 2 nm, respectively, separated by a 30 nm thick AlN layer (Fig.1). EFTEM data were recorded in the Zeiss SESAM microscope [2] between 17 and 26 eV using a 0.3 eV energy slit. A Gaussian function was fitted to the volume plasmon peak at each image pixel. After integration parallel to the QW the plasmon energy profile is obtained (Fig.2). Plasmon energies are plotted in Fig.3 (red symbols) versus the inverse square of the well width. It shows a distinct blue-shift of the plasmon peak position with decreasing QW width. In order to take account of the influence of the AlN/GaN interfaces, we solved the relativistic expressions for the begrenzungs effect given by Bolton et al. [3] and Moreau et al. [4]. The interfaces induce an apparent blue shift of the plasmon with decreasing layer width. However, after correction of this shift the dependence of Ep on QW width is still marked (Fig.3, blue symbols). In a second step we considered the influence of the compressive strain of the GaN layers. Such strain is also known to cause a blue shift of the plasmon [5]. Because the critical thickness is about 3 nm the 2-nm- and 3-nm-QWs are completely strained whereas the strain relaxation in the 4-nm-QW is about 0.12 % [6]. Assuming a square-root dependence of the plasmon energy on unit-cell volume, we corrected the measured data (Fig.3, green symbols). The resulting data show reasonably well a linear trend which is consistent with the concept of quantum confinement. Thus, the use of high-resolution valence electron imaging offers the possibility to distinguish the interplay of different confined properties in strained GaN QWs, which is very promising for understanding and exploiting bandgap engineering of nowadays sophisticated devices.

[1] M. Benaissa et al.: Appl. Phys. Lett. 103 (2013) 021901.

[2] C. T. Koch et al.: Microsc. Microanal. 12 (2006) 506.

[3] J. P. R. Bolton and M. Chen: J. Phys.: Cond. Matter 7 (1995) 3389.

[4] P. Moreau et al.: Phys. Rev. B 56 (1997) 6774.

[5] J. Palisaitis et al.: Phys. Rev. B 84 (2011) 245301.

[6] B. Damilano et al.: Appl. Phys. Lett. 75 (1999) 962.


The research leading to these results has received funding from the European Union Seventh

Framework Programme [FP7/2007-2013] under grant agreement n°312483 (ESTEEM2).

Fig. 1: STEM bright-field image showing a set of 3 GaN quantum wells with nominal widths of 4 nm, 3 nm, and 2 nm, respectively, separated by 30 nm-thick AlN barrier layers.

Fig. 2: Volume plasmon energy profile across the three GaN quantum wells.

Fig. 3: As-measured plasmon energies as a function of the QW width (red). Blue symbols are data after correction for interface effects, green symbols after correction of strain effects. The dashed line is a linear fit to the final data.

Type of presentation: Oral

MS-1-O-2269 HR-STEM investigations of metallic nanoparticles grown with superfluidal He-droplets

Knez D.1, Volk A.2, Thaler P.2, Fisslthaler E.1, Grogger W.1, Ernst W. E.2, Hofer F.1
1Institute for Electron Microscopy and Nanoanalysis, Graz University of Technology, Austria, 2Institute of Experimental Physics, Graz University of Technology, Austria
daniel.knez@felmi-zfe.at

Metallic nanoparticles have attracted more and more interest in recent years as they exhibit completely new physical and chemical properties compared to bulk materials. Over the years numerous synthesis methods, mostly based on wet chemical processes, pyrolysis or evaporation have been developed. In contrast, the nanoparticles used for our investigations were synthesized by using superfluid helium nanodroplets (composed of 103 to 106 helium atoms) at around 0.4 K under ultra-high vacuum (UHV) conditions.1 This approach provides exceptional advantages over conventional methods like sequential addition of a wide range of materials.
Thus, nanoparticles can be synthesized with any composition and different structures, with extremely high purity, which cannot be achieved by other known methods. Figure 1 shows a schematic of the synthesis facility.2 Knowledge of the morphology, dimension and composition of the produced particles are not only essential for understanding their physical and chemical properties, but also for optimizing the synthesis parameters. By using a probe corrected, monochromated FEI Titan3 60-300 equipped with a Super-X detector (EDX) and a Gatan Quantum energy filter we performed analytical high resolution STEM (HR-STEM) in order to characterize metal nanoparticles with respect to their morphology and chemistry.  The HR-STEM investigation of Ag nanoparticles on 3 nm carbon (prepared by the He-droplet method) reveals very small Ag clusters (3-6 nm in size) exhibiting a decahedral structure.

Furthermore, bimetallic clusters with a gold-silver core-shell structure were synthesized. STEM images (a and b in Fig. 3) of a AuAg particle reveal that it grew from single spherical particles inside the He droplet. Elemental analysis of the nanoparticles by EELS and EDX clearly showed that Ag and Au, which were added to the droplet sequentially, are not alloyed. The elemental distribution of this particle is shown in images f and g of Fig. 3.
Finally, the nano-optical properties of these metallic clusters will be studied via low-loss EELS measurements in the plasmon regime, which depend on their size, structure and morphology.3, 4 In order to quantify the influence of the underlying substrate, we will also compare conventional carbon films with mechanical exfoliated monolayer substrates (graphene and hexagonal boron nitride).

References:

1. P. Thaler et al., J. Chem. Phys. 140, 44326 (2014).

2. A. Volk et al., J. Chem. Phys. 138, 214312 (2013).

3. F.-P. Schmidt et al., Nano Lett. 12, 5780 (2012).

4. B. Schaffer et al., Micron 40, 269 (2009)


Our research is supported by the European Union within the 7th Framework Programme (FP7/2007-2013) under Grant Agreement no. 312483 (ESTEEM2) as well as by the Austrian Research Promotion Agency (FFG).

Fig. 1: The helium droplets (blue) are produced in the source (1) by evaporation. After passing the skimmer (2) they collide with atoms or molecules (red) evaporated by the thermal evaporator (3). The particles congregate in the center of the droplet and finally land on the target (4) (e.g. a TEM-grid)

Fig. 2: HR-STEM images (a: HAADF, b: BF) of a silver nanoparticle on a 3 nm amorphous carbon film with decahedral morphology

Fig. 3: a: HAADF image of the AuAg core-shell particle; b: STEM BF image showing that the particle grew from single spherical particles; c-d: EDX elemental maps for Au (c) and Ag (d); e: EELS Ag map (calculated via MLLS fitting); f-g: RGB maps illustrating the elemental core-shell distribution with data from (c) and (e) in (f) and from (c) and (d) in (g)

Type of presentation: Oral

MS-1-O-2278 Fabrication of ordered nanopatterns in AlOx thin films by a single UV laser pulse

Szívós J.1,2, Serényi M.1, Gergely-Fülöp E.1, Sáfrán G.1, Lohner T.1
1HAS, RCNS, Institute for Technical Physics and Materials Science, 2University of Pannonia, Doctoral School of Molecular and Nanotechnologies
szivos.janos@ttk.mta.hu

Nano-scale modification of materials received wide research interest, recently. Most of the techniques for the fabrication of ordered nanostructures suffer from low throughput and high costs. Here we report a fast and cheap method to prepare ordered nanopatterns directly, or to prepare masks and imprint molds for nanolithography.

This applies a template of a monolayer of hexagonally self-assembled silica nanospheres (Langmuir-Blodgett (LB)). The sample surface is treated with a single UV laser (l=248 nm) pulse through the LB film. According to our simulations the nanospheres of the LB film focus the laser light as individual lenses [Fig. 1 (a)]. This provides an array of highly intense spots for the fabrication.

RF and DC magnetron sputtered amorphous AlOx layers, as potential masks, were subjected to laser patterning. Structure, morphology and optical properties of the films were characterized by Atomic Force- (AFM), Transmission Electron Microscopy (TEM) and Ellipsometry.

The intensity distribution of the laser spot was mapped by means of a GaP UV photodiode. The distribution is Gaussian-like, as it is shown in Fig. 1 (b). For a reasonably uniform exposure the inner part of the spot, marked with rectangle, was chosen for the treatment.

According to the Selected Area Electron Diffraction (SAED) and TEM results the RF and DC magnetron sputtered layers are fully amorphous [Fig. 2 (a)] and contain Al nanocrystals (nc-Al) embedded in an amorphous AlOx matrix [Fig. 2 (b)], respectively. Ellipsometry revealed that the absorption coefficient (α) of the nc-Al/AlOx layers is about 3 times higher than that of the fully amorphous layers.

The formation of the observed patterns was revealed by AFM and cross sectional (X) TEM. A wide, shallow pit of ~210 nm diameter obtained in the fully amorphous AlOx can be seen in the XTEM image [Fig. 3 (a)]. It is suggested to form by a volume decrease caused by the implosion of nanovoids due to the intense electromagnetic field and shock wave of the UV laser pulse. An AFM image of the pattern of these pits is shown in Fig 3 (b).

Fig. 4 (a) – (c) illustrates the formation of the patterns, at different laser energies, within the nc-Al/AlOx film: (a) shows a formed hillock that contains separate, small bubbles indicating moderate energy impact. An increase of the local energy is suggested to ignite plasma and gas release blowing up large bubbles (b). Further energy increase causes the burst of the bubble forming a crater in the layer (c). This refers to a series of holes observed by AFM in Fig. 4 (d) that is typical for the patterned nc-Al/AlOx film.

Our results show that by applying silica nanosphere LB films and carefully controlled UV laser pulses masks can be fabricated suitable for nanopatterning various thin films.


Z. Szabó’s help with the simulations and B. Fodor’s contribution with ellipsometry are acknowledged. This work was partially supported by National Development Agency grant TÁMOP-4.2.2/B-10/1-2010-0025.

Fig. 1: (a) The simulated lateral intensity distribution right beneath the LB film. White circles mark the nanospheres of the film. (b) The measured intensity distribution map of the UV laser spot. The black rectangle shows the quasi-homogenous area that had been chosen for patterning.

Fig. 2: Plan view dark field TEM images of the RF sputtered layers (a) and the DC sputtered layers (b). The insets are the Selected Area Electron Diffraction (SAED) of the samples. Inset in (a) represents an amorphous diffraction pattern, while additional rings of nanocrystalline Al can be realized in the SAED in (b).

Fig. 3: (a) The cross-sectional TEM image of a pit fabricated in the fully amorphous AlOx layer. (b) An AFM image (about 10 µm x 10 µm) of the ordered pattern of the obtained pits.

Fig. 4: (a) – (c): Pattern formation in the nc-Al/AlOx layer. Moderate local intensity forms a hillock (a). Bubble on a sample treated with higher fluence (b). Sample treated with even higher intensity: a burst bubble (crater) is created (c). Insets: EELS elemental maps of O (red), Si (green) and Al (blue). (d): Typical AFM image of the pattern obtained.

Type of presentation: Oral

MS-1-O-2359 Temperature-induced core-shell reconfiguration of FexO/CoFe2O4 nanocrystals in ordered 2D nanocrystal arrays*†

Yalcin A. O.1, Tichelaar F. D.1, van Huis M. A.2, Zandbergen H. W.1
1Kavli Institute of Nanoscience, Delft University of Technology, Lorentzweg 1, 2628 CJ Delft, The Netherlands, 2Soft Condensed Matter, Debye Institute for Nanomaterials Science, Utrecht University, Princetonplein 5, 3584 CC Utrecht, The Netherlands
a.o.yalcin@tudelft.nl

A large variety of single- and multi-component nanocrystals (NCs) can now be synthesized and integrated into nanocrystal superlattices.1,2 These superstructures and their components have a limited thermal and temporal stability, though, which often hampers their application as functional devices. On the other hand, temperature-induced reconstructions can also reveal opportunities to manipulate properties and access new types of nanostructures.3-5 In-situ atomically resolved monitoring of nanomaterials provides insight into the temperature induced evolution of the individual NC constituents within these superstructures at the atomic level.6 Here, we investigate the effect of temperature annealing on 2D square and hexagonal arrays of FexO/CoFe2O4 core/shell NCs (Figure 1) as a model for many complex oxides by in-situ heating in a transmission electron microscope (TEM). The FexO core has a rock salt structure with some cation deficiencies (x = 0.83-0.95)7 and the lattice constant varies between 0.4255 nm and 0.4294 nm, depending on the oxidation state.7 The CoFe2O4 shell has a spinel crystal structure (lattice constant 0.846 nm).8 Both structures have a face centered cubic (FCC) oxygen sublattice with a lattice mismatch of only 3 %.7 Both cubic and spherical NCs undergo a core-shell reconfiguration at a temperature of approximately 300 ⁰C, whereby the FexO core material segregates at the exterior of the CoFe2O4 shell, forming ‘snowman’-like particles (asymmetric dumbbells) with a small FexO domain attached to a larger CoFe2O4 domain (Figure 2). During reconfiguration, the core volume is filled by the CoFe2O4 shell material. Upon continued annealing, the segregated FexO domains form bridges between the CoFe2O4 domains, followed by coalescence of all domains resulting in loss of ordering in the 2D arrays. Annealed FexO domains contain Co traces as well (Figure 3).

[1] Redl, F.X. et al. Nature 2003, 423, 968–971.
[2] Talapin, D.V. et al. Nature 2009, 461, 964–967.
[3] van Huis, M.A. et al. Nano Letters 2011, 11, 4555–4561.
[4] Figuerola, A. et al. Nano Letters 2010, 10, 3028–3036.
[5] De Trizio, L. et al. ACS Nano 2013, 7, 3997–4005.
[6] van Huis, M.A. et al. Advanced Materials 2009, 21, 4992–4995.
[7] Pichon, B.P. et al. Chemistry of Materials 2011, 23, 2886–2900.
[8] Song, Q. & Zhang, Z.J. Journal of the American Chemical Society 2004, 126, 6164–6168.

* Yalcin, A.O. et al. Nanotechnology 2014, 25, 055601.

† This work has appeared on the journal cover (Nanotechnology Volume 25, Issue 5) as the featured article (http://ej.iop.org/pdf/nano/vol25/na2505-webcover.pdf).


This work is part of the research programme of the Foundation for Fundamental Research on Matter (FOM), which is part of the Netherlands Organization for Scientific Research (NWO).

Fig. 1: TEM images of different types of 2D arrays. a) cubic FexO/CoFe2O4 core/shell NCs forming a square 2D array. Inset figure was taken with an objective aperture inserted (diffraction contrast). Core-shell contrast is observed better in this way. b) Spherical FexO/CoFe2O4 core/shell NCs array forming a hexagonal 2D array.

Fig. 2: TEM images of FexO/CoFe2O4 ‘initially’ core/shell NC arrays; a) cubic NCs at 335 ⁰C, and b) spherical NCs at 360 ⁰C. The insets in the images were taken when the objective aperture was inserted (diffraction contrast). The (200) spacing of FexO is 0.21 nm, and (220) and (311) spacings of CoFe2O4 are 0.3 nm and 0.255 nm respectively.

Fig. 3: Energy filtered TEM Co-mapping of initially spherical NCs. Figure 3a is the zero-loss image and Figure 3b is the corresponding Co map. The dotted lines were used to clarify different domains. Arrows show the FexO domain with Co presence.

Type of presentation: Oral

MS-1-O-2547 Advanced Transmission Electron Microscopy Investigation of Epitaxy-Enabled Morphology Controling ITO Nanowires

Lebedev O. I.1, Shen Y.2, Turner S.3, Van Tendeloo G.3, Wu T.4
1Laboratoire CRISMAT,UMR 6508, CNRS ENSICAEN, F-14050 Caen, France, 2Division of Physics and Applied Physics, Nanyang Technological University, Singapore 637371 , 3EMAT, Department of Physics, University of Antwerp, B-2020, Antwerpen, Belgium, 4Materials Science and Engineering, King Abdullah University of Science and Technology, Thuwal 23955, Saudi Arabia
oleg.lebedev@ensicaen.fr

Controlling nanowire morphology in bottom-up synthesis and allowing the assembly of nanowires on planar substrates is of tremendous importance for device applications in electronics, photonics, sensing and energy conversion. To date, there has however been only limited success in reliably achieving these goals, hindering both the fundamental understanding of the growth mechanism and the integration of nanowires in real-world technologies. In this work, we will show the impact of transmission electron microscopy (TEM) on this domain, as an extremely versatile and powerful technique.

Novel dual-metal Au-Cu alloy nanoparticles were used as a catalyst for tin-doped indium oxide (ITO) nanowire growth. The enhanced mobility of the catalyst nanoparticles (NPs) enables in situ seeded growth of branched ITO nanowires (NWs). The dynamically tuned chemical potentials in the catalyst NPs selectively stabilize a rare cubic indium-tin-oxide phase (ISO), forming epitaxial heterojunctions within individual NW branches. This methodology of selecting phases and forming compositionally abrupt axial heterojunctions in NWs departs from the conventional synthesis routes, giving unprecedented freedom to navigate phase diagrams and promising novel nanomaterials and devices [1]

Here we report that growth of planar, vertical and randomly oriented ITO nanowires can be realized on yttria-stabilized zirconia (YSZ) substrates via the vapor-liquid-solid (VLS) mechanism, by simply regulating the growth conditions, in particular the growth temperature. [2]. TEM and reciprocal space mapping experiments reveal the indispensable role of substrate-nanowire epitaxy in the growth of oriented planar and vertical nanowires at high temperatures, whereas randomly oriented nanowires without epitaxy grow at lower temperature. Further control of the orientation, symmetry and shape of the nanowires was achieved through use of YSZ substrates with (110) and (111), in addition to (001) surfaces. Based on these insights, we succeeded in growing regular arrays of planar ITO nanowires from patterned catalyst nanoparticles. Overall, our discovery of unprecedented orientation control in ITO nanowires advances the general VLS synthesis, providing a robust epitaxy-based approach towards rational synthesis of nanowires.

[1] –J.Gao, O.I.Lebedev, S.Turner, Y.F. Li, Y.H.Lu, Y.P.Feng, P.Boullay, W.Prellier, G.Van Tendeloo, T.Wu Nano Letters 12 (2012) 275-280

[2] - Y. Shen, S. Turner, P. Yang, G. Van Tendeloo, O. I. Lebedev, T. Wu Nature Communications (Submitted) (2014)


This work was also supported in part by the European Union Seventh Framework Programme under Grant Agreement 312483 - ESTEEM2 (Integrated Infrastructure Initiative–I3).

Fig. 1: (a) Low-magnification STEM-HAADF image of the in-plane nanowires and (b) corresponding SAED pattern (c) GPA patterns along [100] and [010] directions. (d) HR STEM-HAADF cross-section image of ITO nanowire and (e) image of the tri-junctions of the ITO, YSZ and Au particle. (f) STEM-ABF image of ITO / YSZ interface with overlayed structural model

Fig. 2: ) Low magnification ADF STEM (top) and BF TEM(bottom) images of off-plane ITO NWs, (b) - high resolution HAADF-STEM image of the ITO-YSZ interface, (c) ABF-STEM image of the top part of an ITO NW. Notice the presence of the ISO phase.

Type of presentation: Oral

MS-1-O-2563 Quantitatively Following growth processes of CdSe@CdS core-shell particles on the atomic scale 

Mangel S.1, Aronovitch E.1, Enyashin A. N.2, Houben L.3, Bar Sadan M.1
1Chemistry Department, Ben Gurion University of the Negev, Beer Sheba, Israel, 2Institute of Solid State Chemistry UB RAS, Ekaterinburg, Russian Federation, 3Peter Grünberg Institut 5 and Ernst Ruska Centre for Microscopy and Spectroscopy with Electrons, Forschungszentrum Jülich GmbH, 52425 Jülich, Germany
barsadan@bgu.ac.il

Colloidal core-shell crystals of II-IV semiconductors are one of the most extensively researched systems in nanoscience. CdSe@CdS nanoparticles are investigated due to their size dependent optical properties and although they have been known for the last two decades, even today new fabrication routes are still explored to improve their optical properties. While reports on the optical properties of single particles are available, the quantitative characterization of atomic order on a single particle level and the growth mechanism that resulted in that specific rearrangement, are still generally missing. The majority of characterization procedures are performed on ensembles that average properties and may hinder the understanding of fundamental aspects in the colloidal synthesis.
Moreover, atomic resolution analysis, which has emerged with aberration corrected instruments, have mainly provided analysis of few particles per sample. It is now, due to the Cc correction that offers superior resolutions in low voltages that the atomic ordering can be achieved on a routine basis to deliver new statistical data. The application of the low voltage reduces the rate of radiation damage so that both surface and bulk structure. We follow the growth process of the CdSe nanoparticles and the formation of the CdS shell covering them. For the first time, statistical atomic-scale information on dozens of individual nanoparticles is correlated with ensemble measurement data.
Previous research showed a proof of concept of determining the polarity and faceting of the nanoparticles by both TEM and HAADF STEM. The assignment of polarity to individual particles gives detailed understanding of when and where stacking faults form, and the new knowledge can be merged with the known kinetics of the reaction. The Cd- terminated edge, where growth is slower, produces more stacking faults or preserves more of the disorder of the original nucleus. The Se edge, which is the fast growing edge, produces almost a perfect W structure. Upon the deposition of the shell, the core is further annealed and stacking faults concentrate at both edges of the particle. However, the annealed sections may acquire larger fractions of the pressure induced ZB symmetry which is quite rare in the pure CdSe nanoparticles.
This analysis shows that high resolution electron microscopy can serve as a routine tool to understand growth kinetics and it may also be applied to the growth of other hybrid nanoparticle structures where kinetic procedures determine the interfaces nature and properties.


Fig. 1: Phase image of exit-plane wavefunctions reconstructed from through-focus series. (a) CdSe cores and (b) CdSe@CdS core shell particle. The wavefunctions were Fourier filtered to eliminate the background pattern of the periodic hexagonal graphene support film.

Fig. 2: Atomistic models of three possible stacking fault sequences, altering between Wurtzite (W) and Zinc Blende (ZB) crystal structures. Cd in violet, S/Se in green.(right) Overlay of the atomistic models on a phase image of a CdSe@CdS core-shell particle exhibiting the three possible atomic stacking fault configurations. (left)

Type of presentation: Oral

MS-1-O-2583 Atomic Structure and Composition Analysis of Pt0.8Ni De-alloyed Nanocatalysts for Proton Exchange Membrane Fuel Cells – Aberration Corrected STEM Study

Rasouli S.1, Sharman J.2, Martinez A.2, Fongalland D.2, Hards G.2, Yamamoto T.3, Myers D.4, Higashida K.3, Ferreira P.1
1University of Texas at Austin, Austin TX , USA, 2Johnson Matthey Technology Centre, Sonning Common, Reading, UK, 3Kyushu University, Fukuoka, JAPAN, 4Argonne National Laboratories, Lemont, IL, USA
ss.rasouli@gmail.com

Proton exchange membrane fuel cells (PEMFCs) are promising energy conversion devices for transport and stationary applications. Pt nanoparticles are currently used as the catalyst in the anode and cathode of the fuel cell, respectively. However, alloys of Pt with base metals are being investigated to replace Pt on the cathode as a way to improve the efficiency of the fuel cell, and reduce cost [1].
In this work Pt-Ni nanoparticles were de-alloyed by acid leaching to produce 5.8 nm size Pt.8Ni catalyst nanoparticles. In order to better understand the relationship between the elemental distribution and the nanoparticle shape and structure, the nanoparticles were characterized by aberration-corrected scanning transmission electron microscopy (STEM), using high-angle annular dark-field (HAADF) imaging [2]. The Pt and Ni compositional distribution of the de-alloyed nanoparticles was investigated using EDS mapping in STEM mode. In order to better understand the three dimensional shape of the nanoparticles and the carbon support, 3-D electron tomography of the nanoparticles was performed in a JEOL JEM ARM 200F. A total of 61 STEM images were collected over a tilt range of -60 to +60 degrees, with a 2° tilting step. The final tilt series was aligned, reconstructed and visualized using Inspect 3D and Amira 4.1, respectively.
Figure 1 shows two aberration-corrected HAADF STEM images of the de-alloyed Pt0.8Ni nanoparticles projected along the [110] beam direction. Although most of the particles exhibit a truncated octahedron shape (Fig. 1a), there are also some particles with long {111} facets (Fig. 1b). The absence of superlattice reflections in the FFTs (insets of Figs 1a and 1b) shows that Pt and Ni are in solid solution, forming a face-centered cubic structure. However, as shown in Figures 1c and 1d, the nanoparticles exhibit regions of bright and dark contrast, which indicate that the composition of Pt and Ni is not uniform throughout the particle. This heterogeneous distribution among the various particles is a result of the de-alloying process. Moreover, although most of the particles are a Pt-Ni solid solution, there are also some particles exhibiting {100} supperlattice reflections in the FFTs, indicating a partially ordered structure (Figs 1e and 1f). In addition, Pt seems to segregate to the surface of the nanoparticles, as confirmed by EDS Mapping (Fig. 2). Finally, 3-D electron tomography confirms that most of the particles exhibit a truncated octahedron shape (Fig. 3).

References
[1] S. C. Ball et al. ECS Transactions, 11(1) (2007), p.1267.
[2] S. Chen, et al. J. Phy. Chem. C. 113(3) (2009), p.1109.


The authors gratefully acknowledge funding from the U.S. Department of Energy, Office of Energy Efficiency and Renewable Energy, Fuel Cell Technologies Office (Nancy Garland, Technology Development Manager).

Fig. 1: (a) and (b) Aberration corrected HAADF images of Pt0.8Ni nanoparticles and corresponding FFTs (insets). (c) and (d) normalized intensities across the nanoparticle (along the red line). (e) and (f) Aberration corrected HAADF images of the nanoparticles with supperlattice reflections in FFTs (insets).

Fig. 2: HAADF STEM image and EDS mapping of Pt0.8Ni nanoparticles

Fig. 3: 3D reconstructed image of the Pt0.8Ni nanoparticles

Fig. 4:
Type of presentation: Oral

MS-1-O-2588 High Spatial/Energy Resolution Cathodoluminescene Spectroscopy: A Powerful Tool for Profound Characterization of the Physical Properties of the Advanced Nanostructures

FU X.1, FU Q.1, HAN X.1, LIU C.1, ZHANG J.1, LIAO Z.1, YU D.1
1Department of Physics, State Key Laboratory for Mesoscopic Physics, Peking University, and Collaborative Innovation Center of Quantum Matter, Beijing 100871, P. R. China
yudp@pku.edu.cn

    High special/energy resolution cathodoluminescence (CL) spectroscopy is now becoming more and more important in investigations of the optical properties of the low-dimensional structures. Two representative examples of the application of the CL in study were summarized in this presentation as follows:

    Elastic engineering strain has been regarded as a low-cost and continuous variable manner for altering the physical and chemical properties of materials, and it becomes even more important at low-dimensionality because at micro/nanoscale, materials/structures can usually bear exceptionally high elastic strains before failure. The elastic strain effects are therefore greatly “magnified” in micro/nanoscale structures and should be of great potential in the design of functional devices. The purpose of this presentation is to present a summary of our recently progresses in the energy band engineering of elastically strained ZnO micro/nanowires. First, we present the electronic and mechanical coupling effect in bent ZnO nanowires. Second, we summary the bending strain gradient effect on the near-band-edge (NBE) emission photon energy of bent ZnO micro/nanowires. Third, we show that the strain can induce exciton fine-structure splitting and shift in ZnO microwire. Related publications are presented in Figure 1.

    Surface plasmon polaritons (SPPs) show great potential for application in future nanoscale photonic systems due to the strong field confinement at the nanoscale, intensive local field enhancement, and interplay between strongly localized and propagating SPPs. A template stripping method combined with PMMA as a template was successfully developed to create extraordinarily smooth metal nanostructures with a desirable feature size and morphology for plasmonics and metamaterials. The advantages of this method, including the high resolution, precipitous top-to bottom profile with a high aspect ratio, and three-dimensional characteristics, make it very suitable for the fabrication of plasmonic structures. The confined modes of surface plasmon polaritons in these nanocavities have been investigated and imaged by using cathodoluminescence spectroscopy, which has been turned out to be a powerful means to characterize the resonant SPPs modes confined in metal nanocavities. The mode of the out-of-plane field components of surface plasmon polaritons dominates the experimental mode patterns, indicating that the electron beam locally excites the out-of-plane field component of surface plasmon polaritons.


This work was supported by MOST (Nos. 2013CB934600, 2013CB932602), NSFC (Nos. 11274014, 11234001), and the Program for New Century Excellent Talents in University of China (No. NCET-12-0002).

Fig. 1: Recent publications of CL spectroscopy summarized in the first part of the presentations demonstrating the advantages compared to the conventional optical methods.

Fig. 2: Publications related to the second part of the presentations via CL for characterization of the SPP modes confined in metal nanocavities, which is impossible to do it via other optical approaches.

Type of presentation: Oral

MS-1-O-2629 Insight into the structural, electrical and photoresponse properties of individual Fe:SrTiO3 nanotubes

Žagar K.1, Fabrega C.2, Hernandez-Ramirez F.2,3, Prades J. D.3, Morante J. R.2,3, Rečnik A.1, Čeh M.1
1Jožef Stefan Institute, Ljubljana, Slovenia, 2Catalonia Institute for Energy Research, Barceloba Spain, 3University of Barcelona, Barcelona, Spain
kristina.zagar@ijs.si

Titanates are suitable for many applications such as oxygen sensing and tunable high temperature superconducting microwave filters. The potential advantages of the nanostructured forms have been scarcely explored compared to other oxides. We report on the structural and electrical properties of individual iron-doped strontium titanate nanotubes (Fe:SrTiO3). The Fe:SrTiO3 nanotubes were assessed for the first time, showing high stability and reproducibility [1].

Fe:SrTiO3 nanotubes were synthesized using sol-gel electrophoretic deposition (EPD) technique [2]. The Fe:SrTiO3 sol, where 2 mol% of Ti was replaced by Fe, was deposited into the anodic alumina template while a potential was applied between the AAO/Al working electrode and Pt counter electrode. After the deposition samples were annealed at 700 °C for 1 h with subsequent template removal. Resulting Fe:SrTiO3 nanotubes were characterized by electron microscopy techniques. To study electrical properties, Fe:SrTiO3 nanotube devices were fabricated by focused ion beam nanolithography techniques [3].

Obtained Fe:SrTiO3 nanotubes with lengths between 5 and 10 µm and diameters of approximately 200 nm were polycrystalline, dense and made up of cubic grains ranging between 10 and 20 nm in size (Figure 1). Their chemical composition explored by Energy-dispersive X-ray (EDX) analysis showed the presence of Sr, Ti and Fe; and confirmed that Fe was effectively incorporated into the perovskite structure.

For the electrical characterization the prototype device was formed by integration of individual Fe:SrTiO3 nanotubes into simple circuit architecture and the electrical resistivity of approx. 35 ohm∙cm was calculated (Figure 2). This value was significantly lower than the values for intrinsic bulk SrTiO3 samples due to the presence of Fe. This result opens the door to the future synthesis of Fe:SrTiO3 nanotubes suitable for monitoring small trace level of oxygen. Furthermore, some devices were tested as UV-detectors with the final aim to explore the optoelectronics characteristics and validate their suitability for device integration. The dynamic behavior of the photoresponse obtained with a single Fe:SrTiO3 nanotube as a function of different UV photon fluxes is shown in Figure 3. Repeatable and reversible responses were found in all cases, demonstrating that our devices are nice proof-of concept systems showing that an Fe:SrTiO3 nanotube can be used as a UV photodetector.

References:

1. K. Zagar et al., J. Mat. Chem. Phys. 141 (2013), p. 9.

2. K. Zagar et al., Nanotechnology 21 (2010) p. 375605.

3. F. Hernandez-Ramirez et al., Chem. Phys. 11 (2009), p. 7105.


The research was also supported by the Framework 7 program under the project S3 (FP7-NMP-2009-247768) and European Union Seventh Framework Programme [FP7/2007-2013] under grant agreement n°312483 (ESTEEM2).

Fig. 1: (a) Bright-field image of a uniformly shaped and polycrystalline Fe:SrTiO3 nanotube. (b) Higher-magnification bright-field TEM image of polycrystalline Fe:SrTiO3 nanotube with grain sizes in the range from 10 to 20 nm.

Fig. 2: (a) Fe:SrTiO3 nanotube electrically contacted in 2-probe configuration using FIB lithography. (b) I-V curve of the contacted Fe:SrTiO3 nanotube. An ohmic response is found at room temperature and open air atmosphere.

Fig. 3: (a) Photoresponse Iph of a Fe:SrTiO3 nanotube as function of different UV photon intensities (b) Dynamic response of Iph as function of different UV photon fluxes.

Type of presentation: Oral

MS-1-O-2638 Core-shell GaAs/AlGaAs nanowires grown on Si (111)

Kehagias T.1, Florini N.1, Walther T.2, Moratis K.3, Hatzopoulos Z.3, Pelekanos N. T.3
1Physics Department, Aristotle University of Thessaloniki, GR-54124, Thessaloniki, Greece, 2Department of Electronic and Electrical Engineering, University of Sheffield, Mappin St, Sheffield S1 3JD, UK, 3Materials Science & Technology and Physics Departments, University of Crete and IESL/FORTH, GR-71003 Heraklion, Greece
kehagias@auth.gr

Precise control over III-V compound semiconductor nanowires (NWs) growth is crucial for the fabrication of advanced nanoscale electronic and optoelectronic devices. Core-shell GaAs/AlGaAs NWs were grown on Si (111) by plasma assisted molecular beam epitaxy (PAMBE). Initially, the NWs were grown via the vapor-liquid-solid mechanism, using Ga droplets as catalyst, for 20 min. Subsequently, the Ga droplets were removed by exposing the NWs to As flux and growth continued for another 40 min, varying the fluxes of the Al, Ga, and As, in order to form an AlGaAs shell around the GaAs initial core.

The structural features of the NWs were characterized by transmission electron microscopy (TEM) methods. TEM observations and selected area diffraction analysis showed that NWs are zinc-blende (ZB) single crystals grown epitaxially along the [111] direction normal to the Si substrate, despite the presence of a 1-3 nm thick amorphous SiO2 layer on the Si surface. Simultaneously, an interfacial GaAs layer is formed between the NWs, comprising large epitaxial and {111} twin related crystals [Fig. 1(a)]. The emanation point of the NWs is located on small heavily twinned GaAs crystals [Fig 1(b)], which evolve into NWs and finally, through the growth process usually merge with the GaAs crystals of the interfacial layer. In addition, “parasitic” NWs emerging from either the interface, or the original NWs were observed along the inclined <111> and/or the <100> directions. Mirror twins normal to the [111] growth direction can be observed throughout the length of the NWs [Fig. 1(c)]. In fact, NWs grow for several micrometers under a continuous succession of mirror twins. No wurtzite structure was observed.

The weak absorption contrast of high-resolution TEM (HRTEM) in conjunction with the minimal difference of the AlGaAs and GaAs lattice parameters turn HRTEM images unsuitable for visualizing the core-shell structure. Hence, the chemically sensitive 200 reflection for mass contrast TEM imaging [Fig. 1(d)], in addition to annular dark-field (ADF) scanning TEM (STEM) imaging [Fig. 1(e)], were used. These revealed the core-shell configuration of the NWs, where the AlGaAs shell spans from one half to 2/3 of the projected diameter of the NWs ranging from 80 nm to 180 nm. Furthermore, energy dispersive X-ray (EDX) analysis confirmed the core-shell morphology of the NWs and was used to estimate the NW shell composition.


Research co-financed by the European Union (European Social Fund-ESF) and Greek national funds through the Operational Program "Education and Lifelong Learning" of the National Strategic Research Frame (NSRF)-Research Funding Program: ARISTEIA II, project “NILES”.

Fig. 1: (a) TEM image, taken off the [1-10] zone axis, showing the NWs morphology. (b) A NW emerging from a small defected GaAs crystal. (c) HRTEM image, along the [1-10] direction, where mirror twins are depicted by arrows. The amorphous shell is attributed to oxidation. (d)&(e) TEM and ADF STEM images revealing the GaAs/AlGaAs core-shell structure.

Type of presentation: Oral

MS-1-O-2680 Understanding Symmetry Breaking in Anisotropic Nanoparticle Growth

Walsh M. J.1, Barrow S. J.2, Tong W.2, Funston A. M.2, Etheridge J.3
1Department of Materials Engineering, Monash University, Australia, 2School of Chemistry, Monash University, Australia, 3Monash Centre for Electron Microscopy, Monash University, Australia
michael.j.walsh@monash.edu

The highly promising optical, catalytic and electronic properties of gold nanoparticles, particularly nanorods, have made them a major area of research in recent years. In plasmonics, the light absorption and scattering properties of biocompatible Au nanorods make them effective biosensors, whilst their tuneable aspect ratio allows the localised surface plasmon resonant (LSPR) frequency to be shifted into the biologically transparent near infra-red range, opening up potential applications within drug delivery and photothermal hyperthermia treatments [1].

Gold nanorods are typically synthesised via a seed-mediated approach, in which silver ions and halides are used as surfactants [2]. There is currently little agreement on a mechanism for anisotropic growth, and few insights into the fundamental symmetry breaking event that is a prerequisite for shape anisotropy. The question remains as to what causes an essentially spherical seed particle, with a cubic lattice, to develop a preferential growth direction? Here we present an aberration corrected electron microscopy study of nanoparticles at the embryonic stages of growth to provide direct atomic scale insights into the onset of anisotropy.

Seed particles were found to be predominantly single crystal. Overgrowth of these seeds with and without the presence of silver ions is shown in figure 1; clearly revealing that it is the crucial addition of Ag+ that induces symmetry breaking. This event occurs at particle diameters of between 4-6 nm, and only for single crystal structures. We suggest a mechanism for symmetry breaking in which Ag stabilises small {110} truncations in the seed particle structure. The various stages of anisotropic growth are described in figure 2, in which particle size and morphology are characterised at regular intervals during nanorod synthesis. After the initial symmetry breaking event, rapid growth along the rod axis results in a variety of particle sizes. Nano-dumbbells are formed after several minutes, corresponding to a maximum redshift in the longitudinal LSPR frequency. Reduced growth in the length direction cause the dumbbell to transition to a nanorod morphology, with a subsequent blueshift in LSPR.

These observations are consistent with a silver under-potential deposition mechanism, and we discuss this and possible synergistic effects of AgBr complexing as driving forces for anisotropic growth. Precise understanding of the growth mechanism and resulting structure-property relationships of nanoparticles should allow for high yields of particles with size, shape and crystal surfaces tailored to a variety of potential applications.

References
[1] S. Abalde-Cela et al. J. R. Soc. Interface. 2010, 7, 435
[2] S. E. Lohse and C. J. Murphy. Chem. Mater. 2013. 25, 1250


This work was supported by the Australian Research Council (ARC) grant DP120101573 and used microscopes at the Monash Centre for Electron Microscopy funded by ARC Grant LE0454166

Fig. 1: Standard seed particles (left), overgrown seeds (centre) and seed particles overgrown in the presence of Ag ions (right).

Fig. 2: The stages of seed mediated Au nanorod growth after 0, 2, 20 and 60 minutes, with corresponding UV-Vis spectra.

Type of presentation: Oral

MS-1-O-2712 Spectral unmixing of electron energy-loss spectra of ~5 nm InP/ZnS nanocrystals

Duchamp M.1, Xi L.2, Lam Y. M.3, Dunin-Borkowski R. E.1, Kardynal B.2
1Ernst Ruska-Centre for Microscopy and Spectroscopy with Electrons, Forschungszentrum Jülich, Germany, 2Peter Grünberg Institute 9, Forschungszentrum Jülich, Germany, 3Institute of Materials for Electronic Engineering II, RWTH-Aachen, Sommerfeldstr. 24, D-52074 Aachen, Germany
martial.duchamp@gmail.com

We examine InP/ZnS nanocrystals (NCs), with possible core-shell or alloyed structures [1][2]. As InP and ZnS are both cubic and have very similar lattice constants, their compounds are difficult to distinguish using high-resolution transmission electron microscopy (TEM). This information is important, since the presence of a ZnS shell is essential in reducing surface recombination but also prevent oxidation of the InP [3].
Scanning TEM and electron energy-electron spectroscopy (EELS) were performed at 80 kV, in order to quantify the distribution of In- and Zn-containing compounds in individual NCs and their oxidation states after oxidation of the NCs in an oxygen/argon plasma to simulate aging process. Data analysis involved spectral unmixing using vertex component analysis (VCA) [4][5], in order to improve the signal-to-noise ratios of In and Zn elemental maps and to extract spectral signatures at the O K edge.
Figures 1a and b show conventional background-subtracted elemental maps measured for the In M5,4 and Zn L3,2 edges. VCA was applied over the energy range 850-1100 eV. A spectral component corresponding to the Zn L3,2 edge can be identified. The corresponding map (Fig. 1e) is similar to Fig. 1b but less noisy. Surprisingly, an abondance map corresponding to In is also obtained over this energy range, where no In edge is present. This signal originates from the background of the In M3 edge (678 eV) and is similar to Fig. 1a, which was extracted at the In M5,4 edge.
Figure 2 shows results obtained using VCA over the energy range 450-600 eV. Of the three extracted components, component 1 corresponds to an In signal. The abundance map associated with this component is similar to the In maps as a result of the presence of the In M5,4 edge in this energy range. More interestingly, the O K edge shows a shoulder at ~533eV. This shoulder is in a similar position but is less intense than the first peak in an In2O3 reference spectrum, suggesting that the NCs are partially oxidised. Although a clear signature of the oxidation of ZnS is not observed, oxidation of the In could be explained by the presence of either an incomplete ZnS shell or a ZnS layer that oxidized during the plasma treatment.
Our results show that a thin ZnS shell does not protect InP from oxidation sufficiently well for long-term applications. Our use of the VCA algorithm allows the oxidation of In in sub-5-nm InP/ZnS NCs to be identified with much greater confidence than using conventional background-subtraction methods.

[1] H. Borchert et al. Nano Lett. 2 (2002) 151; [2] K. Huang et al. ACSNano 4 (2010) 4799; [3] J. Jasinski et al. Solid State Commun. 141 (2007) 624; [4] M. Duchamp et al. Appl. Phys. Lett. 102 (2013) 133902; [5] C. Boothroyd et al. Ultramicroscopy, in press (2014)


The authors acknowledge financial support from the European Union under the Seventh Framework Programme (project references 312483 - ESTEEM2 and NWs4LIGHT).

Fig. 1: (a, b) Background-subtracted elemental maps corresponding to the In M5,4 and Zn L3,2 edges; (d, e) Corresponding maps obtained by applying VCA over the energy range 850-1100 eV. (c) Spectra extracted from the spectrum image at the positions marked in (a); (f) Spectral components associated with the maps shown in (d, e)

Fig. 2: Spectral components extracted over the energy range 450-600 eV, which includes the In M5,4 and O K edges. The corresponding abundance maps are shown on the right. A reference In2O3 spectrum is also shown. The scale bar is 15 nm.

Type of presentation: Oral

MS-1-O-2715 Investigation of Detailed Monolayer MoS2 Edge Structure and Defect Configuration by Atomic Resolution Scanning Transmission Electron Microscopy

Li K.1, Hong J.2, Jin C.2, Zhang X.3
1Imaging and Characterization Core Lab, King Abdullah University of Science and Technology, Thuwal 23955-6900, Kingdom of Saudi Arabia, 2Department of Materials Science and Engineering, Zhejiang University, China, 3Division of Physical Science, King Abdullah University of Science and Technology, Thuwal 23955-6900, Kingdom of Saudi Arabia
kun.li@kaust.edu.sa

Recent intense research and development activities in graphene have regained strong research interest in its related two dimensional (2D) counterparts, especially BN and transition metal dichalcogenides (TMDCs), with MoS2 under the focus. The idea has been fortified again that in addition to composition and structure dimensionality also plays a very important role in determining the fundamental properties due to quantum confinement effect. These materials possess exotic properties that are absent in their bulk forms. For MoS2, understanding its edge and defect configuration at atomic level is critical for realizing its application potentials as electronic devices and catalysis, as it affects the electronic band structure of MoS2 and the catalytic behavior.

Here we employ a probe Cs corrected Titan scanning transmission electron microscope (STEM) operated at 80 kV to study the defect and edge structure of monolayer MoS2. Z-contrast STEM technique is used to differentiate between Mo and S atom columns, with Mo showing brighter contrast and S showing darker contrast. Vacancies with different size are found in our study, ranging from mono vacancy to large triangular-shaped vacancies. Fig. 1 shows a triangular-shaped vacancy with a lateral length of 3 unit cells, revealing Mo-terminated Klein edge. Mo-terminated zigzag edge is found in a bigger triangular-shaped vacancy with a lateral length of 4 unit cells, as shown in Fig. 2. The results suggest that S atoms are easier than Mo atoms to be removed under electron beam illumination and Mo-terminated edge is the most often found type. We also find in our experiment that when a nano-ribbons is formed, one side of it is Klein edge, and the other side is zigzag, where removed Mo atoms tend to be absorbed. Based on this observation we believe that for catalysis application, zigzag edge should also be the preferred absorption site for noble catalytic atoms. Dislocation pairs are also found in this study with atomic resolution (Fig. 3) and the Burgers vector is also defined (Fig. 4); it is a 60-degree dislocation.

With no doubt the capability of unambiguously identifying defect and edge configuration at atomic level will facilitate process optimization for realizing the promising applications of MoS2 and its related TMDCs.


Fig. 1: Atomic resolution defect configuration of triangular-shaped vacancies with the lateral length of three unit cell and Mo-terminated Klein edge.

Fig. 2: Atomic resolution defect configuration of triangular-shaped vacancies with the lateral length of four unit cell and Mo-terminated zigzag edge

Fig. 3: Dislocation pairs in MoS2 Monolayer.

Fig. 4: Dislocation in Figure 3 with Burgers vector identified

Type of presentation: Oral

MS-1-O-2842 Gold repartition at surfaces and interface in silicon nanowires: A TEM / APT confrontation

Grillet N.1, David T.1, Roussel L.1, Neisius T.2, Cabie M.2, Gailhanou M.1, Alfonso C.1, Charaï A.1, El Kousseifi M.1, Hoummada K.1, Descoins M.1, Mangelinck D.1
1Aix-Marseille Université – CNRS, IM2NP, Faculté des Sciences de Jérôme, F-13397 Marseille, France, 2Aix-Marseille Université, CP2M, Faculté des Sciences de Jérôme, F-13397 Marseille, France
nadia.grillet@im2np.fr

The development of new methodologies is needed for precise measurement of concentration and localization of different species insmall objects like nanowires(NWs). Transmission Electron Microscopy (TEM) and Atom Probe Tomography (APT) are two major techniques which give complementary informations at atomic scale.
This work deals with post-growth repartition of the gold used as catalyst during Molecular Beam Epitaxy (MBE) growth of silicon NWs. Electron microscopy analysis was performed using Scanning Transmission Electron Microscopy - High Angle Annular Dark Field (STEM-HAADF), electron tomography, and Energy-dispersive X-ray spectroscopy (EDS) to collect complementary information on morphology, structure and chemistry. Atom probe tomography has ability for both 3D imaging and chemical composition measurements at the atomic scale. The two techniques were used to characterize the same samples containing silicon nanowires deposited by MBE.

We observed in STEM-HAADF that the NWs have a hexagonal (only {112} faces) and/or dodecagonal section ({112} and {110} faces) and present a 'saw-tooth' faceting on one over two {112} faces (cf. Figure 1 a) and b)).
We found that gold clusters are spread on the surfaces of the NWs, but no gold was observed in the NW bulk. Furthermore, the gold coverage is uneven on the different faces of the NW. Its repartition is very homogeneous on the ‘flat’ {112} faces. But, when the {112} saw-tooth faceting is present, the gold spreads preferentially on the {113} facets, rather than on the {111} facets (cf. Figure 1 c)) [1].
A rough estimation considering that gold clusters are hemispherical allowedus to estimate that {113} facets are covered by 3 to 6ML of gold which is largely higher than coverage found in literature [2].

The three dimensional distribution of Au has been determined in the volume and on the surface of the Si NW also by APT. These results show no gold detection in the bulk of the NW (Figure 2), which is in good agreement with previous studies [3]. Moreover, Au clusters detected on the Si-NW surface by APT are in accordance with TEM measurements.

An investigation of the interfacial region between catalyst and silicon nanowire has also been done using EDS mapping and APT. Both techniques indicate that a silicon oxide layer was formed between the gold catalyst and the Si NW. In addition, ATP measurements showed the presence of a SiAu alloy layer containing less than 1% gold underneath the catalyst droplet.

[1] T. David et al., Journal of Crystal Growth, 383, 151-157 (2013)
[2] C. Wiethoff et al., Nano Letters 8, 3065-3068(2008)
[3] J.E. Allen et al., Nature Nanotechnology 3, 168-73 (2008)


Fig. 1: Tomography STEM-HAADF showing the different facets and the repartition of gold on them

Fig. 2: STEM-HAADF / EDS mapping versus APT at the Au/Si interface

Type of presentation: Oral

MS-1-O-3052 Structure and Phase Analyses of Nanoparticles using Combined Analysis of TEM scattering patterns

Boullay P.1, Lutterotti L.2, Chateigner D.1
1CRISMAT, CNRS UMR 6508, CAEN, France, 2Department of Industrial Engineering, University of Trento, TRENTO, Italy
philippe.boullay@ensicaen.fr

The development of materials science at the nanoscale questions the characterization techniques on their ability to describe small objects, either individually or as large assemblies. Transmission electron microscopy (TEM) appears as one of the techniques able to provide quantitative results using imaging, spectroscopic or diffraction methods. Aiming the structure, size and phase analysis of nanoparticles, a TEM approach would ideally combine these methods at the nanometer scale but analyses on individual particles are not ideal if one wants a representative statistical analysis.

Another approach would be based on the quantitative analysis of electron diffraction intensities similarly to what is done in X-ray Powder Diffraction (XPD). Selected Area Electron Diffraction patterns of an assembly of nanoparticles exhibit ring patterns analogous to those from XPD, hereafter called Electron Powder Diffraction patterns (EPD). Phase identification and structure refinement of such powder diffraction patterns can be reached by search-match routines followed by Rietveld analysis [1-2] or PDF (Pair Distribution Function) [3-4] methods. Besides the phase identification and structure refinement issue, we will show that the average size and shape of the crystallites (Fig. 1) as well as quantitative texture analysis (Fig. 2) can be obtained from EPD [5]. Using Rietveld analysis within the Combined Analysis methodology, almost routine analyses of nanoparticles in the form of powders and thin films can be achieved. Complementary measurements can be added, for instance Energy Dispersive X-ray Spectroscopy in order to constrain the refinements in cases for which elemental variations are of matter, and PDF, in order to quantify even amorphous structures [3,6].

This reciprocal space approach allows a fast access to statistically meaningful information about the average size and shape of an assembly of nanoparticles (agglomerated or not). It is thus very complementary to direct imaging of isolated nanoparticles. Fast and insensitive to sample drift, this approach shall be advantageously used for gaining quantitative information from in-situ environmental studies of dynamic processes involving nanoparticles.

[1] T.E. Weirich, M. Winterer, S. Seifried, H. Hahn and H. Fuess, Ultramicroscopy 81 (2000) 263.
[2] A.M. Tonejc, I. Djerdj and A. Tonejc, Mat. Sci. Eng. C 19 (2002) 85.
[3] T. Takagi, T. Ohkubo, Y. Hirotsu, B.S. Murty, K. Hono and D. Shindo, App. Phys. Lett. 79 (2001) 485.
[4] A.M.M. Abeykoon, C.D. Malliakas, P. Juhás, E.S. Božin, M.G. Kanatzidis, S.J.L. Billinge, Z. Kristallogr. 227 (2012) 248.
[5] P. Boullay, L. Lutterotti, D. Chateigner and L. Sicard, Acta Cryst. A (2014) in press.
[6] D.J.H. Cockayne and D.R. McKenzie, Acta Cryst. A 44 (1988) 870.


LL and DC warmly thank the Conseil Régional de Basse-Normandie and FEDER for financing LLs' Chair of Excellence at CRISMAT, and the Université de Caen Basse-Normandie for two months as invited professor of LL. PB and DC thanks the project FURNACE funded by the French research agency (contract ANR-11-BS08-0014).

Fig. 1: a) Mn3O4 nanoparticle’s aggregates. Associated EPD in b) and 1D plot in c) representing the full integration along the Debye rings. The profile is fitted (Rw=2.06% and RBragg=1.55%) considering the sample contribution and compared with XPD in d). e) TEM bright field image of isolated particles. f) Average size and shape of the Mn3O4 nanoparticles.

Fig. 2: a) EPD for 2 extreme and 0° sample tilts obtained on a Pt thin film deposited on a Si single crystal substrate. b) corresponding 1D patterns using Dh = 10° and for h = 180°. c) 2D plots for the 35 1D-patterns of each 2D pattern in a). Experimental data (bottom) and fits (up) are represented, Pawley pattern matching. Square root intensity scales.

Type of presentation: Oral

MS-1-O-3066 Atomic scale studies of individual catalyst nanoparticles with atom probe tomography

Cairney J. M.1, Felfer P. J.1, Eder K.1, Maschmeyer T.2, Masters A.2
1Australian Centre for Microscopy & Microanalysis, The University of Sydney, Sydney, NSW 2006 Australia, 2School of Chemistry, The University of Sydney, Sydney, NSW 2006 Australia
julie.cairney@sydney.edu.au

From sunscreen to optoelectronics, sensors, catalysis and drug delivery, nanometer-scale particles play an important role in a rapidly growing range of applications. An important example of the commercial application of nanoparticles is in the field of catalysis. By maximising the surface area of catalytic metals through the use of nanoparticles, catalytic reactivity can be greatly enhanced and the selectivity strongly influenced. Bi- or multimetallic particles offer even greater scope for fine-tuning.

To better understand their catalytic performance, one must gain a detailed understanding of the size, shape, composition and, most importantly, the arrangement of atoms within and on the surface of the particles. While some atomic scale information on the structure of nanoparticles has long been accessible through electron microscopy [1], identifying the chemical nature and 3D location of the individual atoms remains a challenge using such techniques.

The potential for APT to provide microstructural information for catalysis is well-recognized, and experiments on nanoparticles have proven to be promising, but experimentally challenging [2,3]. Here, we demonstrate how high-resolution atom probe tomography (APT) can be used to quantitatively determine the three-dimensional distribution of atoms within a Au@Ag core-shell nanoparticle with a resolution of +/- 0.5 nm. Specifically, we will describe several major advances in atom probe techniques in recent years that are specifically suited to the study of nanoparticles. These include new specimen preparation techniques that overcome the conventional barrier to the study of nanoparticles by atom probe, and recent advances in the available methods to extract information about the segregation of atoms to three dimensional surfaces that allow mapping of the distribution of the shell species in core-shell particles.

By using these new tools, we reveal that the elements in the Au@Ag nanoparticles are not evenly distributed across the surface and that this distribution is related to the surface morphology and residues from the particle synthesis. Access to this type of information is a revolutionary step forward for the rational design of nanoparticles.

References:

[1] Kiely, C., Electron microscopy: New views of catalysts. Nature Materials, 2010. 9: p. 296-297.

[2] Xiang, Y., et al., Long-chain terminal alcohols through catalytic CO hydrogenation. Journal of the American Chemical Society, 2013. 135(19): p. 7114-7117.

[3] Tedsree, K., et al., Hydrogen production from formic acid decomposition at room temperature using a Ag-Pd core-shell nanocatalyst. Nat. Nano, 2011. 6(5): p. 302-307.


The authors acknowledge the facilities and the scientific and technical assistance of the Australian Microscopy & Microanalysis Research Facility at the Australian Centre of Microscopy & Microanalysis

Type of presentation: Oral

MS-1-O-3178 Imaging Nano Segregation in Advanced Pt Alloy Fuel Cell Electrocatalysts

Gan L.1, 3, Heggen M.2, Cui C.1, Rudi S.1, Strasser P.1
1The electrochemical Catalysis, Energy and Materials Science Laboratory, Department of Chemistry, Technical University Berlin, 10623 Berlin, German, 2Ernst Ruska Center for Microscopy and Spectroscopy with Electrons, Forschungszentrum Juelich GmbH, 52425 Juelich, Germany, 3Division of Energy and Environment, Graduate School at Shenzhen, Tsinghua Unviersity, Shenzhen 518055, PR China
lgan.thu@gmail.com

Segregation is an important physical phenomenon in alloy materials and has significant influences on the physical and chemical properties. In particular, segregation at the surface or subsurface can drastically change the molecular adsorption properties of alloy surfaces and thus becomes a promising way to design highly active catalysts.1 Atomic understanding of segregation effect in nanoscale alloy catalyst particles is therefore crucial yet still challenging for future catalyst designs.

In this talk, we will highlight some of our recent works on understanding and controlling nano segregation effect in advanced Pt alloy catalysts for fuel cell technologies.2-5 We demonstrate how alloy composition (Fig. 1), particle size (Fig. 2), and particle shape (Fig.3) can result in different segregation behaviors in Pt alloy nanoparticles and thus drastically influence their catalytic activity and stability. Focus will be placed on the atomic imaging of nano segregation by using state-of-the-art aberration-corrected scanning transmission electron microscopy (STEM) and electron energy loss spectroscopy (EELS) and in situ experiments. In particular, by using STEM-EELS elemental profile/mapping, we revealed novel compositional segregation patterns in PtxNi1-x core-shell nanoparticles, showing an unexpected Ni-segregated inner shells depending on the bulk alloy composition (Fig. 1).2 Furthermore, we discovered a distinctly different compositional segregation in octahedral PtxNi1-x nanoparticles, which featured a surprising Pt segregation at the edges/corners and Ni segregation at the facets (Fig. 3).4 We explored the physical origin for these distinct segregation behaviors and their impact on the catalytic activities and stability for fuel cell reactions. This will be further complemented by in-situ STEM-EELS experiments to study the structural and compositional evolution of Pt alloy nanoparticles during nanoparticle synthesis, post thermal annealing, and solution-phase electrocatalysis, shedding important light for catalyst designs with desired segregation patterns and chemical properties.

References:

(1)Stamenkovic, V. R. et al. Science 2007, 315, 493-497. (2)Gan, L., Heggen, M., Rudi, S. & Strasser, P. Nano Lett 2012, 12, 5423-5430. (3)Gan, L., Heggen, M., O'Malley, R., Theobald, B. & Strasser, P. Nano Lett 2013, 13, 1131-1138. (4)Cui, C., Gan, L., Heggen, M., Rudi, S. & Strasser, P. Nature Mater 2013, 12, 765-771. (5)Gan, L., Cui, C., Rudi, S. & Strasser, P. Top Catal 2014, 57, 236-244.


This work was supported by U.S. DOE EERE award DE-EE0000458 via subcontract through General Motors and by Ernst Ruska Center for Microscopy and Spectroscopy with Electrons, Forschungszentrum Juelich GmbH, Germany. 

Fig. 1: Figure 1. Aberration-corrected STEM-EELS elemental profiles of dealloyed PtNi (a, d), PtNi3 (b, e) and PtNi5 (c, f) core-shell NPs, showing near-surface Ni-rich inner shells.2

Fig. 2: Figure 2. STEM images and EELS line profiles of size-selected spherical PtNi3 catalyst after stability test. Nanoporous particles formed at larger sizes (ca. 10 nm) and, consequently, lower near-surface Ni content as well as larger Pt shell thickness.3

Fig. 3: Figure 3. EELS elemental mapping of octahedral PtNi1.5 nanoparticles along (a) <110> direction and (b) <100> direction, showing that Pt segregated at the edges and corners whereas Ni segregated at the facets. (c) The revealed structural model. 4

Type of presentation: Oral

MS-1-O-3257 Atomic Level In-situ Characterization of Metal/TiO2 Photocatalysts under Light Irradiation in Water Vapor

Zhang L.1, Crozier P. A.1
1Arizona State University, Tempe, USA
liuxian.zhang@asu.edu

Photocatalysts have potential applications for solar fuel generation either through water splitting or CO2 reduction. It is now recognized that atomic level in situ observations are critical for understanding the structure-reactivity in photocatalysts in the presence of reactant and product species and during in-situ light illumination. TiO2 is a promising photocatalyst used for self-cleaning, pollutants degradation and water splitting etc. Metal particle co-catalysts such as Pt are coupled to the semiconductor to provide chemically active sites and attract excited electrons preventing charge recombination. Herein we use TiO2 as a model material to develop in situ photocatalytic experimental methodology and explore structure changes of metal/semiconductor photocatalysts. We employ a modified ETEM with a broadband light source to study the behavior of metal particles on TiO2 semiconductor surfaces under photoreaction conditions.
Well defined anatase nanoparticles were prepared following a hydrothermal method. Metal co-catalysts such as Pt, Ag and Cu were loaded onto the anatase nanobars using methods such as dry impregnation, photo-deposition and sputtering. An FEI Tecnai F20 ETEM was modified to allow samples to be illuminated with light with intensity up to 10 suns [1]. In situ analysis was performed tracking structure changes in photocatalytic vapor phase water splitting reactions. Ex-situ experiments were performed to compare or confirm in-situ observations under exposure to a 450W xenon lamp with a mirror reflecting 360nm to 460nm light. TEM images for ex-situ experiments were taken from an FEI aberration corrected Titan TEM.
The initial anatase particles shown in Figure 1a appear crystalline on the surface and the surface is smooth and atomically abrupt. Figure 1b shows a crystal after 40 hrs exposure to water and light without prior exposure to the electron beam. When the titania is exposed to light and water vapor, the initially crystalline surface converts to an amorphous phase one to two monolayers thick [2]. 5% wt Pt particles were loaded onto anatase nanoparticles and well dispersed through proper heat treatment. Figure 2a also shows initial Pt on TiO2 materials. After exposure to a xenon lamp in liquid water for 6hrs, Pt particles show significant sintering as shown in Figure 2b. Pt/TiO2 sample shows significant surface disordering (Figure 2b). In-situ experiments were performed to study the evolution of the Pt sintering, TiO2 surface roughening and the Pt/TiO2 interface changes. Other metal co-catalysts will also be discussed in the presentation.
References:
[1]. Miller, B.K.; Crozier, P.A. Microscopy and Microanalysis 2013, 19, 461-469
[2]. Zhang, L.; Miller, B. and Crozier P. A. Nano Lett. 2013 13 (2), 679–684


The support from US Department of Energy (DE-SC0004954) and the use of ETEM at John M. Cowley Center for HR Microscopy at Arizona State University is gratefully acknowledged.

Fig. 1: a) Initial anatase crystal at 150°C without exposure to water or light. b) After 40 hrs in 1 Torr H2O, 12 hrs light exposure.

Fig. 2: a) Initial 5%wt Pt on anatase particles, b)after ex-situ 6hrs exposure to 360nm-460nm light in liquid water; (with insertions low magnification images).

Type of presentation: Oral

MS-1-O-3315 Quantitative Z-contrast Imaging of Zeolite-supported Metal Clusters and Single-metal-atom Complexes With Single-Atom Sensitivity

Xu P.1, Yang D.1, Martinez-Macias C.1, Kistler J. D.1, Chotigkrai N.2, Gates B. C.1, Browning N. D.3
1Department of Chemical Engineering and Materials Science, University of California, Davis, Davis, CA 95616, USA, 2Department of Chemical Engineering, Chulalongkorn University, Bangkok 10330, Thailand, 3Fundamental and Computational Science Directorate, Pacific Northwest National Laboratory, Richland, WA 99352, USA
amy.pinghongxu@gmail.com

Supported metal catalysts, in particular zeolites and oxide supported noble metals, are widely applied in industrial processes, such as petroleum refining, automobile exhaust conversion and petrochemical conversion. A crucial challenge in studying these catalysts lies in structural nonuniformity, which hinders the tuning of catalytic properties for control of selectivity and activity. In this work, we report investigation of supported metal catalysts with small, uniform structures on highly crystalline supports to gain a fundamental understanding of supported catalysts. Here we present direct measurements of such catalysts, using aberration-corrected scanning transmission electron microscopy (STEM) at atomic resolution with single atom sensitivity. STEM is well-suited for characterizing these catalysts, as the high atomic number difference between the support and the supported metal provides a strong contrast in the STEM images. Quantitative analysis was performed on metal sizes, shapes and their bonding locations within the cavities of zeolite structures, which is essential in studying structure-property relationship. Our results also demonstrated that STEM technique is a very powerful when complemented by extended X-ray absorption fine structure and infrared spectroscopies for identification of ligands, determination of metal-metal distances and coordination number.

High electron-dose (105-108 e-/A2) STEM imaging was applied for these highly beam sensitive materials, with images taken quickly before significant destruction or modification of of structures. High signal-to-noise ratio for this approach provides the advantage of easy image interpretation. Results of characterization of a range of supported heavy metals, including iridium, platinum and rhodium supported on various zeolites will be presented. Zeolites were chosen as the support material because of their wide applications in industry and their high degrees of crystallinity, which provide well-defined structures for imaging. Our results demonstrated high contrast of the STEM images characterizing these catalysts, which is unattainable from previous studies, as exemplified by Figure 1 showing mononuclear iridium species supported on zeolite Y. Detailed determination of the configuration of the metal species in the structure and the interaction between these metal complexes and the ligands, including the support, will be presented.


This work was supported in part by the United States Department of Energy (DOE) Grant No. DE-3-BDOE797 through the University of California, Davis, the Laboratory Directed Research and Development Program (LDRD): Chemical Imaging Initiative at Pacific Northwest National Laboratory (PNNL), and the Environmental Molecular Sciences Laboratory (EMSL), a national scientific user facility at PNNL.

Fig. 1: Aberration-corrected HAADF-STEM image of zeolite Y supported iridium species. Bright features encircled are examples of iridium species, mostly mononuclear complexes.

Type of presentation: Oral

MS-1-O-3345 Microscopic experiments with radial junction solar cells based on silicon nanowires

Fejfar A.1, Hývl H.1, Ledinský M.1, Vetushka A.1, Kočka J.1, Misra S.2, Foldyna M.2, Lin Wei Yu2, Roca i Cabarrocas P.2
1Institute of Physics, Academy of Sciences of the Czech Republic, Prague, Czech Republic, 2Laboratoire de Physique des Interfaces et des Couches Minces (LPICM), Ecole Polytechnique, CNRS, Palaiseau, France
fejfar@fzu.cz

Microscopic measurements of Si thin films and nanostructures can provide interesting insights for their applications, e.g., for an operation of corresponding solar cells. This was the case for amorphous and microcrystalline Si films, but also for structures of polycrystalline Si on glass.

Radial junctions based on silicon nanowires (SiNWs) are an example of modern nanostructured solar cell designs with excellent light trapping and efficient photogenerated charge collection. A single pump-down process used to prepare a randomly grown matrix of SiNWs and conformal p-i-n radial junctions led to cells with efficiencies over 8% [1]. Considerable influence of irregularities in SiNWs lengths, orientations, shapes and mutual interaction on the photovoltaic action can be expected. Direct measurement of these effects requires microscopic measurements of photoresponse. This is possible using atomic force microscopy (AFM) with conductive cantilever which serves as a contact to individual radial junctions [2]. At the same time the cantilever can measure the local nanomechanical properties, including local stiffness of the wires, which can only sustain contact forces up to ~1 nN. Resulting conductivity maps show substantial variation of the local electronic properties. The AFM tip cannot reach deeper into the SiNWs matrix and correlation with scanning electron microscopy of the identical nanowires was sought in order to identify the reason for conductivity variations. The results are discussed in terms of random photodiode arrays connected in parallel with overall performance limited by weak diodes.

[1] S. Misra et al., Sol. Energy Mat. Sol Cells. 118 (2013) 90–95.

[2] A. Fejfar et al., Sol. Energy Mat. Sol. Cells. (2013) 228–234.


This work was supported by Czech Science Foundation projects 13-12386S,  and 14-15357S and the LNSM (Laboratory of Nanostructures and Nanomaterials) infrastructure framework LM2011026 supported by Ministry of Education, Youth and Sports of the Czech Republic.

Fig. 1: SEM view of the radial junction solar cells based on Si nanowires (top) and scheme of the conductive AFM characterization (bottom).

Fig. 2: Map of local current observed by C-AFM within 5x5 micrometers area (top) and local current superposed on the topography (bottom).

Type of presentation: Oral

MS-1-O-3374 In-situ Observations of Pt Nanoparticle Growth Using Aberration-corrected TEM and Graphene Liquid Cells

Ercius P.1, Yuk J. M.2, 3, 4, Park J.3, Kim K.2 3, Lee J. Y.4, Zettl A.2, 3, Alivisatos A. P.3
1NCEM, Lawrence Berkeley National Laboratory, Berkeley, CA 94720, 2Department of Physics, UC Berkeley, CA 94720, 3Materials Sciences Division, Lawrence Berkeley National Laboratory, Berkeley, CA 94720, 4Department of Materials Science and Engineering, KAIST, Daejeon 305-701, Korea
percius@lbl.gov

TEM aberration correctors now allow for structures to be investigated at atomic resolution with high contrast, which could greatly benefit in-situ observations of physical, chemical and biological phenomena. Recently, in-situ TEM observations of the growth of nanoparticles in liquids revealed new phenomena during the formation of colloidal inorganic particles, but the relatively thick SiN windows and even the liquid material trapped inside degrades TEM resolution and SNR [1]. The growth mechanisms of such materials could greatly benefit from high-resolution characterization although such experiments were difficult due to thick SiN windows. Graphene sheets have been successfully used as a single-atom thick substrate for high-contrast HR-TEM imaging, and its high flexibility, mechanical strength and impermeability allows the encapsulation of liquid under TEM vacuum conditions [2-3]. We introduce a type of liquid cell using graphene sheets to entrap a colloidal growth solution for in-situ HR-TEM imaging.

Graphene liquid cells (GLC) were created by superimposing two graphene sheets grown on separate grids. A Pt growth solution is pipetted on top of the opposing graphene substrates. The Pt growth solution intercalates between the graphene sheets and stays trapped after drying in air. Figure 1 shows a low-magnification TEM image of the encapsulated solution and an illustration of the GLC. For in-situ experiments, pockets of Pt-growth material are first identified using a low electron dose at low magnification. The electron dose is increased to 103 – 104 A/m2 at high magnification to reduce the Pt precursor and begin nanocrystal growth. We observed growth and coalescence of colloidal Pt nanoparticles at atomic resolution at 3.85 fps using TEAM I at 80keV and a Gatan US1000 CCD camera. The combination of a 2 atom thick membrane to contain a small amount of liquid and the chromatic aberration corrector (C-COR) provide unprecedented resolution during Pt nanoparticle growth.

Resolution and SNR are sufficient to image atomic columns, facets and twins of individual nanoparticles in each movie frame. Figure 2A) and B) show particle coalescence resulting in a FCC single crystal and a twinned FCC crystal, respectively. The twin structure remains for the duration of our observations.

Liquid encapsulated between graphene sheets provides an ideal in-situ system to study nanoparticle growth and coalescence with atomic resolution. The technique can be readily applied to study a diverse range of systems in-situ in a HR-TEM.

References:

[1] H. Zheng et al, Science 324, 1309 (2009)

[2] Z. Lee et al., Nano Lett. 9, 3365 (2009)

[3] J.M. Yuk et al., Nano Lett. 11, 3290 (2011)


GLC work supported on DOE contract no. DE-AC02-05CH11231. NCEM is funded by DOE contract no. DE-AC02-05CH11231.

Fig. 1: Low resolution TEM image of a 100 nm diameter liquid pocket and illustrations of a graphene liquid cell.

Fig. 2: Growth of Pt nanoparticles via coalescence along <111> directions imaged by HR-TEM in a GLC with relative times from the start of the reaction. A) Two particles coalesce to form a FCC single crystal. B) Two particles form a twinned nanocrystal as indicated in the included FFT pattern. 2 nm scale bars.

Type of presentation: Poster

MS-1-P-1458 EFFECT OF INTERGROWTH DISTRIBUTION AND PRESENCE OF DEFECTS ON CATALYTIC PERFORMACE OF MFI/MEL MATERIALS

Gomes M. E.1, Imbert F.2, Gonzalez G.1
1Lab. Materiales, Centro de Ing. Materiales y Naonotecnologia, Instituto Venezolano de Investigaciones Científicas, Caracas, Venezuela, 2Laboratorio de Cinética y Catalisis, Departamento de Química, Facultad de Ciencias, Universidad de Los Andes, Mérida 5101 - Venezuela
megomez@ivic.gob.ve

The impact of zeolites as catalyst for refining processes has been mainly associated to the regular distribution of channels and cavities that determines the distribution of product species on terms of geometry and size. In the search of new structures with new topologies, the controlled synthesis of zeolites with stacking defects forming intergrowth structures looks very promising.
The nature and distribution of intergrowth structures, and recurrently defect structures, like twins, as well as new types of structural imperfections can have a dramatic effect on the catalytic performance of these materials, due to the formation of new cavities in the intersections, resulting in new product distribution.
The structural disorder and its correlation with the catalytic activity have not been very well studied, neither the controlled synthesis of structural disorder. This is the fundamental base for the controlled design of microporous materials of disorder structures.
In the present work, the effect of different of intergrowths domain distribution of MFI/MEL and the presence of defect structures, on the catalytic performance of these materials has been studied using of n-decane hydroconversion as a model reaction.
The synthesis of the intergrowths of MFI/MEL was carried out by the combination of TBA and TPA organic molecules, TBA first added to colloidal silica solution and after 2 h of agitation TPA was incorporated, and hydrothermal treatment was followed at 150C and 90 C, obtaining proportions of 80MFI/20MEL and 70MFI/30MEL respectively.
The amount of intergrowth formation was determined by XRD, by fitting the experimental patterns to simulated patterns of intergrowth structures generated using the software DIFFaX.
The resulting distribution of products for MEL and 70MFI/30MEL and 80MFI/20MEL (Fig.1), clearly indicates that a new configuration of channels and cavities is present, forming larger cavities at the intersections that allows an increased amount of bulky products (di-branched) to be obtained. HRTEM was an essential technique to comprehend this behavior. For the intergrowth materials, a random distribution of intergrowths domains and numerous defects are present: vacancies, pore coalescence, that were responsible for the product distribution obtained in the catalytic reaction. (Fig.2). While the MEL presents more ordered and homogeneous structure (Fig. 3)
In this work, it was shown that the properties of a catalyst are governed by its microstructure and chemistry on an atomic scale, and electron microscopy methods were essential to directly analyze these properties.


Fig. 1: Fig. 1 Distribution of di-branched and mono-branched products reached at maximum isomerization in the catalytic reaction

Fig. 2: Fig. 2 HRTEM image of intergrowth 20%MEL/80%MFI

Fig. 3: Fig. 3 HRTEM image of MEL

Type of presentation: Poster

MS-1-P-1432 Morphological changes induced by reaction in a RuFe bimetallic catalyst: a BF-TEM and XED-Spectrum Imaging investigation

Teixeira-Neto E.1, Vignado C.2, Jordão E.2, Figueiredo F. C.2, 3, Carvalho W. A.3
1Laboratório de Microscopia Eletrônica, LNNano, CNPEM, C.P. 6192, 13083-970, Campinas - SP, Brazil, 2Faculdade de Engenharia Química, UNICAMP, C.P. 6066, Campinas, SP, 13083-970, Brazil., 3Centro de Ciências Naturais e Humanas, UFABC, Santo André-SP, 09210-170, Brazil.
erico.teixeira.neto@gmail.com

Catalyst deactivation is a major challenge for the catalysis community. The proposed mechanisms of catalyst deactivation include sintering, re-oxidation of metal components and surface reconstruction and mechanical deactivation through attrition. The bimetallic RuFe system has been investigated and employed as an interesting alternative catalyst in many applications. [1]
In this work we show results on the determination of morphological changes of an alloyed (1:1) RuFe/TiO2 (6% m/m) catalyst. This material was prepared by impregnating the TiO2 support with a solution of Ru3+ and Fe3+ and subsequently drying the suspension in a rotatory evaporator. The material was oxidized in ambient atmosphere at 600C for 2 h and then thermally processed in a H2-rich atmosphere at 400C for 1 h. The as prepared catalyst was used in the hydrogenation of dimethyl adipate in a Parr reactor at 250C and 50 atm of H2 for 15h. After the reaction, the catalyst was recovered and analyzed by TEM and XED-SI. The images and spectroscopic information shown here are representative of a detailed investigation of this system.
Figure 1 shows bright field (BF-TEM) images of RuFe particles deposited on the surface of TiO2 support. The as prepared bimetallic nanoparticles appear as dispersed dark hemispheres. A marked morphological change is observed in the catalyst after the reaction: in Fig. 1-A, small dark particles are seen embedded in an irregular shaped gray matrix. The inset in B show fragmented particles in detail.
In Figure 2, BF-TEM images and XED-SI chemical maps of the as prepared catalyst show correlated distributions of Fe and Ru, which is evidence of the formation of a RuFe solid solution. After the reaction, the Fe content is distributed throughout the gray matrix observed in Fig. 1 and Ru is concentrated at positions associated with the small dark gray particles seen in BF.
The decrease in the catalytic performance observed during the reaction can be attributed to the change in the distribution of metallic domains within individual particles. The initial morphology of the hemispherical RuFe solid solution particles changes to Ru-rich particles embedded into a matrix of iron oxide. This morphological description will provide new arguments to the understanding of the observed catalytic performance.

1. Nikolaos E. Tsakoumis, Magnus Rønning, Øyvind Borg, Erling Rytter, Anders Holmen, Catalysis Today 154 (2010) 162-182.


The authors thank Fapesp (2013/11298-0) and LME-LNNano-CNPEM (JEOL JEM 2100).

Fig. 1: BF-TEM images of RuFe particles deposited on the surface of TiO2 support. The as prepared bimetallic nanoparticles appear as dispersed dark hemispheres. A marked morphological change is observed after the reaction: in A, small dark particles are seen embedded in an irregular shaped gray matrix. The inset in B show fragmented particles in detail.

Fig. 2: XED-SI chemical maps of the as prepared catalyst show correlated distributions of Fe and Ru, which is evidence of the formation of a RuFe solid solution. After the reaction, the Fe content is distributed throughout the observed gray matrix (Fig. 1) and Ru is concentrated at positions associated with the small dark gray particles seen in BF.

Type of presentation: Poster

MS-1-P-1438 Plasmonic Photocatalyst Ag/AgCl Nanohybrids on Titanate Thin Film for Photocatalytic Application

Tang Y. X.1, Cheng Z.1, Dong Z. L.1
1School of Materials Science and Engineering, Nanyang Technological University, Singapore
zldong@ntu.edu.sg

Semiconductor photocatalysts have been extensively studied for the removal of organic compounds in waste water using solar energy [1, 2]. In this work, we demonstrate a novel plasmonic photocatalyst silver/silver chloride nanohybrids on the titanate thin film obtained via a facile and cost-effective approach [3, 4], which involves the following steps. Firstly, the sodium titanate thin film is prepared using a traditional hydrothermal method at 200 oC for 6h. Secondly, the Na+ ions in the interlayer of the titanate is replaced by Ag+ ions through an ion-exchange process. Then the obtained silver titanate readily reacts with HCl vapor to form the AgCl particles on the titanate thin films. Finally, the visible-light-driven plasmonic photocatalyst Ag/AgCl/titanate is obtained by partially reducing the Ag+ ions from the AgCl particles with the aid of Xe lamp illumination.

Typical FESEM and TEM images of the titanate film and the AgCl/titanate film are shown in Fig. 1 and Fig. 2 respectively. The as-prepared titanate film shows porous honeycomb-like features (Fig. 1a). Each honeycomb consists of 3~6 sided walls, inside which intertwined titanate nanowires are present. The diameter of a single nanowire is in the range from 40 to 50 nm. The X-ray diffraction pattern from the titanate film is shown in Fig. 1b, and the peaks are indexed as coming from orthorhombic titanate phase Na2Ti2O5. After the reaction with HCl, new peaks corresponding to the cubic AgCl phase are observed. Electron microscopy studies indicate that the dense AgCl nanoparticles are uniformly distributed on the surface of each titanate nanowire without agglomeration (Fig. 1c and 1d), and the particles size is around 50 nm (Fig. 2b). Fig. 3 shows that the as-prepared Ag/AgCl/titanate film photocatalyst exhibites higher activity in the visible region of the solar spectrum for the degradation of phenol solution, while the titanate thin film shows negligible activity for the phenol removal. This room-temperature synthesis route could be easily extended to prepare various solar light responsive semiconductors via metal ion exchange and gas reaction process for photocatalytic applications.

References

[1] X.C. Wang, K. Maeda, A. Thomas, K. Takanabe, G. Xin, J.M. Carlsson, K. Domen, M. Antonietti, Nat. Mater. 8 (2009) 76-80.

[2] X. Chen, S.S. Mao, Chem. Rev. 107 (2007) 2891-2959.

[3] Y. X. Tang, V. P. Subramaniam, T. H. Lau, Y. K. Lai, D. G. Gong, P. D. Kanhere, Y. H. Cheng, Z. Chen, Z. L. Dong, Appl. Catal. B: Environ., 106 (2011) 577.

[4] Y. X. Tang, Z. L. Jiang, J. Y. Deng, D. G. Gong, Y. K. Lai, H. T. Tay, I. T. K. Joo, T. H. Lau, Z. L. Dong, Z. Chen, ACS Appl. Mat. Interfaces, 4(2012) 438.


The authors thank the Environment and Water Industry Programme Office (EWI) under the National Research Foundation of Singapore (grant MEWR651/06/160) for the financial support.

Fig. 1: (a) FESEM image of the as-prepared titanate thin film showing honeycomb-like structure, (b) X-ray diffraction patterns of titanate film and AgCl/titanate film, (c) FESEM image showing the morphologies of the AgCl/titanate film, and (d) FESEM image showing uniform distribution of AgCl nanoparticles on titanate nanowire surface.

Fig. 2: The TEM images of (a) titanate nanowires, and (b) AgCl/titanate nanowires. The samples are obtained from the titanate film and AgCl/titanate film via ultrasonic treatment in water.

Fig. 3: Comparison of photocatalytic activityof titanate film and Ag/AgCl/titanate film samples for the photocatalyticdecomposition of phenol in water under the visible light illumination. Thelight intensity is around 115 mW/cm2.

Type of presentation: Poster

MS-1-P-1442 POLYMORPHS EVOLUTION DURING CRYSTALLIZATION OF BETA ZEOLITE

Sagarzazu A.1, Gonzalez G.1
1Centro de Ing. Materiales y Nanotecnología, Instituto Venezolano de Investigaciones Científicas, Caracas, Venezuela
asagarza@ivic.gob.ve

Molecular sieves are open-framework materials that can separate a mixture of different molecules on the basis of molecular size and shape. Among them zeolite beta  is     one of the most important with numerous industrial uses as a result of its structure of polymorphs  and a large pore openings of 7 Å.  It is formed by the intergrowth of two or three polymorphs (A, B and C) related by a combination of different stacking planes.  Searching for new framework topologies in such materials, with specific chemical and physical properties, a method to obtain different proportions of polymorphs has been employed in order to control the microstructure and chemistry on an atomic scale, therefore electron microscopy characterization is essential to directly analyze these structures.  Hydrothermal synthesis was used as crystallization method. The syntheses of intergrowths are based on the combination of different concentrations of   the organic template (TEAOH) and different silica/alumina ratios.     High resolution Scanning and Transmission Electron Microscopy, electron diffraction and X-ray powder diffraction techniques have been used for characterizing the structures.  To identified the different polymorphs the Multislice method was used to generate HRTEM images with specific crystallographic orientation for different focal and thickness series, using  the software Cerius2 and  the proportion of intergrowths was calculated by fitting the XRD data to the simulated patterns using DIFFaX,  a software design to calculate stacking defects in zeolites.
The mechanisms behind the formation of different proportions of polymorphs depended strongly on the synthesis parameters studied.  Although, most of the authors reported that the proportion of polymorphs in beta is 60%B-40%A, in the present work it was found this proportion fluctuates with the crystallization time for the different synthesis conditions employed, the lowest (51%B) and the highest (68%B) content of polymorph B was obtained for   SiO2/Al2O3 = 100 and TEA2O/ SiO2 =0.27 and 0.75 respectively. Therefore, it seems that high template content stabilizes polymorph B. Fig. 1 shows   HRTEM images for these samples.
Electron microscopy was an essential technique to understand the mechanism of formation of these materials and to understand their structure. The complementary use of all the microscopy techniques provided a wealth of unique information for the extensive characterization of these solids.


Fig. 1: Fig. 1 HRTEM images for beta zeolites obtained with different synthesis conditions: a., b. SiO2/Al2O3 =100,  TEA2O/ SiO2 =0.27,  2d. c., d. SiO2/Al2O3 =100, TEA2O/ SiO2 =0.75, 12d.

Type of presentation: Poster

MS-1-P-1461 EFFECT OF SYNTHESIS PARAMETERS ON THE MESOPOROUS STRUCTURE OF SBA-15 AND SBA-16

Soto D. A.1, Gomes M. E.1, Gonzalez G.1
1Laboratorio de Materiales. Centro de Ingeniería de Materiales y Nanotecnología, Instituto Venezolano de Investigaciones Científicas. Caracas, Venezuela.
damarysoto@gmail.com

SBA-15 and SBA-16 exhibiting arrangements of mesopores, have received particular attention for applications involving selective host-guest interactions or diffusion of large molecules.
SBA-15, prepared with the triblock copolymer PEO20PPO70PEO20 (Pluronic P123), consists of a mono-dimensional channel system distributed in a two-dimensional hexagonal structure. SBA-16 consists of two non-interpenetrating three dimensional
channel systems with spherical mesoporous cage-like cavities connected through windows.
A good understanding of how the synthesis conditions of these materials affect the meso- and macro-structure characteristics is important for their applications, since it allows to control  their properties.
In the present work the effect of variation of different synthesis conditions in the structure of both materials have been systematically investigated.
The synthesis was carried out in acidic medium from the triblock copolymers surfactants Pluronic F127 (EO106PO70EO106) and Pluronic P123 (PEO20PPO70PEO20) to obtain SBA16 and SBA 15, respectively  and TEOS was used as silica source. The synthesis parameters studied were temperature (70 to 110 ̊C), time (24 to 72 hours) and agitation. 
The synthesis carried out using Pluronic 123, resulted in an increase in pore diameter with temperature and time, from 5.6 nm at 70C, 24h to 8.3 nm at 110C, 72h and therefore  a decreases in mesoporous area and increase in  pore volume. It was observed that the microporosity was lost at high temperatures, and the mesoporous wall thickness decreases.
On the other hand, the synthesis carried out with Pluronic 127 showed a very small variation in  pore size with temperature and time. However, the mesoporous area increased, and the microporosity decreased with temperature showing some disorder on mesoporous arrangement at 100C, as it is observed by HRTEM.
Fig 1 shows HRTEM images of mesoporous materials SBA 15 synthesized at 90 °C, 48h and Fig 2 and 3 SBA 16 synthesized for 48h at 90 and 110 C, respectively, showing the disorder structure obtained at high temperatures.


Saidi Duno, Paola Patete, Edgar Cañizales for TEM facilities, Lisbeth Lozada for TEM sample preparation.

Fig. 1: SBA-15 synthesized with F123 at 90°C, 48h

Fig. 2: SBA 16 synthesized with F127 at 90°C, 48h.

Fig. 3: SBA 16 synthesized with F127 at 110°C, 48h.

Type of presentation: Poster

MS-1-P-1537 Characterization of silica-coated Au/Fe2O3 nanoaggregates

Krumeich F.1, Sotiriou G. A.2,3, Starsich F.2, Pratsinis S. E.2
1Laboratory of Inorganic Chemistry, ETH Zurich, Zurich, Switzerland, 2Particle Technology Laboratory, ETH Zurich, Zurich, Switzerland, 3Department of Environmental Health, Harvard University, Boston, USA
krumeich@inorg.chem.ethz.ch

The plasmonic characteristics of metals like Au or Ag dramatically change with particle size. The increased light absorption of nanoparticles (NPs) moreover depends on the wavelength and is maximized when the electrons in the conduction band are in their resonance state. Relaxation processes turn this oscillation energy into phonons, with an efficiency that depends on various parameters such as the particle size, shape and aggregation [1]. The thereby generated heat can be utilized for various applications, including cancer treatment. If Au NPs are selectively taken up by cancer cells, they can be activated photothermally by laser irradiation and the resulting heat can destroy these cells [2]. Here we report on the electron microscopy characterization of hybrid agglomerates (50 - 100 nm in diameter) consisting of SiO2-coated Fe2O3 and Au nanoparticles that show promising plasmonic and superparamagnetic properties [3]. This hybrid material was synthesized by enclosed flame spray pyrolysis, a very flexible and scalable technology [4].
TEM images (Figure 1) confirm that the Au/Fe2O3 NPs are indeed coated by an amorphous SiO2 shell which is ca. 2.5. nm thick here. The dark disks correspond to Au NPs with diameters between 10 – 40 nm, while the Fe2O3 NPs appear gray similar to the silica layer which encloses both types of NPs. The crystalline Au and Fe2O3 NPs furthermore show some lattice planes. STEM is employed for detailed characterization of these aggregates (Figure 2). In the HAADF-STEM (Z contrast) image, bright disks correspond to the Au NPs whereas faint gray areas indicate the presence of the less heavy scatterers (i. e., Fe2O3 and SiO2), as additionally confirmed by EDXS analysis of small areas (Figure 3a) and EDXS mapping (Figure 3b,c). Note that the crystalline Fe2O3 NPs are also detectable as areas showing lattice fringes in the PC-STEM image (Figure 2b) [5].
These results reveal that the Au and Fe2O3 NPs are predominantly located next to each other forming Janus-, or dumbbell-like nanoaggregates and that they are encapsulates by SiO2. The comprehensive characterization of the aggregates is important as the distance between the Au NPs determines the plasmonic interparticle coupling and this distance can be finely tuned by closely controlling the SiO2 shell thickness [3].

[1] P. K Jain et al. Accounts Chem. Res. 41 (2008) 1578.
[2] L. R. Hirsch et al., Proc. Natl. Acad. Sci. USA, 100 (2003) 3549.
[3] G. A. Sotiriou et al., Adv. Funct. Mater., 2014, http://dx.doi.org/10.1002/adfm.201303416.
[4] A. Teleki et al., Sens. Actuators, B, Chem. 119 (2006) 683; A. Teleki et al., Langmuir 24 (2008) 12553; G. A. Sotiriou et al., Adv. Funct. Mater. 20 (2010) 4250.
[5] F. Krumeich et al., Micron 49 (2013) 1.


Electron microscopy was performed at the electron microscopy center of ETH Zurich (ScopeM).

Fig. 1: TEM images of silica-coated Au-Fe2O3 aggregates revealing the coating of the Au and Fe2O3 NPs by an amorphous silica layer (microscope: Tecnai F30 (FEI), FEG, operated at 300 kV).

Fig. 2: HAADF-STEM (Z contrast) (a) and PC-STEM (phase-contrast) (b) images of the silica-coated Au-Fe2O3 aggregates (microscope: HD2700CS (Hitachi) with probe corrector (CEOS), cold FEG, operated at 200 kV).

Fig. 3: HAADF-STEM images (a,b) with the results of EDXS measurements of the indicated areas in (a) and EDXS elemental mapping (c) of (b). Au: green; Fe: blue; Si: red. (microscope: HD2700CS (Hitachi) with EDX spectrometer (EDAX Gemini system)).

Type of presentation: Poster

MS-1-P-1554 Photo-induced lattice accommodation in Ag/Cu composite nanoparticles

Yasuda H.1
1Research Center for Ultra-High Voltage Electron Microscopy & Graduate School of Engineering, Osaka University, Osaka, Japan
yasuda@uhvem.osaka-u.ac.jp

Nanoparticles exhibit specific structural and optical properties which are different from those of the corresponding bulk materials. The lattice softening is one of the specific properties originated from the shallow interatomic potential. On the other hand, localized plasmon in metallic nanoparticles is recently focused on the optical properties. The electric field induced by the localized plasmon may interact with the lattice vibration and enhance the lattice softening.
We confirmed in our previous research that a lattice accommodation takes place in a two-phase nanoparticle which has a lattice misfit. The lattice accommodation is induced not only on the interface between the two phases but also all over the nanoparticle. In such a lattice-accommodated two-phase nanoparticle, if only one phase is resonantly excited by the localized plasmon using well-defined photo-illumination, that is, the electron-phonon interaction is induced in the only one phase, how will the accommodated lattice behave to photo-illumination ?
In the present work, photo-induced lattice accommodation in Ag/Cu composite nanoparticles has been studies in situ by laser-coupled TEM with a double source evaporator, in order to see an electron-phonon interaction in the nanoparticles.
Fig. 1(a) shows a BFI of Ag/Cu composite nanoparticles and the corresponding DFI taken from Ag 111 reflection. Two kinds of morphologies are observed as shown schematically in the figure. One is core-shell structure (type A), and the other is particle-connected structure with a planer interface (type B). The amount of type A is larger than that of type B. Fig. 1(b) shows electron diffraction profiles from the nanoparticles before, during and after photo-illumination with the energy of 2.3 eV. All the diffractions are identified as Debye-Scherrer rings of the fcc silver and copper. The lattice constant of copper with and without photo-illumination are 0.370 and 0.366 nm, respectively. The changes in the lattice constant take place reversibly. The fact that no changes in the lattice constant are induced by photo-illumination in pure silver or copper nanoparticles denies an effect by the thermal expansion.
It was evident that photo-induced lattice accommodation takes place in Ag/Cu composite nanoparticles. Photo-illumination with the energy of 2.3 eV resonantly enhances the localized plasmon with the energy of approximately 2.0 eV in copper nanoparticles. An enhancement of the local electric field in copper nanoparticles may induce the lattice vibration and the subsequent lattice softening in the copper core region. Consequently, it is considered that the lattice of the silver shell accommodated by the copper core is relaxed to increase toward the lattice constant close to that of pure silver.


Fig. 1: (a)A BFI of Ag/Cu composite nanoparticles and the corresponding DFI taken from Ag 111 reflection. (b)Electron diffraction profiles from the nanoparticles before, during and after photo-illumination.

Type of presentation: Poster

MS-1-P-1562 HAADF-STEM of Road Aged Diesel Oxidation Catalysts

Ward M. R.1, Hyde T.4, Boyes E. D.1,2, Gai P. L.1,3
1Department of Physics and the York Nanocentre, University of York, UK, 2Department of Electronics and the York Nanocentre, University of York, UK, 3Department of Chemistry and the York Nanocentre, University of York, UK, 4Johnson Matthey Technology Centre, Sonning Common, UK
michael.ward@york.ac.uk

Diesel oxidation catalysts (DOCS) reduce the amount of pollutants emitted by diesel automobiles. It is well known that CO, NOx, hydrocarbons and soot are harmful to the environment (1). DOCs generally use nanoscale Pt nanoparticles supported on a γ-Al2O3 wash-coat which is held on a macro-scale monolithic structure in the car’s exhaust. Figure 1 shows a diagram of this type of structure. Pt is expensive and rare so it is a necessity to reduce the amount of Pt used in DOCs but also improve their long-term stability. One solution is to add Pd to Pt (1, 2). In addition to sintering, constituent atomic species in bimetallic nanoparticles can segregate over time. There are no detailed studies into this aging mechanism from real commercial DOCs. Here, we describe the aging mechanisms of a 57,000 km road aged bimetallic-DOC (PtPd-Al2O3) with HAADF-STEM (3).

A double aberration corrected JEOL 2200FS with Thermo Scientific Si(Li) window EDX was used (3). The specimens were prepared by first slicing and crushing monolith channels. The debris from this process was suspended in ethanol before being despotised onto a holey-C film Cu TEM grid (3). A fresh, unused monolith and the 57,000 km aged variant were supplied. Despite being used in a real automobile, there were no major issues with contamination in HAADF-STEM or HRTEM with the aged variant.

HAADF-STEM is ideally suited to image the DOC material due to the large difference in atomic number between the nanoparticles and the support. The nanoparticles were found to have grown by almost 400 % when comparing the aged DOC to the fresh DOC. The nanoparticles in the fresh DOC were 2.50 nm in diameter on average and 11.00 nm in the aged variant. In general, rounded surfaces of the nanoparticles are present in the fresh DOC. In the aged DOC, the rounded surfaces remain in the majority but the proportion of faceted nanoparticles had increased. Furthermore, HAADF-STEM and EDX was able to show that intensity variations in a minority of the nanoparticle images were attributed to the segregation of Pt and Pd (3). Pd was found at the edge of large segregated nanoparticles as predicted by theory (4). In some cases, the Pt/Pd appeared to form bands within the nanoparticles suggesting partial segregation which has not been observed from a commercial DOC before.

References

1. M. V. Twigg, Catal Today 117, 407 (2006).

2. A. Russell, W. S. Epling, Catal Rev 53, 337 (2011).

3. M. R. Ward, T. Hyde, E. D. Boyes, P. L. Gai, Chemcatchem 4, 1622 (Oct, 2012).

4. A. B. Shah et al., Nano Lett 13, 1840 (Apr, 2013).


The authors thank the EPSRC for support from critical mass grant EP/1018058/1.

Fig. 1: Structure of a DOC monolith (a) inside exhaust, (b) magnified view of the monolith channels and (c) magnified view of the wash-coat (Al2O3) and nanoparticles

Fig. 2: Typical HAADF-STEM images of the (a) fresh and (b) road aged DOC showing Pt-Pd nanoparticles

Type of presentation: Poster

MS-1-P-1586 Characterisation of mesoporous silica nanoparticles for chemotherapeutic applications

Young N. P.1, Huang X.2, Townley H. E.2
1Department of Materials, University of Oxford, Parks Road, Oxford, U.K. , 2Department of Engineering Science, University of Oxford, Oxford, U.K.
neil.young@materials.ox.ac.uk

Mesoporous silica nanoparticles present a wide range of applications, amongst which is an attractive means of delivery for pharmaceuticals within the body. Utilising the high internal surface area, tunable size and low toxicity of these nanomaterials, drugs may be targeted to sites within the body, yielded improvements over conventional treatment methodologies. In this study we have investigated the suitability of a number of different mesoporous silica nanoparticle structures for carrying a drug cargo [1]. Nanoparticles were characterised in terms of their physical parameters; size, surface area, internal pore size and structure. Additionally these were compared to properties specific to successful application in drug delivery; namely the loading and unloading profiles for a model therapeutic, and also the response of nanoparticles to conditions similar to those found inside the body. This data allows an informed decision to be made on the optimum nanoparticle structures required to maximise cargo capacity and optimise temporal control of the unloading. Controlled capping of the pores was also found to improve on the drug delivery capability.

Figure one shows SEM and TEM images of three of the classes of mesoporous silica nanoparticles investigated in the present study. These were named hexagonal mesoporous silica nanoparticles (a,b), blackberry-like mesoporous silica nanoparticles (c,d), and finally chrysanthemum-like mesoporous silica nanoparticles (e,f) on the basis of their structures. Overall the hexagonal particles were found to be ideally suited to drug delivery following confirmation of the properties described above. High-resolution TEM and tilt-series HAADF-STEM were used to fully characterise the internal pore structure and arrangement within these nanoparticles. Importantly this was found to be ordered with pore channels that were continuous throughout the volume of the nanoparticles, contributing to the high cargo carrying potential and efficient unloading profile.

[1] X. Huang, N.P. Young, H.E. Townley, Nanotechnology and nanomaterials 4 (2014) 1. DOI: 10.5772/58290


Fig. 1: A range of mesoporous silica nanoparticles imaged via SEM and TEM. Hexagonal (a) and (b), blackberry-like (c) and (d) and chrysanthemum-like (e) and (f).

Type of presentation: Poster

MS-1-P-1589 The effect of Ce0.8La0.2O1.9 support modifiers on the microstructure and N2O decomposition (de-N2O) performance of γ-Al2O3 supported Ir catalysts

Delimitis A.1, Pachatouridou E.1,3, Papista E.2, Iliopoulou E. F.1, Marnellos G. E.1,2, Konsolakis M.3, Yentekakis I. V.3
1CPERI / CERTH, Thermi, Thessaloniki, Greece, 2University of Western Macedonia, Kozani, Greece, 3Technical University of Crete, Chania, Crete, Greece
andel@cperi.certh.gr

N2O has been widely recognized as a hazardous greenhouse gas exhibiting 300 times higher Global Warming Potential compared to CO2, as well as an ozone layer destruction contributor. One of its major sources is fossil fuels and biomass combustion and, consequently, several methodologies have been considered towards its end-of-pipe emission control. Catalytic decomposition represents the most promising method, due to lower energy demand and cost. Currently, Ir-based catalysts have gained considerable interest as promising alternatives for de-N2O process. Enhancement of the Ir active phase intrinsic features via support-mediated promotional effects comprises the subject of the present study. In particular, the effect of Ce0.8La0.2O1.9-modified γ-Al2O3 support (AlCeLa) on the Ir nanostructural characteristics and its de-N2O activity is investigated, using a combination of electron microscopy (TEM, HRTEM) and image analysis methods.
The morphology of the unmodified 0.5 wt% Ir/γ-Al2O3 sample is depicted in Fig. 1. IrO2 catalyst adopts both a medium-size (up to 70 nm), crystalline rectangular particle morphology (a), as evidenced by the Selected Area Diffraction (SAD) pattern in (b), and a smaller and disordered particles one (c), densely aggregated on top of γ-Al2O3. Supporting Ir on AlCeLa results in the exclusive formation of larger size, highly crystalline IrO2 particles, as illustrated in Fig. 2(a) and (b), although the Ir loading is identical in both catalysts. The particles’ mean size is up to 500 nm in Ir/AlCeLa. Their high crystalline quality is presented in Fig. 3(a), where the edge of a IrO2 particle is shown, viewed along its [001] zone axis. Measurements of the lattice spacings resolved in the image resulted in d110=0.317 nm and d200=0.223 nm, in good agreement with their theoretical values. This is further confirmed by the Geometric Phase Analysis (GPA) results in Fig. 3(b). The strain map reveals a uniform distribution, even at the surface crystal edges, where any contamination by impurity elements or crystal defects formation may be more pronounced. Strain leaps, white arrowed in Fig. 3, were only measured at regions of crystal misorientations due to particle inclination.
The superior structural quality of Ir/AlCeLa catalyst was reflected in its outstanding ~100% and 90% N2O conversion records, in the absence and presence of O2, respectively. This is most probably a result of the trend of oxygen, formed by N2O decomposition, to desorb more easily from highly crystalline, clear IrO2 surfaces rather than from defected cites, mainly present in disordered, poorly crystalline small Ir particles. This inevitably leads to higher N2O decomposition activity in the former case, rendering Ir/AlCeLa a highly efficient de-N2O catalyst.


Financial support by the program “THALES” (MIS 375643), co-financed by the Greek Ministry of Education and Religious Affairs and the European Social Fund is acknowledged.

Fig. 1: (a) and (c) TEM images from the Ir/γ-Al2O3 catalyst, revealing the two distinct morphologies; (b) typical SAD pattern from the area in (a). Reflections attributed to IrO2 are denoted in (b).

Fig. 2: (a) TEM image and (b) [001] SAD pattern of a typical IrO2 particle in the Ir/AlCeLa catalyst. The difference in size and crystallinity is outlined.

Fig. 3: (a) HRTEM image, viewed along [001] and corresponding GPA strain map (b) from the edge of an IrO2 particle in the Ir/AlCeLa catalyst. A uniform distribution of strain is illustrated; peaks are only observed at regions of crystal inclination, as shown by the strain profile inset in (b).

Type of presentation: Poster

MS-1-P-1619 High resolution HAADF investigation of Ga droplets on Si(001) surfaces

Beyer A.1, Werner K.1, Stolz W.1, Volz K.1
1Philipps-Universität Marburg, Faculty of Physics and Materials Science Center, Marburg, Germany
andreas.beyer@physik.uni-marburg.de

The growth of III/V material on Si substrates opens a wide field of new applications and devices [1]. The initial stages of the nucleation and especially the type of element (III or V) it is started with may have a crucial impact on the interface structure and therefore a device´s performance [2]. In this study Ga was deposited on Si substrates without the presence of group V elements to investigate the processes occurring during these early stages of growth.
The samples were grown via metal organic vapor phase epitaxy in an AIX 200 GFR reactor. To investigate the impact on the morphology two different precursors, triethylgallium and trimethylgallium, were used and the growth temperature was varied between 400 and 500°C. Electron transparent samples were prepared along an <110> axis of the silicon by conventional mechanical grinding and final ion milling in a Gatan PIPS. The samples were characterized in a double C S-corrected JEOL JEM 2200 FS scanning transmission electron microscope (STEM) operating under high angle annular dark field (HAADF) conditions resulting in Z-contrast.
On the surface of the Si Ga droplets form which can be identified in HAADF images by their bright contrast, due to the higher atomic number of Ga with respect to Si (Fig. 1). The number of droplets clearly scales with amount of supplied precursor during the growth. Moreover, the droplets are not only confined to the surface but penetrate into the Si forming a pyramidal structure with boundaries on {111} lattice planes. Complementary energy dispersive X-ray measurements confirm that these pyramids contain Ga. The observed morphology can be explained by the fact that the liquid Ga etches the Si at the growth temperature. By addition of a precursor for group V after the Ga deposition, crystallization of the droplets can be enforced. Therefore, the droplets can serve as nucleation sites for the growth of low dimensional materials like nanowires.
This contribution will show how HAADF STEM can be used to investigate etching processes on an atomic scale.

References

[1] Liebich et al., Appl. Phys. Lett. 99, 071109 (2011).
[2] Volz et al., J. Cryst. Growth 315, 37 (2011).


Funding of the DFG in the framework of GRK 1782 is gratefully acknowledged.

Fig. 1: High resolution HAADF image of Ga droplet formed on the Si surface. The droplet penetrates into the Si and is framed by boundaries on {111} lattice planes which are indicated by broken lines.

Type of presentation: Poster

MS-1-P-1623 Plasmonic properties of hollow AuAg nanostructures by STEM-EELS

Genç A.1, Arenal R.2, 3, Patarroyo J.4, Henrard L.5, Gonzalez E.6, Puntes V.4, 7, 8, Arbiol J.1, 8
1Institut de Ciència de Materials de Barcelona, CSIC, Campus de la UAB, 08193 Bellaterra, Spain., 2ARAID Fondation, 50018 Zaragoza, Aragon, Spain, 3Laboratorio de Microscopias Avanzadas(LMA), Instituto de Nanociencia de Aragon (INA), Universidad de Zaragoza, 50018 Zaragoza, Spain., 4Catalan Institute of Nanotechnology (ICN), Campus de la UAB, Edifici Q (ETSE), 08193 Bellaterra, Barcelona, Spain, 5Department of Physics, University of Namur, rue de Bruxelles 61, B-5000 Namur, Belgium, 6Instituto Geofísico, Facultad de Ingeniería, Pontificia Universidad Javeriana, 110231, Bogota, Colombia, 7Universitat Autònoma de Barcelona (UAB), Campus de la UAB, 08193 Bellaterra, Barcelona, Spain, 8Institucio Catalana de Recerca i Estudis Avançats (ICREA), 08010 Barcelona, Catalonia, Spain.
agenc@icmab.es

The surface plasmon resonances are the collective oscillation of the conduction electrons of a metal excited by an electromagnetic radiation. During the last decade, plasmonic properties of metal nanoparticles have been attracted great interest owing to their potential applications in different fields such as electronics, photonics, biotechnology and Raman spectroscopy. Characteristics of the surface plasmon resonances, hence plasmonic properties, are known to be affected by the small modifications in size, shape and composition of the nanostructures, therefore it is essential to be able to directly correlate the surface plasmon resonances with the structural properties at the nanoscale. In this study, we have obtained the in-plane 2D distribution of the surface plasmonic resonances of hollow AuAg nanostructures [1], by means of low loss electron energy loss spectroscopy (EELS) in an aberration corrected scanning transmission electron microscope (STEM), equipped with a monochromator, with sub-eV and sub-nanometer resolutions. The studied complex nanoparticles are nanoengineered from solid Ag cubes to different hollow AuAg nanostructures such as nanoframes and multi walled nanoboxes [1]. We have investigated the local plasmonic property modulations on each nanostructure and correlated them their structural features. We have also correlated the obtained experimental results with models performed in the frame of discrete dipole approximation.

[1] E. González, J. Arbiol, V. F. Puntes, Science, 334, 1377 (2011).


Aziz Genç acknowledges the Ministry of National Education of Turkey for the PhD scholarship. 

Fig. 1: Figure 1: (a) background substracted EEL spectra extracted from the selected areas in the inset EELS SI. (b) is the plasmon energy map between 1.9 and 2.4 eV and (c) shows plasmon intensity maps between 1.8 and 3 eV with 0.2 eV windows (please note that the intensities are normalized for all maps).

Type of presentation: Poster

MS-1-P-1717 Enhanced Phytochemical approach for fabrication of Cobalt Nanoparticles

Debut A.1, Kumar B.1, Cumbal L.1
1Centro de Nanociencia y Nanotecnologia, Universidad de las Fuerzas Armadas ESPE, Av. Gral. Rumiñahui s/n Sangolqui, P.O. BOX 171-5-231B, Ecuador
apdebut@espe.edu.ec

Fabrication of cobalt nanoparticles using Passiflora tripartita var. mollissima fruit extract is an ecofriendly approach; produces various sizes and morphologies, including spherical, hexagonal and triangular. The P. tripartita fruit, known in Ecuador as “taxo & tumbo” belongs to the Passifloraceae plant family comprises around 530 species originated from temperate and tropical South America. Two different sonication conditions were employed for the synthesis of the cobalt nanoparticles and their growth recorded, in order to analyze the effect of the phytochemical synthesis and the ultrasonic irradiation on the morphology and size of the final product. The synthesized nanoparticles were characterized by U.V.-Vis, Dynamic Light Scattering, Transmission Electron Microscopy (TEM) with Selected Area Electron Diffraction (SAED) and X-ray diffraction. TEM analysis showed polydispersed nanoparticles with size ranges from 110 nm to 10 nm at different time interval and ultrasonic irradiation power. The X-ray diffraction analysis revealed the face-centered cubic geometry and SAED confirmed partial crystalline or amorphous nature of cobalt nanoparticles. Infrared spectrum measurements were carried out to hypothesize the possible biomolecules (flavonoid C & O-glycosides) responsible for stabilization the cobalt nanoparticles using P. tripartita. This simple, ecofriendly, and significantly low-cost protocol can be employed at ease and compatibility for pharmaceutical and biomedical applications.


This scientific work has been funded by the Prometeo Project of the National Secretariat of Science, Technology and Innovation (SENESCYT), Ecuador.

Fig. 1: Cobalt nanoparticles at different sonication conditions

Type of presentation: Poster

MS-1-P-1738 TEM/STEM investigations of silver nanoparticles embedded in titanium oxides for photochromic applications

Pailloux F.1, Diop D. K.1,2, Babonneau D.1, Simonot L.1
1Pprime Institute, UPR 3346 CNRS-Univ. Poitiers, France, 2LHC, UMR CNRS 5516, St Etienne, France
frederic.pailloux@univ-poitiers.fr

Nanocomposite films composed of Ag nanoparticules (NP) within a TiOx matrix present photochromic properties [1]. The permanent or reversible changes of color occurring under illumination by UV/visible lasers rely on the control of the localized surface plasmon resonance (SPR) of Ag NP. They result from the tuning of the NP size/shape distribution through photo-activated redox reactions occurring specifically with the TiOx matrix.

The morphology of Ag:TiOx nanocomposite is investigated by high-angle annular dark-field HAADF-STEM, structural informations are obtained by energy filtered electron diffraction (EFED) coupled with electron energy loss spectroscopy (EELS). Samples grown under different deposition conditions (TiOx thickness, O2 pressure in the chamber, ...) are investigated.

For TiOx grown under a metallic sputtering mode (low oxygen pressure), the STEM images reveal an homogeneous distribution of Ag nanoparticles with a rather large size distribution and various shapes. Their cristallinity is assessed by energy filtered electron diffraction. The porosity of the TiOx matrix is revealed by the HAADF images. Whereas the diffraction pattern of TiOx would suggest an amorphous structure the ELNES recorded on the O K and Ti L23 edges suggest the presence of a short range ordering of theTiO6 octaedra. The change of growing mode (under higher oxygen pressure) for TiOx, produces dramatic changes in the morphology of the nanocomposite: the Ag particles, if still present, have not been clearly resolved.

The influence of the thickness of the TiOx capping layer is investigated as well. It reveals that a threshold thickness exists below which the samples become sensitive to the electron beam, which promote morphological changes of the nanoparticles.

The morphological and structural insights are further compared with in-situ reflectance measurements [2].

References

[1] L. Nadar, N. Destouches, N. Crespo-Monteiro, R. Sayah, F. Vocanson, S. Reynaud, Y. Lefkir, J. Nanopart. Res. 15 (2013) 2048

[2] V. Antad, L. Simonot, D. Babonneau, Nanotechnology 24 (2013) 045606


This work is supported by ANR Photoflex project

Type of presentation: Poster

MS-1-P-1748 Nano-Branched Free Standing Gold Foils

Rodríguez-González B.1, Vázquez-Vázquez C.2, Ameneiro Prieto O.2, Correa-Duarte M. A.2
1CACTI, University of Vigo, E-36310 Vigo, Spain, 2Department of Physical Chemistry, University of Vigo, E-36310 Vigo, Spain
jbenito@uvigo.es

Herein we present results about the obtaining, characterization and formation mechanism of nano-branched free standing gold foils. Dendritic and nano-branched gold foils have potential applications as SERS active substrates in sensing.[1] Foils were prepared by a facile wet chemical synthesis method using gold salt and a reducing agent derived from the formaldehyde molecule. The overall method is cost-effective and allows for a facile transfer to a wide range of substrates for different sensing applications. As example, we have achieved the transfer of the nano-branched gold foil to paper, glass and silicon wafers.

In Figure 1(A) we show a low magnification TEM image of one of the obtained foils, it is clear the nano-branched structure displayed by the gold foil. This structure looks convenient for the flowing of liquids or gases through the large openings between the branches; this opens a wide field of versatile applications. Figure 1(B) shows the branches in more detail; they are formed by polycrystalline gold with multiple grain boundaries and twinning planes. The presence of those defects is a direct consequence of the foil formation mechanism.

In order to understand the formation mechanism, and the origin of the final structure of the foils, we have conducted electron microscopy studies over samples removed from the reaction vessel at different reaction times. These studies allow us to propose a multi-step formation mechanism. The first step is the reduction of the gold followed by the immediate formation of small gold nanoparticles or clusters in the reaction medium. These nanoparticles migrate to the liquid surface where the foil starts to develop due to an aggregation process. The arrangement in branches and gaps are due to a particular disposition of the particles along the aggregation process.

[1] Y. Wang, M. Becker, L. Wang, J.Liu, R. Scholz, J. Peng, U.h Gösele, S. Christiansen, D. H. Kim, and M. Steinhart, Nano Letters 2009 9 (6), 2384-2389


Fig. 1: (A) Low magnification TEM image of the nano-branched gold foil. (B) TEM image showing the grain boundaries and defects of the branches.

Type of presentation: Poster

MS-1-P-1779 In-situ observation of morphological changes of gold nanorods under near infrared pulsed laser irradiation

Matsumura S.1, Sumimoto N.1, Yamamoto T.1, Yasuda K.1, Niidome Y.2
1Department of Applied Quantum Physics and Nuclear Engineering, Kyushu University, Fukuoka, Japan, 2Department of Chemistry and Bioscience, Kagoshima University, Kagoshima, 890-8580, Japan
syo@nucl.kyushu-u.ac.jp

The controllable optical properties of metal nanoparticles have been an active research field because of their potential technological applications. Gold nanorods are ultrafine particles 20?150 nm in length and 5?20 nm in diameter. Their anisotropic shape gives rise to a surface plasmon (SP) absorption band corresponding to the longitudinal SP mode along the long axis in the near infrared region, in addition to a SP band for the transverse mode in the visible light region. The characteristic wavelength of the former SP mode is controlled by the aspect ratio of the rods. The longitudinal SP band is usually much more pronounced than the transverse SP band, and it is exploited for potential technological applications unique to gold nanorods. When irradiated with pulsed laser light, gold nanorods are deformed into different shapes, such as spheres, Φ-shape and elongated rods. Qualitatively, such deformations are considered to result from heating due to light absorption. However, the deformation behavior of gold nanorods remains largely unknown because most of the experiments have been performed ex-situ in irradiated aqueous solution. Recently, we erected a pulsed laser-light illumination system attached to a high-voltage electron microscope (HVEM) to observe light-induced behaviors of nano objects. In the present study, we observe in-situ the structural transformation in gold nanorods induced by irradiation. The JEM-1300NEF HVEM was operated at an accelerating voltage of 1250 kV, and laser pulses of λ= 1064 nm with 6-8 ns duration were simultaneously illuminated.

 Figure 1 reveals a sequential structural change in gold nanorods irradiated by 0, 1, and 7 laser pulses at 310 J/m2/pulse. One may notice that most of the gold nanorods have transformed their shape after a single laser pulse. However, additional laser pulses induce little further change in the nanorods, as shown in Fig. 1 (c). This attenuation of deformability can be explained in terms of the blue shift of the longitudinal SP band with the decrease of aspect ratio. HRTEM imaging reveals that the outer deformation is accompanied by total atomic restructuring in the nanorod interiors, involving generation and annihilation of planar defects, as shown in Fig. 2.


The present study was partly supported by Gant-in-Aid for Challenging Exploratory Research (#23656387) and for Scientific Research (B) (#25289221) from JSPS.

Fig. 1: In-situ observation of gold nanorods irradiated with laser pulses at 310 J/m2/pulse. Before irradiation (a), after 1 pulse (b), and after 7 pulses (c).

Fig. 2: HRTEM images of a gold nanorod before laser irradiation (a), after exposure to 1 pulse (b) and 2 pulses (c). Laser intensity is 425 J/m2/pulse.

Type of presentation: Poster

MS-1-P-1849 Atomic stoichiometry of the industrial-style Co-promoted MoS2 nanocatalysts

Zhu Y.1, Ramasse Q. M.2, Brorson M.1, Moses P. G.1, Hansen L. P.1, Kisielowski C. F.3, Helveg S.1
1Haldor Topsøe, Nymøllevej 55, DK-2800 Kgs. Lyngby (Denmark), 2SuperSTEM Laboratory, STFC Daresbury, Keckwick Lane, Daresbury WA4 4AD (UK), 3National Center for Electron Microscopy and Joint Center for Artificial Photosynthesis, Lawrence Berkeley National Laboratory, 1 Cyclotron Road, Berkeley, CA 94708 (USA)
yuaz@topsoe.dk

The knowledge of the position and the chemical identification of atoms at non-periodic sites in nanostructured catalysts is essential for the understanding of their catalytic functionality and can eventually lead to rational materials design. In the field of oil refining, transport fuels with ultra-low sulfur contents are produced by catalytic hydrodesulfurization (HDS) processes that rely on nanocrystalline MoS2–based catalysts.1 The HDS activity of MoS2 nanocrystals can be significantly promoted by transition metals such as Co and this catalytic activity enhancement is commonly attributed to the so-called “Co-Mo-S” phase, having Co located at the edges of the Mo plane of the MoS2 nanocrystals.2 However, detailed structural information regarding the Co promoter for the industrial catalysts is lacking.3

Recent advances in high-resolution (scanning) transmission electron microscopy ((S)TEM) imaging have opened up the possibility to study industrial-style MoS2 nanocatalysts with atomic-level resolution and sensitivity.4 By means of high-resolution TEM and high-angle annular dark-field (HAADF) STEM, the elemental distribution in unpromoted single-layer MoS2 nanocrystals was resolved5 and allowed for a distinction of the edge terminations.6 A combination of aberration-corrected HAADF imaging and electron energy-loss spectroscopy (EELS) is a promising approach for atomic chemical analysis; however, few characterizations have been possible owning to experimental challenges, such as the fine balance between interpretable signals and electron beam damage.

Here, we employed atomic-resolved HAADF-STEM and EELS at low primary electron energy to obtain the first site-specific identification of single-atom Co promoter and the associated S reconstruction in doped single-layer MoS2 nanocrystals (Fig. 1). Interestingly, single-atom Fe contaminants were unambiguously identified, competing with Co for the same sites at the S-edge. The present analytical capability of pinpointing local stoichiometry atom-by-atom with one atomic number sensitivity could be highly beneficial for improving the accuracy of the knowledge on complex nanostructures.

1 F. Besenbacher, et al., Catalysis Today 130, 86 (2008).

2 H. Topsøe, et al., Hydrotreating Catalysis (Springer, 1996).

3 O. Sorensen, et al., Applied Catalysis 13, 363 (1985).

4 C. O. Girit, et al., Science 323, 1705 (2009).

5 C. Kisielowski, et al., Angewandte Chemie, International Edition 49, 2708 (2010).

6 L. P. Hansen, et al., Angewandte Chemie, International Edition 50, 10153 (2011).

7 Y. Zhu, et al., (2014) submitted.


This work is supported in part by the UK Engineering and Physical Sciences Research Council, the Office of Science, Office of Basic Energy Sciences of the U.S. Department of Energy, the Danish Council for Independent Research (grant HYDECAT) and for Strategic Research (grant CAT-C).

Fig. 1: Fig. 1. a) High-resolution HAADF image of the S-edge of a monolayer Co-Mo-S nanocrystal. Corresponding EEL elemental maps of b) the combination of Mo (blue) and Co (red) and c) of S (yellow). d) Sum of EEL spectra, integrated over a 3 x 3 pixel window (probe size). f) Industrial-style Co-Mo-S ball model, with a side view of the Co-promoted S-edge.7

Type of presentation: Poster

MS-1-P-1853 Nanostructured cobalt ferrite for gas sensing

Leroux C.1, Madigou V.1, Giorgio S.2, Lopes-Moriyama A. L.3, Pereira de Souza C.3
1University of Toulon, CNRS, La Garde, France, 2Aix Marseille University,CNRS, Luminy, France, 3Universidade Federal do Rio Grande do Norte, Natal, Brazil
leroux@univ-tln.fr

Nowadays, measurement and control systems for pollutant and toxic gas emissions gain increasing importance in the frame of sustainable development. Although gas sensors devices are widely commercialized, they still suffer from drawbacks like lack of selectivity and stability. Parameters like time of reaction, time of recovery, reproducibility, working temperature, should also be considered. To overcome some of these disadvantages, nanostructured materials are investigated. The detection function of the sensing material is dependant of a high surface to volume ratio, but also to the exposed crystallographic facets. Nanoparticles present high surface to volume ratio, but tend to agglomerate. One way to overcome to some extent this phenomenon is to build hierarchical and hollow oxide nanostructures [1]. The transduction function of the sensing material is more linked to the composition and structure. It should be possible to tailor the reactivity and sensitivity of the sensing materials by controlling their composition, their structure, phase, shape, size, and size distribution [2]. Hence, we were interested in studying cobalt ferrites as nanoparticles or thin films for applications in gas sensors. The cobalt ferrite (CoFe2O4) attracts considerable attention due to its good chemical stability, mechanical hardness, magnetic behavior and catalytic activity [3-4].

Octahedron-like nanoparticles of CoFe2O4 were synthesised using a hydrothermal technique. Several microscopy techniques like SEM, conventional TEM coupled with EDS, high resolution TEM, environmental TEM, were carried out in order to understand the mechanisms involved in the growth of the grains and their reaction under gas. The particles have an octahedral shape, with sizes around 20 nm (Figure 1). Samples were observed in a TEM under 1mbar gas pressure and were submitted to H2 -O2 cycles, at ambient temperature. The {111} facets became more rounded under oxygen (Figure 2). Before that, the {100} facets extended which led to truncated octahedra. The same phenomenon was already observed in case of metallic nanoparticles [5].

References

[1] J.-H. Lee, Sensors and Actuators B 140 (2009) 319–336

[2] C. Wang, L. Yin , L. Zhang, D. Xiang and R. Gao, Sensors 2010, 10, 2088-2106

[3] D.S. Mathew, R.S. Juang, Chem. Eng. J. 129 (2007) 51–65.

[4] L. Ajroudi,S. Villain,V. Madigou,N. Mliki,Ch. Leroux, J. Cryst. Growth 312 (2010) 2465–2471.

[5] M. Cabié, S. Giorgio, C.R Henry, M. Rosa Axet, K Philippot, B. Chaudret,J. Phys. Chem. C 114 2160-2163, 2010


This work was done in the general framework of the CAPES COFECUB Ph-C 777-13 and ARCUS PACA BRESIL french-brazilian cooperation projects.

Fig. 1: The same CoFe2O4 nanoparticle viewed along a [110] direction, and viewed along a [111] direction after tilting, along with a drawing of the octahedron projection.

Fig. 2: Evolution of one CoFe2O4 nanoparticle under O2.

Type of presentation: Poster

MS-1-P-1856 Covellite nanocrystals and their evolution by addition of metals atoms: HRTEM and Exit Wave study

Bertoni G.1,2, Riedinger A.1, Xie Y.1, Brescia R.1, Pellegrino T.1, Manna L.1
1Istituto Italiano di Tecnologia, Via Morego 30, 16163 Genova, Italy, 2CNR-IMEM, Parco Area delle Scienze 37/A, 43124 Parma, Italy
giovanni.bertoni@imem.cnr.it

Copper sulphides (Cu2-xS) and related nanocrystals are promising candidates in optoelectronics device, due to their intrinsic p-type doping and tunability of their band-gap with stoichiometry.
In particular, the covellite structure (CuS) has one third of the Cu atoms in triangular coordination and two thirds of Cu atoms in tetrahedral coordination. At the same time, two thirds of S atoms form disulfide groups and one third are single sulfide ions.(1) At the tetrahedral sites, the Cu atoms are bound to the S atoms of the disulfide bonds. These different sites can be resolved in HRTEM (here we use negative C3 imaging conditions) or Exit Wave reconstructions (EWR) from side views (as [100] or [210] orientations) of Cu1.1S nanocrystals (see Figure 1). Indeed, at opportune values of defocus ΔZ, the S-S disulfide layers connected appear bright in the image, permitting to directly visualize them.
Our group demonstrated how the evolution from Cu1.1S (covellite type) to Cu2S (chalcocite) was accompanied by the red-shifting of the localized surface plasmon resonance (LSPR) generated by free holes generated in the valence band (and with a dominant in-plane mode).(2) The resonance gradually disappears as the Cu2S stoichiometry is reached (i.e. no copper deficiencies left). These states seem then to be linked to the disulfide bonds present of the CuS covellite structure (see Figure 1). In general, we expect by adding Mx+ atoms, the following transformation:

Cu1.1S + γMx+ + γe- → Cu1.1MγS ,

in which S(-1) is reduced to S(-2), possibly by breaking the S-S bonds, and the metal keeps its (x+) oxidation state. In this presentation we focus on Pd(II) doped CuS nanocrystals synthesized by chemical methods. Pd is added by using Palladium(II)-acetylacetonate plus ascorbic acid. We see how by increasing the amount of Pd atoms, the CuS structure is progressively lost, as the number of disulfide layers is gradually reduced, as can be seen from the HRTEM images (see Figure 2). Consequently, the density of holes in the valence band is lowered and the plasmon resonance is consequently red-shifted and reduced.
This is indeed a demonstration of the tunability of the LSPR with the amount of metal atoms in Covellite type nanocrystals.

[1] Pattrick R.A.D. et al. Geochim. Cosmochim. Acta, 1997, 61 (10), pp. 2023–2036
[2] Xie Y. et al. J. Am. Chem. Soc., 2013, 135 (46), pp. 17630–17637


European Union FP7/2007-2013 Grant Agreement 240111 (ERC Grant NANO-ARCH) and European Union FP7 Grant Agreement 312483 - ESTEEM2 (Integrated Infrastructure Initiative–I3).

Fig. 1: a) EWR phase together with the experimental image (ΔZ = -34 nm, C3 = -26 um) and the simulated image from [210] side view of a CuS nanodisk. At about -30/+20 nm defocus the intensity is transferred to the planes with Cu-S in tetrahedral coordination, giving a direct visualization of S disulfur planes (see the intensity profile in b).

Fig. 2: The nanocrystals maintain their hexagonal shape while adding Pd metal atoms. However, the number of S-S layers is progressively reduced as Pd increases, as well as the LSPR is reduced in intensity and red shifted.

Type of presentation: Poster

MS-1-P-1874 Plasmon Enhanced Fluorescence Imaging on Linear Arrays of Metal Half-Shells

Farcau C.1, 2, Astilean S.1, 2
1Institute for Interdisciplinary Research in Bio-Nano-Sciences, Babes-Bolyai University, Cluj-Napoca, Romania, 2Faculty of Physics, Babes-Bolyai University, Cluj-Napoca, Romania
cosmin.farcau@phys.ubbcluj.ro

Recent studies on the fluorescence emission of fluorophores located nearby metallic
nanostructures allowed the observation of peculiar phenomena such as modification of
radiative transition rates, enhancement of emission quantum yield, or directional emission.
The mentioned effects are induced by coupling between emitter dipole and surface plasmon
polariton excitations. Controlling these interactions with ordered metal nanostructures
bearing well-defined plasmon modes offers the means for advancing applications requiring
e.g., an enhanced emission, improved photostability, or larger FRET distances.

Here we discuss our studies of Surface Enhanced Fluorescence (SEF) on metal-coated
colloidal uni-axial arrays. These hybrid plasmonic-photonic crystals were obtained
by colloidal convective self-assembly (CSA) on DVD templates and metal film evaporation [1].
Their morphology is resolved by electron microscopy and atomic force microscopy (see Figure 1),
while their polarization-sensitive optical response is evidenced by transmission and
reflectivity micro-spectroscopy. SEF of fluorophores adsorbed on top of a spacing layer
are studied by both steady-state and time-resolved fluorescence. Furthermore, fluorescence
lifetime imaging (FLIM) is performed to highlight the correlations between topography / optical
response / SEF. These results on fluorescence emission enhancement, plasmon-controlled emission
polarization, lifetime modification and imaging can be of interest both fundamentally, for
better understanding of plasmon-coupled emission, but also from the application point of view, for the
design of sensors or other light-emitting devices.

[1] V. Saracut, M. Giloan, M. Gabor, S. Astilean, C. Farcau, ACS Appl. Mat. Interf. 2013, 5, 1362.


This work was supported by a grant of the Babes-Bolyai University, under the contract GTC_34057/2013.

Fig. 1: AFM image of the hybrid plasmonic-photonic crystal consisting of gold half-shells onto polystyrene microsphere linear arrays.

Type of presentation: Poster

MS-1-P-1906 EDXS on MoS2-base/Al2O3 HDS catalysts: A chemical distribution study of silicon

Angeles-Chavez C.1, Toledo-Antonio J. A.1, Cortes-Jacome M. A.1
1Mexican Institute of Petroleum, Molecular Engineering, Distrito Federal, MEXICO
cangeles@imp.mx

Electron microscopy (SEM and TEM) is a powerful tool for the characterization of a wide range of solid catalysts. Both microscope types give direct evidence of the morphology, chemical composition and crystalline structure in the different scales (micrometer to nanometer). The improvements of the instruments in the spatial resolution, energy resolution, efficiency of the detectors and data collection, has improved very much the quality of the obtained results and the silicon dispersion on gamma alumina particles used for the preparation of MoS2-based (hydrodesulfurization) HDS catalysts shows these new capabilities.
MoS2-base/Al2O3 HDS catalysts are widely used to remove sulphur from hard-to-desulfurize compounds such as 4,6-dimethyldibenzothiophene. Their catalytic performance is directly related to the dispersion of the MoS2 structure and the current scientific research is focused on achieving higher dispersion of Co-Mo-S active catalytic sites. In this work, we add silicon atoms on the surface of the alumina particles to modify their acidic properties and increase the dispersion of the Co-Mo-S structure. The Si atoms were aggregated to tri-lobular extruded of alumina by an incipient wet process using a silicon solution. Subsequently, the extruded were calcined and characterized by SEM and TEM.
The concentration of O, Al and Si in the sample, obtained by EDXS, was 46.90, 49.74 and 3.36 wt% in average, respectively. This chemical quantification indicates that the Si was integrated as SiO2 in the sample to a concentration around 7.2 %. The silicon permeation in the extruded was revealed by a composition study through the cross section of extruded. The result obtained is shown in Figure 1. A homogeneous concentration of silicon inside the extruded is observed. Therefore, this sample was the strongest candidate to impregnate the active phases (P, Co and Mo). Their dispersion was evidenced by concentration profiles (Figure 1) and chemical mapping, see Figure 2. Homogeneous dispersion of Co and Mo is appreciated in both results. However, the dispersion silicon was heterogeneous. This sample was subsequently sulfided to produce the Co-Mo-S structures. The result obtained is illustrated in Figure 3. MoS2 structures fully dispersed on the Al2O3-7. 2%SiO2 surface in HRTEM images is observed. Therefore, from these first results, the presence of SiO2 on gamma-alumina contributes to the formation of the MoS2 structures. However, still it is necessary improve the spreading of silicon in the extruded.


The authors acknowledge financial support to IMP through project D.00447.

Fig. 1: Concentration profiles in the tri-lobular extruded. Before impregnation (Si graph) and after impregnation (Co, Mo and P graphs).

Fig. 2: Chemical mapping of Si, Co and Mo in the tri-lobular extruded after impregnation.

Fig. 3: HRTEM image showing the MoS2 structures on the Al2O3-7.2%SiO2 surface.

Type of presentation: Poster

MS-1-P-1970 High resolution HAADF-STEM imaging of MoS2 nanolayers in industrial hydrotreating catalysts

GAY A. S.1, GIRLEANU M.2, TALEB A. L.1, BAUBET B.1, HUGON A.1, DEVERS E.1, ERSEN O.2
1IFP Energies Nouvelles - Rond point de l'échangeur de Solaize - BP 3 - 69360 Solaize (France), 2IPCMS-UMR 7504 CNRS-Univ. de Strasbourg - 23 rue du Loess - BP 43, 67034 Strasbourg cedex 2 (France)
anne-sophie.gay@ifpen.fr

Current severe environmental legislations constrain a strong decrease of sulfur concentration in fuels. Thus, the improvement of hydrotreating catalysts is of major importance. Co-promoted MoS2 based catalysts supported on alumina are known to be industrially used in the selective gasoline hydrodesulfurization (HDS) process. The challenge is to increase the selectivity of these catalysts. Catalyst performance (in particular selectivity) is suspected to be related to the local structure (ie 2D morphology) of the active phase, composed of MoS2 nanolayers promoted by cobalt. The equilibrium morphology is usually well predicted by theoretical approaches based on density functional theory (DFT) [1,2]. For since, it has also been visualized in model materials supported on gold or graphite by STM [3] and HAADF-STEM [4].
For that study, we observed MoS2 and CoMoS industrial catalysts, supported on delta-alumina by HAADF-STEM using a JEOL TEM 2100F with a Cs-corrected condenser. MoS2 and CoMoS catalysts were prepared by incipient wetness impregnation and sulfided under pure H2S at atmospheric pressure, either at 550°C or 700°C.
In MoS2 catalyst sulfided at 550°C, nanolayers present mainly a truncated triangle shape, in good accordance with DFT predictions [1]. Nevertheless, the morphologies are quite irregular : some nanolayers present a more isotropic shape. Some clusters are also observed. In MoS2 catalyst sulfided at 700°C (Fig 1), nanolayers are larger, well crystallized and morphologies are more homogeneous : mainly truncated triangles and some isotropic multi-facetted slabs.
CoMoS catalyst sulfided at 500°C present mainly hexagonal or irregular shape. No truncated triangle morphology is present. Some clusters are present. In addition, slabs are more stacked and aggregated than in non-promoted catalyst. At higher sulfidation temperature (Fig 2), the morphology of the nanolayers is homogeneous : all slabs are large, well crystallized, isotropic with many edges. No cluster is present. This observation is attributed to a combined effect of temperature and promoter edge decoration impacting the resulting 2D morphologies of CoMoS slabs [2].
In conclusion, this study highlights that HR-HAADF-STEM is a powerful technique to observe MoS2 nanolayers, even supported on alumina in industrial catalysts. In perspective of this work, changes of 2D morphology of nanolayers will be correlated to selectivity measured by catalytic tests.

[1] H. Schweiger, P. Raybaud, G. Kresse, H. Toulhoat. J. Catal. 207, 76-87 (2002).
[2] E. Krebs, B. Silvi, P. Raybaud. Catal. Today 130, 160-169 (2008).
[3] J. V. Lauritsen et al. Journal of Catalysis 197, 1–5 (2001)
[4] L. P. Hansen et al., Angew. Chem. Int. Ed. 2011, 50, 10153-10156


The authors thank P. Raybaud for helpful discussions about DFT.

Fig. 1: MoS2/alumina catalyst sulfided at 700°C under pure H2S

Fig. 2: CoMoS/alumina sulfided at 700°C under pure H2S

Type of presentation: Poster

MS-1-P-1973 Ti-assisted polarity inversion in ordered GaN nanorods investigated by transmission electron microscopy and density functional theory

Kong X.1, Li H.1,2, Albert S.3, Bengoechea-Encabo A.3, Calleja E.3, Draxl C.2, Trampert A.1
1Paul-Drude-Institut für Festkörperelektronik, Hausvogteiplatz 5–7, D-10117 Berlin, Germany , 2Institut für Physik and IRIS Adlershof, Humboldt-Universität zu Berlin, D-12489 Berlin, Germany, 3Dpto. Ingenieria Electronica, ETSI Telecomunicacion, Universidad Politecnica, Ciudad Universitaria, 28040 Madrid, Spain
x.kong@pdi-berlin.de

GaN nanorods are considered as promising building blocks for the realization of novel high performance light-emitting diodes due to their superior structural perfection. Ordered arrays of uniform GaN nanorods and axial (In,Ga)N/GaN heterostructures were achieved on Ti masked GaN (0001) templates by selective area growth using plasma-assisted molecular beam epitaxy. Here, we will report on the unexpected observation of polarity inversions accidentally found in those nanorods that we have analyzed by convergent beam electron diffraction and electron energy-loss spectroscopy (EELS). The inversion domains (IDs) with diameter of less than 10 nm cross the entire nanorod and originate at the homo-junction bounded by a stacking faults-like planar defect. Lattice imaging based on high-resolution transmission electron microscopy (TEM) is applied to determine an extra (0002) lattice plane in connection with this basal plane ID boundary. Spatially resolved EELS measurements reveal the presence of Ti impurities, which is possibly responsible for the formation of planar defects and the associated polarity inversion. In order to prove this assumption, to clarify the Ti lattice site occupation on the GaN (0002) plane and to explain the polarity inversion effect, we have performed first-principles total-energy calculations within the framework of density functional theory (DFT). They show that Ti monolayer adsorption on the GaN basal plane generates an energetically favorable atomic configuration that contains a planar defect resulting in a polarity inversion. The calculations perfectly match with the TEM observations. The influence of the polarity on the growth of axial (In,Ga)N nanorods is further discussed in detail.


Type of presentation: Poster

MS-1-P-1999 Systematic comparison of catalytic properties and applicability of d-elemental nanoclusters inside SWNT in-situ and on the atomic scale by means of AC-HRTEM

Zoberbier T.1, Chamberlain T.2, Biskupek J.1, Bichoutskaia E.2, Khlobystov A.2, Kaiser U.1
1Electron Microscopy of Materials Science, Ulm University, Germany, 2School of Chemistry, University of Nottingham, United Kingdom
thilo.zoberbier@uni-ulm.de

Catalysis on the nanoscale plays an important role on the one hand in terms of an enormous increase of efficiency/transformation rate on the other in the formation of defined nanostructures and functionalized materials. In both the morphologic properties of the catalyst plays an important role as well as the metal specific chemical properties. The possibility of controlled regulation and adaption of these parameters will lead to maximum gain and highest selectivity in chemical reactions and tap the full potential of catalysis in industry and the development of novel nanostructures. However neither the catalytic mechanisms are understood on an atomic level nor is there systematic studies on a fundamental base to understand why different metal type deviate in their applicability.

In our experiments we perform atomically resolved in-situ imaging of chemical reactions between d-elemental metals and a carbon environment inside SWNTs by means of aberration-corrected high-resolution transmission electron microscopy (HRTEM). The experiments aim to fundamentally understand and compare the catalytic properties of different catalysts by variation of a large range of transition-metals in equivalent experiments. This enables a detailed study of the processes essentially characterizing the aptitude and applications of the different metals, such as formation of metastable transient structures, formation of ordered carbon networks in different morphologies, annealing and reorganization processes and ability to assimilate carbon source material. Moreover the investigations allow a study of intermediates and the underlying chemical properties of Pi- and Sigma- bonding, metal-cohesive forces and solubility of carbon in metal.


Type of presentation: Poster

MS-1-P-2015 TEM study of TiO2 photocatalyst layers deposited on carbon nanosheet templates by atomic layer deposition

Kurttepeli M.1, Deng S.2, Verbruggen S. W.3, 4, Guzzinati G.1, Cott D. J.5, Lenaerts S.3, Detavernier C.2, Bals S.1, Van Tendeloo G.1
1EMAT, University of Antwerp, Groenenborgerlaan 171, B-2020 Antwerp, Belgium, 2Department of Solid State Science, Ghent University, Krijgslaan 281/S1, B-9000 Ghent, Belgium, 3Department of Bio-science Engineering, Sustainable Energy and Air Purification, University of Antwerp, Groenenborgerlaan 171, B-2020 Antwerp, Belgium, 4Department of Microbial and Molecular Systems, Center for Surface Chemistry and Catalysis, KU Leuven, Kasteelpark 11 Arenberg 23, B-3001 Heverlee, Belgium, 5Imec, 75, Kapeldreef, B-3001 Leuven, Belgium
mert.kurttepeli@uantwerpen.be

Due to its distinctive physical properties, chemical stability, bio-compatibility, non-toxicity and low cost, titanium dioxide (TiO2) is of great interest for a wide range of applications [1], [2]. The great potential of TiO2 nanostructures is obvious, but the desired physical and chemical properties of the materials will only be reached if a complete understanding of the relation between the activity and the structure of the materials has been obtained. In order to perform a detailed characterization of such nanostructures, transmission electron microscopy (TEM) is an ideal tool. Not only structural, but also chemical and electronic information can nowadays be obtained, even atomic column by atomic column [3], [4]. Nevertheless, one should take into account that conventional TEM images are only two-dimensional (2D) projections of three-dimensional (3D) objects. Therefore, TEM has been expanded to 3D, which is referred to as "electron tomography".
Hereby, we present results from different TEM characterization techniques to investigate the effects of annealing in helium environment on the structure of TiO2 layers deposited onto carbon nanosheets (CNSs) using atomic layer deposition (ALD). Using monochromated STEM-EELS, areas with TiO2 in anatase and amorphous form have been identified. From these maps, it is observed that the coating is mostly in anatase form, and there is only a low amount of amorphous TiO2 after annealing (see Fig 1). The graphite distribution map additionally indicated the presence of graphite throughout the layer. To investigate the 3-D structure of the material, HAADF-STEM electron tomography was applied (see Fig 2). The volume renderings proved both the homogeneity of the ALD coating throughout the CNSs layer, and the porosity of the complete film.
[1] A. Fujishima and K. Honda, Nature 238, 37 - 38 (1972).
[2] A. Kay and M. Grätzel, Solar Energy Materials and Solar Cells 44, (1996).
[3] K. W. Urban, Nature Materials 8, 260 - 262 (2009).
[4] D. A. Muller, Nature Materials 8, 263 - 270 (2009).


The authors acknowledge financial support from European Research Council and Sim-Flanders.

Fig. 1: HAADF-STEM image of the sample. Colored elemental STEM-EELS map with (B) anatase-TiO2 (red) and (D) graphite (green) is embedded. The amorphous-TiO2 elemental map is given in (C).

Fig. 2: Visualizations of the 3-D reconstruction of the sample depicted along different orientations are given in (A) and (B). A slice (orthoslice) through the 3-D reconstruction is presented in (C).

Type of presentation: Poster

MS-1-P-2323 Preparation and characterization of systems with plasmonic metal nanoparticles for fluorescence-lifetime imaging microscopy

Kokoskova M.1, 2, Pavlova E.1, Hromadkova J.1, Slouf M.1, Sloufova I.2, Sutrova V.2, Vlckova B.2, Kapusta P.3, Hof M.3, Michl M.4
1Institute of Macromolecular Chemistry, Academy of Sciences of the Czech Republic, Heyrovsky Sq. 2, 162 06 Prague 6, Czech Republic, 2Dept. of Physical and Macromolecular Chemistry, Charles University in Prague, Hlavova 8, 128 40 Prague 2, Czech Republic, 3J. Heyrovsky Institute of Physical Chemistry, Academy of Sciences of the Czech Republic, Dolejškova 2155/3, 182 23 Prague 8, Czech Republic, 4Faculty of Nuclear Sciences and Physical Engineering, Czech Technical University in Prague, V Holešovičkách 2, 180 00 Prague 8, Czech Republic
marketa.kokoskova@natur.cuni.cz

Luminescence of luminophores localized in a close proximity of plasmonic nanoparticles (NPs) such as Au or Ag are known to be completely or at least partially quenched [1]. However, in the case of larger distances from the metal surface a luminescence enhancement may be observed [2]. In order to investigate nanoparticle-luminophore distance effects, we focused our attention on reproducible preparation of homogeneous and well-defined model samples.

We studied systems with gold as well as silver NPs prepared by different techniques. The first set of samples were [substrate–AuNPs–spacer–luminophore] systems differing by AuNPs morphology. The substrates were microscopic cover glasses with constant thickness (specimens for fluorescence lifetime imaging microscopy, FLIM) and carbon-coated copper grids (specimens for TEM, controls). AuNPs were sputter-coated on the substrate; their morphology was controlled by sputtering time and subsequent thermal treatment. The spacer layer was created by thermal evaporation of carbon. The testing luminophores were widely used quantum dots and Ru(II) tris(2,2’-bipyridine); those were both sprayed and drop deposited onto the sample. The second set of samples comprised [substrate–AgNPs–luminophore] systems. In this case the AgNPs were prepared chemically by reduction of silver nitrate by hydroxylamine hydrochloride and in a form of a single aggregate deposited onto microscopic cover glass [3]. The testing luminophore /Ru(II) tris(2,2’-bipyridine) was drop deposited. The luminescent signal was measured immediately after deposition.

The size and the shape of Au and Ag nanoparticles were monitored by TEM and FEGSEM. We demonstrated that the combination of sputter coating and thermal treatment could yield NPs ranging from 5 nm up to several mm. The average size of the AgNPs was ~30 nm. The presence and homogeneity of luminophore on the surface was verified by TEM and EDX. Preliminary FLIM experiments of the systems with AuNPs showed quite inhomogeneous distribution of fluorescence lifetimes. Parallel TEM investigations suggested that the luminophores were deposited in multiple layers. Therefore, the different distances of luminophores from different layers might explain the observed distribution of FLIM signal. In the case of Ag NPs systems and drop deposition of luminophore, the surface-enhanced luminescence was observed.

References: [1] Geddes CD et al., Fluoresc. 2002, 12, 121., [2] Lakowicz JR, Anal. Biochem. 2001, 298, 1., [3] Sutrova, V. bachelor thesis, PřF UK, Praha 2013.


GACR P208/10/0941, TACR TE01020118 and GAUK 558213.

Fig. 1: TEM images of sputtered and thermally treated AuNPs (ts = sputtering time, thermal treatment: 450 °C/15 min).

Fig. 2: Fluorescence lifetime images and elastic scattered light images of AuNP/C/QD 510 system (A, B) and AgNP/Ru(bpy)3 system (C, D).

Type of presentation: Poster

MS-1-P-2031 Strain Concentration in Fivefold Twins

Yu R.1, Wu H.1, Zhu J.1
1Tsinghua University, Beijing, China
ryu@tsinghua.edu.cn

Multiple twinning is widespread in both natural and synthesis matter. The two types of multiple twinning, lamellar and cyclic, have attracted much attention due to their unique structures and properties. Lamellar twinning was shown to give a combination of high strength and toughness in copper [1], and highest creep resistance in titanium aluminide alloys [2]. Cyclic twinning, as another type of multiple twinning, occurs in an even wider range of materials, including not only inorganic small particles and thin films, but also proteins and virus [3]. The fivefold twinning is the most common form for multiple cyclic twinning [3].

The fivefold twinning has also attracted attention from the viewpoint of symmetry, which is an important concept in modern science. In fact, the fivefold rotational symmetry is geometrically forbidden in periodic crystals, although widely found in quasicrystals. Due to the geometrical imcompatibility, the fivefold twins have to be strained relative to the single-crystalline counterpart. Various models for the strain distribution have been proposed, including the linear homogeneous strain, angular and radial homogeneous strain, and the inhomogeneous strain models.

In this work, the atomic structure of the fivefold twins in diamond and silicon have been investigated by combining aberration-corrected transmission electron microscopy and first-principles calculations. In contrast to the strain distribution in metallic systems, which has small inhomogeneity, the strain in the fivefold twins of semiconductors depends significantly on the Pugh’s ratio of shear modulus to bulk modulus. For diamond with very high Pugh’s ratio, the strain is highly concentrated at the twin boundaries. Correspondingly, the frontier orbitals are located at the surfaces, in contrast to the case of silicon, where the frontier orbitals are close to the center.

References:

1. L. Lu et al., Science 304, 422 (2004).

2. F. Appel, and R. Wagner, Mater. Sci. Eng. R. 22, 187 (1998).

3. H. Hofmeister, Cryst. Res. Techno. 33, 3 (1998).


Acknowledgement: This work was supported by National Basic Research Program of China (2011CB606406), NSFC (51071092, 51371102, 11374174, 51390471, 51390475), and Program for New Century Excellent Talents in University. This work used the resources of the Beijing National Center for Electron Microscopy and Shanghai Supercomputer Center.

Fig. 1: (a) Aberration-corrected TEM image of and (b) strain distribution in diamond five-fold twins.

Type of presentation: Poster

MS-1-P-2094 Direct observation of Ti vacancies in Ti0.87O2 nanosheet using low-voltage monochromated TEM

Ohwada M.1, Kimoto K.1, Mizoguchi T.2, Ebina Y.1, Sasaki T.1
1National Institute for Materials Science, 2The University of Tokyo
kimoto.koji@nims.go.jp

Titania nanosheets [1] are two-dimensional single crystals of a titanium oxide with a thickness of one titanium or two oxygen atoms (Fig. 1a), and they show attractive material properties, such as photocatalytic reactions. The titania nanosheets are synthesized from a layered titanate K0.8Ti1.73Li0.27O4 through a soft chemical procedure (i.e., delamination), and the atomic arrangement of Ti-O layers in the parent crystal are basically preserved in the titania nanosheets. The nanosheets have the composition of Ti0.87O2, including cation vacancies at Li-substituted Ti sites of the parent crystal. In general, atomic vacancies affect the stability of crystal structures and material properties; therefore, it is important to reveal the atomic structure around Ti vacancies and their distribution in the nanosheets.

The observation of atomically thin materials requires not only high spatial resolution but also high sensitivity and low irradiation damage. We found that oxide nanosheets are substantially beam-sensitive, in contrast to a graphene and related materials. For instance we reported the topotactic reduction of a Ti0.87O2 nanosheet to Ti2O3 nanosheet [2].

We performed low-voltage and low-dose TEM observation using Titan3 at 80 kV with an image corrector (CEOS, CETCOR) and a monochromator, whose energy spread is 0.1 eV (FWHM). Attainable information limit under this condition was found to be 90 pm [3]. Figure 1b shows a high-resolution TEM image observed under underfocused condition [4]. The TEM image shows several bright areas as indicated by arrows, and we integrated these portions of the TEM image contrast (Fig. 2a). Based on the experimental results we constructed Ti vacancy structure models, and the atomic positions were optimized using first-principles calculation (the CASTEP code) as shown in Fig. 2b. The multislice simulation result based on the model successfully reproduces the experimental result (see Fig. 2c), and we found that the two oxygen atoms near the Ti vacancy are considered to be desorbed during the TEM observation [4].

[1] T. Sasaki, et al., J. Am. Chem. Soc. 118 (1996) 8329. [2] M. Ohwada, et al., J. Phys. Chem. Lett. 2 (2011) 1820. [3] K. Kimoto, et al., Ultramicrosc. 134 (2013) 86. [4] M. Ohwada, et al., Scientific Reports 3 (2013) 2801.


We thank Drs. Nagai, Ishizuka, Inoke, Lazar, Freitag, Sato and Suenaga for invaluable discussions. This work is supported by Nanotechnology Platform of MEXT and Research Acceleration Program of JSPS.

Fig. 1: A crystal structure model of Ti0.87O2 nanosheet (a), and a high-resolution TEM image of a Ti0.87O2 nanosheet. The TEM image is observed under underfocused condition. White arrows in Fig. 1b indicate several bright areas, suggesting Ti vacancies.

Fig. 2: (a) Experimental TEM image which is obtained as an average of the portions in Fig. 1b. (b) The atomic arrangement near the Ti vacancy optimized using the first-principles calculation. (c) The multislice simulation of a TEM image based on the optimized structure model.

Type of presentation: Poster

MS-1-P-2140 Structure of Plasmonic Nanocomposites obtained by Thermal and Laser Annealing of AlN:Ag Multilayers grown by Magnetron Sputtering

Bazioti C.1, Dimitrakopulos G. P.1, Kehagias T.1, Komninou P.1, Siozios A.2, Lidorikis E.2, Koutsogeorgis D. C.3, Patsalas P.1, 2
1Physics Department, Aristotle University of Thessaloniki, GR-54124 Thessaloniki, Greece, 2Department of Materials Science and Engineering, University of Ioannina, GR-45110 Ioannina, Greece, 3School of Science and Technology, Nottingham Trent University, NG11 8NS, Nottingham, United Kingdom
gdim@auth.gr

Localized surface plasmon resonances (LSPR) exhibited by plasmonic nanoparticles (PNPs) depend on PNP size, shape, distribution, and on the refractive index of the surrounding matrix. In this regard, efforts are undertaken to elucidate and control these parameters depending on growth conditions and post-growth treatment.
The structural properties of AlN thin films containing Ag PNPs were studied using TEM/HRTEM methods, and the results were correlated to the optical response. Magnetron sputtering (MS) was employed to deposit initially AlN:Ag multilayers with either amorphous (a-AlN) or nanocrystalline wurtzite-structured matrix (w-AlN) [1]. In one set of samples (series A), laser annealing (LA) using up to 700 mJ/pulse at 193 nm was employed in order to photomodulate the PNPs. In a second sample series (series B), flash thermal annealing (TA) was applied sequentially after MS deposition of each Ag layer, followed by LA in order to tailor the final microstructure.
In sample series A, LA dissolved the multilayer structure up to approximately half of its initial thickness, as shown in Figs. 1(a) and 1(b). This influence was more intense in the a-AlN case. Sample series B comprised just four 3 nm thick Ag interlayers embedded between AlN layers of 12 nm nominal thickness. TA led to complete structural reorganization resulting in dissolution of the layers and to a rather homogenous PNP distribution in the a-AlN case [Fig. 2(a)]. In the w-AlN case, an inhomogeneous PNP distribution was obtained, as larger PNPs were confined into two zones, one close to the substrate and one close to the surface. After LA, the homogenous PNP dispersion was destroyed in the a-AlN case, and larger PNPs were created [Fig. 2(b)]. For the w-AlN matrix, the PNP-zone close to the substrate was not dissolved, but still an improved PNP arrangement was obtained.
PNP enlargement by LA was described as an Ostwald ripening phenomenon. Larger PNPs were found close to the film surface, due to the enhanced Ag surface diffusivity. TA promoted Ag segregation, leading to even larger PNPs. The w-AlN crystallinity appeared almost unaffected, due to the strong ionic character of the atomic bond. Crystallinity was found to limit PNP enlargement as shown in Fig. 3, due to the resistance of the lattice to deformation. Overall it was demonstrated that controlled annealing processes can be employed to modulate the LSPR signal depending on the initial structure of the samples, as well as on the matrix crystallinity.


Work partially supported by the EU FP7 Project ‘SMARTRONICS’, Grant Agreement No 310229.

Fig. 1: Cross sectional bright field (BF) TEM overall images of multilayer a-AlN:Ag nanocomposites. (a) The nanocomposite prior to LA, comprising twenty Ag layers of 3 nm thickness with a 7 nm periodicity. (b) The nanocomposite after LA showing dissolution of the top half layers and ripening of PNPs.

Fig. 2: Cross sectional BF TEM overall images of a-AlN:Ag nanocomposites (a) after combined MS-growth plus sequential TA, and (b) after post-growth LA of the sample.

Fig. 3: HRTEM image showing Ag PNPs embedded in nanocrystalline w-AlN following LA treatment. Ag{111} and AlN{01.0} d-spacings are indicated.

Type of presentation: Poster

MS-1-P-2154 Detection of neuroendocrine tumor markers using nanostructured biosensors based on Au nanoparticle / Au film sandwich architecture

Boca-Farcau S.1, Farcau C.1, Astilean S.1
1Nanobiophotonics and Laser Microspectroscopy Center, Interdisciplinary Research Institute on Bio-Nano-Sciences, Babes-Bolyai University, 42 Treboniu Laurian St., 400271 Cluj-Napoca, Romania
sanda_c_boca@yahoo.com

Neuroendocrine tumors as are pheochromocytomas are dangerous tumors that require consideration in a large number of patients. Currently, the biochemical diagnosis of neuroendocrine tumors is based on plasma or urinary measurement of the direct secretory products of the adrenomedullary-sympathetic system or their metabolites, specifically catecholamines or their metanephrine derivatives. However, the techniques used for analysis of plasma free metanephrines, i.e. high-performance liquid chromatography (HPLC) or HPLC coupled with mass-spectrometry, are technically-demanding and time consuming which limit their availability [1]. Nano-biosensors based on colloidal gold or silver nanoparticles have proved their applicability for the accurate detection of tumor markers using Surface-Enhanced Raman Scattering (SERS) technique [2]. Recently special interest has been devoted to a type of biosensing platform made of individual gold nanoparticles (AuNPs) over gold films. This structure showed an increased sensitivity compared to self-assembled Au nanoparticles on other solid substrates [3], which was attributed to the additional electric field enhancement by the electromagnetic coupling between the nanoparticles and their supporting metal films.

In the present work we demonstrate a simple, fast and low-cost method for deposition of Au nanoparticles onto flat Au films with the aim of creating a SERS sensing platform with good enhancement and high signal reproducibility. Gold nanoparticles of tunable size and shape were synthesized by simple or seed mediated growth method. The structure and surface morphology of the nanoparticle-film sandwiched structure was characterized by scanning electron microscopy (SEM), atomic force microscopy (AFM), and UV–vis spectroscopy. Methanephrine metabolite was dropped into the sandwich structure and the SERS enhancement as a function of the deposited particle properties was measured using for excitation three laser lines (532, 633 and 785 nm). The obtained results demonstrate that the resultant Au-nanoparticle film exhibit noticeable SERS amplification of the adsorbed metabolite and can be used in the design of efficient, stable SERS-active substrates for the detection and identification of specific tumor markers.

References:

1. M. Procopiou, H. Finney, S.A. Akker, S.L. Chew, W.M. Drake, J. Burrin, A.B. Grossma, Eur J Endocrinol. 2009 ,161,131-40.

2. H. Hwang, H. Chon, J. Choo, J.-K. Park, Anal. Chem., 2010, 82, 7603-7610.

3. C.L. Du, C.J. Du, Y.M. You, C.J. He, J. Luo, D.N. Shi, Plasmonics, 2012, 7,475-478.


This work was financially supported by Babes-Bolyai University, Cluj-Napoca, Romania under the Research Grant for Young Scientists, contract GTC-UBB No. 34056/2013.

Type of presentation: Poster

MS-1-P-2165 Diffusion effects investigations of self-organized Gold nanostructures on Ge(001) surface by Electron Microscopy

Jany B. R.1, Nikiel M.1, Szajna K.1, Indyka P.2, Krok F.1
1Jagiellonian University, Marian Smoluchowski Institute of Physics, Reymonta 4, PL30059 Krakow, Poland, 2Jagiellonian University, Faculty of Chemistry, Ingardena 3, PL30060 Krakow, Poland
benedykt.jany@uj.edu.pl

The self-organized gold nanostructures on Ge(001) surface are currently of special interest due to their applications for mono-molecular electronic devices [1,2]. The understanding of electrical as well as physical properties of the system is of great importance.
The Ge(001) substrate samples were cleaned to achieve atomically flat terraces by low energetic ions bombardment and by annealing. Next, 6 ML of Au was deposited by the Molecular Beam Epitaxy in room temperature. Later, the sample was post-annealed to temperature from 473 K to 770 K. The gold self-organizes to create island structures on Ge surface as depicted in Figure 1.
The morphology of Au/Ge(001) samples was measured for different post-annealing temperatures with SEM FEI Quanta 3D FEG. The island surface density and their sizes were measured providing the information on surface diffusion effects. The autocorrelation analysis shows that there exists preferred island orientation along crystallographic directions on the substrate surface.
Cross sections from the Au/Ge(001) samples were prepared using FIB technique for transmission electron microscopy measurements conducted with TEM FEI Tecnai Osiris 200 kV equipped with Super-X EDX detector. The TEM measurements show that some island are submerged in germanium substrate. The chemical composition of the islands was mapped by the STEM/EDS measurements. This uncovered core/shell structure of the islands, with germanium shell on top. The crystalline nature was first studied by Selected Area Electron Diffraction (SAED) diffraction and Dark Field imaging. Later, detailed investigations were performed by Nano Beam Diffraction (NBD) measurements in STEM micro-Probe. This showed differences in crystalline structure of the islands.
The electron microscopy gives the possibility to fully study the creation dynamics and to completely characterize the fabricated nanostructures. Surface diffusion effects are investigated by the SEM as well as effects of diffusion processes into the bulk Ge crystal are measured by the TEM cross sections. This gives the unique scientific possibilities to fully investigate the evolution of the self-organized systems. The results and used experimental techniques will be discussed.

[1] C. Joachim et. al., Nature 408, 2000
[2] M. Wojtaszek et. al., Advances in Atom and Single Molecule Machines, Vol.1, 2012


The authors gratefully acknowledge financial support from the Polish National Science Center, grant no.DEC-2012/07/B/ST5/00906. The research was carried out with equipment purchased with financial support from the European Regional Development Fund in the framework of the Polish Innovation Economy Operational Program (Contract No. POIG.02.01.00-12-023/08).

Fig. 1: Self-organized gold island grown on Ge(001) surface, In center: TEM image of cross section through the gold island, Top: SEM secondary electron image shows the gold islands on the germanium surface.

Type of presentation: Poster

MS-1-P-2172 Equilibrium shape changes of PtCu alloy nanoparticles across the order-disorder transformation - An in-situ TEM study

Chatterjee D.1, Ravishankar N.1
1Materials Research Centre, Indian Institute of Science, Bangalore, India
dipanwita.chatterjee06@gmail.com

Platinum alloy nanoparticles find application in the field of electrocatalysis and gas phase reaction catalysis. The activity and stability of the catalyst
nanoparticles depend on the exposed crystal facets. Hence, the shape changes of the catalyst nanoparticles affect their surface properties like adsorption of gases and
hence the catalytic activity. Shapes assumed by the particles can be either kinetic or thermodynamic. If the particles are equilibrated under fixed conditions of
temperature, pressure, volume and composition, they attain their thermodynamic or equilibrium shape which is given by the Wulff construction. While the shape changes
with temperature has been studied for monometallic particles, there are very few studies on alloy nanoparticles. In particular, the shape changes associated with
changes in ordering of nanopartciles has not been investigated.

We have chosen a class of bimetallic A50B50 type alloy system which undergoes order to disorder phase transformation with the increase of temperature. The equilibrium
shape changes of the alloy systems having a varied range of heat of mixing across the transition from B2 ordered to A2 disordered structure and L1o ordered to A1
disordered structure have been studied theoretically. Shape change in terms of change in the area of the exposed facet was observed with changing degree of order in
the alloy system (Figure 1).

Experimentally equilibrium shape changes have been observed for PtCu system which undergoes transformation from ordered rhombohedral structure to disordered cubic
structure using in-situ heating techniques in transmission electron microscope (TEM). The system has been designed such that the alloy nanoparticles nucleated on the
MgO cubes appear edge-on on the face of the cube when tilted to its [001] zone axis and the shapes of the equilibrated alloy particles at different temperatures
implying different order parameters have been imaged at high resolution. The facet lengths of the equilibrated particles were observed to change monotonically with
decreasing order parameter that corresponds to the facet area change observed in the theoretical study (Figure 3).
The equilibrium shape change with the degree of order in an alloy system has been observed for the first time and its implication is immense, not only in terms of its
fundamental basis but also in terms of the application of the alloy particles undergoing order-disorder phase transformation where the property of equilibrium shape
change with degree of order could be exploited in the field of catalysis.


Financial support from DST is acknowledged. The electron microscopes are a part of the Advanced Facility for Microscopy and Microanalysis at IISc.

Fig. 1: Figure 1. Theoretically derived equilibrium shapes for B2 ordered to A2 disordered phase transition wherein the relative area of 110 and 100 facets vary with decreasing order parameter. The effect of different heat of mixing is also shown in this figure.

Fig. 2: Figure 2. (a) Low magnification image of MgO cubes with PtCu alloy nanoparticles nucleated on the cubes. (b) and (c) faceted ordered alloy nanoparticles appearing edge-on on MgO substrate.

Fig. 3: Figure 3. PtCu alloy nanoparticle on MgO substrate equilibrated at three different temperatures showing changes in the facet length in two-dimensional projection which translates to facet area change in the three-dimensional particle.

Type of presentation: Poster

MS-1-P-2178 PtBi Alloy Nanoparticles on Nitrogen-Functionalized Reduced Graphitic Oxide Support for Electrocatalysis

Tripathi S.1, Ravishankar N.1
1Materials Research Centre, Indian Institute of Science, Bangalore, India
shalinitripathi2307@gmail.com


The use of Pt-based electrocatalyst for methanol and formic acid oxidation reactions suffers from coarsening and deactivation of catalyst due to adsorption of CO. In this context, alloying Pt with a non-noble element has been shown to reduce the poisoning effect of CO significantly. Also,the presence of a conducting catalyst support can inhibit the catalyst coarsening without compromising the facile electron transfer. In this work, we report a microwave-based wet chemical approach for alloying Pt nanoparticles with Bi on a reduced graphitic oxide (RGO) support, which enhances the stability of the catalyst. Furthermore, we illustrate a way to improve the electron transfer by increasing the conductivity of the support through nitrogen functionalization of RGO. This wet-chemically synthesized graphitic oxide sheets facilitated the doping of the nitrogen at a very low temperature compared to the other reported physical processes, which can be attributed to the numerous localized defects in the GO sheets. Detailed transmission electron microscopy has been applied to understand the underlying mechanism in order to engineer the composition and morphology of the catalyst alloy nanoparticles. Our study also suggests that the alloying happens by nucleation of Bi on pre-formed Pt nanoparticles. Furthermore, the lower melting point of bismuth and its higher diffusivity facilitates the formation of intermetallic PtBi structure at such a low temperature. Thus, a microstructure-based thorough mechanistic understanding of the catalyst fabrication presented in this work imparts a control over shape, size and composition of the catalyst.


TEM facilities provided by Advanced Facility for Microscopy and Microanalysis (AFMM), and XPS facility of CENSE, Indian Institute of Science, Bangalore, India.

Fig. 1: Fig1. (a) LM and (b) HRTEM images of Pt nanoparticles, (c) PtBi alloy nanoparticles on RGO support; (d) SAED showing PtBi phase; (e) BF image of Pt NPs on NGO support; (f) shows the corresponding DF image; (g) LM and (h) HRTEM showing ordered PtBi phase over NGO support

Fig. 2: Fig.2: (a) High resolution XPS spectrum of Pt4f and (b) Bi4f from fabricated catalyst; (c) quantification shows a 1:1 atomic ratio of Bi and Pt

Fig. 3: Fig.3: Schematic showing the MW-based mechanism of selective heterogeneous nucleation for fabrication of alloy catalyst on support

Type of presentation: Poster

MS-1-P-2194 Nucleation texture of metal nanoparticles on amorphous substrates

Chatterjee D.1, Akash R.1, Kamalnath K.1, Ravishankar N.1
1Materials Research Centre, Indian Institute of Science, Bangalore, India
dipanwita.chatterjee06@gmail.com

Crystals nucleating homogeneously tend to adopt their
equilibrium shapes to minimise the barrier for
nucleation. For heterogeneous nucleation on a substrate,
the Wulff shape of the crystal itself is translated,
rotated and truncated by the substrate. The orientation
and level of truncation determimes the volume of the
so-called Winterbottom shape. We have calculated the
preferred orientation of heterogeneous nucleation on an
amorphous (isotropic) substrate by assuming that the
wetting of the solid nucleus on the substrate is
constant for the different orientations of nucleation.
Under the given conditions, the preferred orientation of
nucleation is the one for which the exposed volume of
the crystal on the substrate is minimum, as for such an
orientation the nucleation barrier is the minimum.

Theoretical calculations for obtaining minimum energy
Winterbottom shapes of nuclei of FCC metal at their
preferred orientations for a range of wetting conditions
have been done and the results are shown in Figure 1.
Experimentally, we have attempted to estimate the
orientation of nucleation of few hundreds of FCC metal
nuclei in order to statistically conclude the preferred
direction of orientation of heterogeneous nucleation on
an amorphous carbon substrate. Precession Electron
Diffraction (PED) technique is being used to scan over
regions containing a good number of nuclei and obtain an
orientation map from which the nucleation orientation of
the metal nuclei is to be determined.

Very fine nuclei of Au or Pt nanoparticles have been
nucleated on functionalized amorphous Carbon coated
Copper grid by microwave reduction of the precursor
salts in ethylene glycol medium. Electron diffraction
pattern obtained from such fine nuclei do not contain
enough number of spots for a reliable indexing of the
pattern using standard diffraction patterns for the
particular metal. So, the orientation map obtained from
the sample has a very low reliability index. For
optimization of conditions to obtain reliable
orientation mapping, PED scan on homogeneously nucleated
Au particles of around 8 nm diameter [inset of Figure 2
(a)] have been carried out and the resulting orientation
map has been shown in Figure 2(b). Here the 8 nm
particles could be resolved properly, as can be seen in
the virtual bright field image of the scanned area in
Figure 2(a) but the reliability index is poor because of
the polycrystalline nature of the Au nanoparticles.
Results on the nucleation texture of different
nanoparticles will be presented with detailed analysis
of the suitable microscopy conditions required for the
same.


Financial support from DST is acknowledged. The electron microscopes are a part of the Advanced Facility for Microscopy and Microanalysis at IISc.

Fig. 1: Figure 1. Preferred orientation of nucleus of FCC crystal heterogeneously nucleating on amorphous substrate at different wetting condition defined by Δs. Δs is related to the difference in the subtrate-vapour and substrate-particle interfacial energy.

Fig. 2: Figure 2. (a) Virtual bright field image of the Au nanoparticles on amorphous Carbon generated after the PED scan, inset showing a low magnification image of the area scanned, (b) Orientation map, different colours designating definite directions of the crystals in the scanned area.

Type of presentation: Poster

MS-1-P-2222 Quantification of PtIr Catalyst Nanoparticles Using ADF STEM

MacArthur K. E.1, Jones L. B.1, Lozano-Perez S.1, Ozkaya D.2, Nellist P. D.1
1Department of Materials Science, University of Oxford, Oxford, UK , 2Johnson-Matthey Technical Centre, Reading, UK
katherine.macarthur@materials.ox.ac.uk

Bimetallic catalyst nanoparticles for hydrogen fuel cell applications contain less Pt and exhibit higher catalytic activity than pure Pt particles.1 We have analysed Pt/Ir alloy particles which have shown improved resistance to CO poisoning. In order to understand these systems further it is necessary to examine their 3-dimensional atomic structure. Quantification of annular dark-field scanning transmission electron microscope images uses atomic resolution images as data sets for extracting composition and thickness information. Calculating the scattering cross-section (CS) of each atomic column provides robustness to many experimental parameters2 providing greater flexibility when imaging such challenging samples.
An automated code3 carries out detector normalisation,4 peak finding, background subtraction and column wise integration making it now possible to analyse and compare many particles. The measured CSs are assigned to atom counts through comparison with a simulation library. Simulations were carried out using the QEP μSTEM software matching experimental conditions of a 300kV microscope,5 with detector angles 34.9-190mrad and a probe convergence angle of 20.2mrad, with 30 phonon configurations. Due to their proximity in atomic number Pt and Ir are indistinguishable below 14 atoms thickness, Figure 1. Above 15 atoms the CS trend of each species begins to diverge; this is also the thickness where the accuracy of the atom-count assignments is greater than ±1 atoms making the error too large for accurate counting. Armed with the number of atoms within each column and their x-y coordinates, we can reconstruct the 3-dimensional structure from a single experimental image by assuming no vacancies and minimising surface steps.
To validate the experimental nanoparticle structure, the theoretical Wulff shape for a Pt/Ir alloy particle, was constructed using the Wulffman code,6 Figure 2, and orientated to a comparable viewing direction. The energies of the different alloy surface facets were assumed to be a linear combination of the pure elements.7 The considerable similarity between experiment and the Wulff shape demonstrates the accuracy of the atom counting results. Deviations can be explained by the quantised nature of such small length scale facets and the surface steps which are thought to be critical for catalytic activity.

1 Z Liu et al, Catalysis Review 55 (2013), p255-88
2 H E et al, Ultramicroscopy 133 (2013), p109-119
3 The Absolute Integrator code is free for academic use from www.lewysjones.com/software/
4 J M LeBeau et al, Nano Letters 10 (2010), p4405-8
5 B D Forbes et al, Physical Review B 82, (2010) 104103
6 A R Roosen et al, Computational Materials Science 11 (1998) p16-26
7 B D Todd and R M Lynden-Bell, Surface Science 281 (1993), p191-206


The research leading to these results has received funding from the European Union Seventh Framework Program under Grant Agreement 312483 - ESTEEM2 (Integrated Infrastructure Initiative–I3), and from the EPSRC (grant number EP/K032518/1)

Fig. 1: Simulated library of the scattering cross-section of a pure Pt or Ir crystal with increasing sample thickness. With only 1 atomic number between Pt and Ir they are indistinguishable below 14 atoms. However, the elements have different channelling lengths; this produces a deviation at much higher atom counts.

Fig. 2: Reconstruction of an experimental Pt/Ir particle, left, and equivalent Wulff plot, right. The Wulff plot has been orientated to similar orientation for comparison, (111) faces are purple (with the close packed hexagonal arrangement in the hard sphere model), (100) faces are blue (with square arrangement), and the (110) faces are red.

Type of presentation: Poster

MS-1-P-2226 Mixed FeOx-CeO2-x nanomaterials for chemical looping characterized by transmission electron microscopy and spatially resolved EELS

Turner S.1, Meledina M.1, Galvita V.2, Poelman H.2, Marin G. B.2, Van Tendeloo G.1
1EMAT, University of Antwerp, Groenenborgerlaan 171, 2020 Antwerp, Belgium, 2Laboratory for Chemical Technology, Ghent University, Technologiepark 914, 9052 Ghent, Belgium
stuart.turner@uantwerpen.be

Mixed FeOx-CeO2-x nanomaterials are promising candidates for use as oxygen storage materials in the production of H2 by chemical looping. The technology of chemical looping is based on periodic reduction/re-oxidation cycles of metal oxides, designed to convert hydrocarbons to hydrogen with a quality that exceeds the requirements of all types of fuel cells.1,2 In this work, a series of mixed FeOx-CeO2-x with varying Fe/Ce content are characterized using a combination of advanced imaging techniques and spatially resolved EELS, in order to characterize the presence and nature of the constituting components. The oxide materials are studied throughout the oxidation/reduction cycle, paying special attention the morphology and surface features of the FeOx/CeO2-x material.

Low iron content materials (e.g. 5wt.% FeOx/CeO2-x) typically consist of ceria nanoparticles with sizes ranging from approximately 20 to 60 nm. Electron diffraction and imaging show no evidence for the presence of a separate Fe2O3 (or FeOx) phase in this material. The ceria nanoparticles do show the presence of nanometer-sized voids, which have previously been observed in nanosized ceria. Spatially resolved EELS maps show that both voids and ceria surfaces are decorated with isolated Fe atoms, and that the surface atoms of the ceria nanoparticles and the voids are in a reduced state compared to bulk CeO2.3 Particular attention has been paid to possible changes in the oxidation state and clustering of these Fe species upon oxidation and reduction. The high iron content materials consist of α-Fe2O3 nanoparticles decorated by significantly smaller ceria nanoparticles. In these samples, both structural and valency changes at the FeOx/CeO2-x interface upon cycling have been studied in detail.

1) V. Galvita et al., Topics in Catalysis 2011, 54, 907.

2) V. Galvita et al., Ind. Eng. Chem. Res. 2013, 52, 8416

3) S. Turner et al. Nanoscale, 2011, 3, 3385


S.T. gratefully acknowledges financial support from the Fund for Scientific Research Flanders (FWO). V.G. and H.P acknowledge financial support from the 'Long Term Structural Methusalem Funding by the Flemish Government'.

Fig. 1: (a) Overview HAADF-STEM image of a 5wt.% FeOx–CeO2-x sample. (b) HR-HAADF-STEM showing strong faceting and the presence of voids. (c) Overview HAADF-STEM image and corresponding EELS maps: the Fe is enriched at the ceria surface and within the voids.

Fig. 2: (a) High resolution EELS references for Ce4+ and Ce3+. (b) Overview HAADF-STEM image and (c) Ce3+/Ce4+ map showing surface reduction in the ceria nanoparticles. (d) HAADF-STEM image of the surface of a ceria nanoparticle with (e) corresponding EELS spectra from the surface (black spectrum) and near-surface (blue spectrum) regions.

Type of presentation: Poster

MS-1-P-2252 3D STEM of highly anisotropic insertions in nitride nanorods: a challenge to FIB preparation techniques and transmission electron tomography

Niehle M.1, Trampert A.1
1Paul-Drude-Institut für Festkörperelektronik, Berlin, Germany
niehle@pdi-berlin.de

The ongoing request for innovative semiconductor devices for opto-electronics motivates the growth of low-dimensional objects such as nanowires or nanorods. The realization of designed heterostructures  based on axial or radial symmetry depends on the nanorod's shape, i.e. on its surface facets, and growth conditions. The understanding of the physical properties of the resulting low-dimensional heterostructures necessitates the detailed three-dimensional (3D) microstructure information. Consequently, there is a demand to further establish transmission electron tomography as a feasible tool in materials science – especially for nanoscale semiconductor heterostructures – along with the challenging site specific preparation of adequate samples.
The investigation of inclined GaN nanowires grown on a non-polar (11-22) GaN template with (In,Ga)N insertions at the top by scanning transmission electron microscopy (STEM) tomography is presented in this work. The objects' geometrical arrangement (Fig. 1b) requires a sophisticated sample preparation technique in a dual-beam device comprising a scanning electron microscope (SEM) and a focused ion beam (FIB). On the one hand, the technique allows to isolate the target within an electron transparent lamella (Fig. 1a). On the other hand, the positioning of the lamella realized by the incorporated  micromanipulator and the versatile sample stage enables the chemical sensitive high-angle annular dark field (HAADF) STEM imaging along a <11-20> direction (Fig. 1c) that is not straight forwardly available in conventionally prepared samples. The mounting of the sample with its [0001] orientation along the tilt axis will be discussed.
To access the complex morphology (facets, layer thickness, In content) of (In,Ga)N insertions in the GaN based objects, a HAADF STEM tilt series has been acquired over a tilt range of 165°. The 3D reconstruction reveals the shape (Fig. 2a) and the anisotropic occurrence of (In,Ga)N insertions in layers parallel to the facets of the object (Fig. 2b). The isosurface rendered volume shows that the object is limited by the hexagonal m- and rplanes as well as a rough cap parallel to the c-plane. The r-planes close to the substrate normal are only weakly developed whereas the other four are clearly formed. The cross-sections through the reconstructed 3D volume show high abundance of In in red color whereas the parts dominated by green belong to the GaN core and shell.
This study demonstrates the unique access to complex three-dimensional morphological and chemical information of nanoscale semiconductor heterostructures by HAADF STEM tomography. The requirement of a sophisticated sample preparation technique has to be underlined.


We gratefully acknowledge Enrique Calleja Pardo providing samples for the presented investigations.

Fig. 1: (a) The SEM image represents the target object within the lamella suitable for tomography. (b) The schematic of the target object illustrates its special geometry which challenges TEM sample preparation. (c) The HAADF STEM image exhibits the lamella in cross-section. The white arrow in image (a) and (c) marks the object that is presented in Fig 2.

Fig. 2: (a) Isosurface representation of the three-dimensionally reconstructed object along the direction perpendicular to the substrate and the view onto a (1-100) side facet. (b) The cutaway of a cube from the object (schematic) offers the view onto three ortho-slices parallel to low indexed lattice planes providing chemical information.

Type of presentation: Poster

MS-1-P-2257 Characterisation of air and water stable Cobalt nanorods

Marcelot C.1,2, Lentijo-Mozo S.1, Hungria T.1, Gatel C.2, Fazzini P. F.1, Cormary B.1, Tan R.1, Respaud M.1, Soulantica K.1
1Université de Toulouse; INSA, UPS, CNRS, LPCNO 135 avenue de Rangueil, 31077 Toulouse, France., 2Centre d’Elaboration de Matériaux et d’Etudes Structurales (CNRS), 29, rue Jeanne Marvig, 31055 Toulouse, France
cgarcia@insa-toulouse.fr

The synthesis of hybrid nanoobjects containing a metallic magnetic core and a shell constituted by a noble metal is highly desirable because of their potential use in the fields of electronics, optics, catalysis, biology and medicine. In these nanoparticles, the magnetic core provides the possibility to manipulate the nanoparticle by a magnetic field and the noble metal shell offers protection from oxidation, a surface for functionalization by biomolecules and depending on the metal core additional properties (catalytic, plasmonic etc). In this context, Co anisotropic nanoobjects such as nanorods and nanowires are of special interest for applications in which hard magnetic materials are required. However the development of a continuous shell of a noble metal around Co nanoparticles is a challenge due to incomplete covering by the noble metal. We will describe new hybrid Co-metal core-shell nanorods of different shell composition and thicknesses. The growth of a complete shell is accomplished by introduction of a buffer layer between Co and the noble metal, compatible with the two otherwise immiscible materials. The complete shell protects the Co nanorods from oxidation, as demonstrated by HRTEM (Fig.1) and EDS (Fig.2) analysis and corroborated by the magnetic measurements. These results prove that the magnetic properties of Co, which are very sensitive to oxidation, are stable after exposition of the nanorods to the air for several weeks. Furthermore when the metal shell is thick, it can provide oxidation protection of the Co-core in aqueous solutions for prolonged periods of time. After ligand exchange these nanorods can be transferred from organic solvents into aqueous solutions.


The authors thank the the European Commission for the FP7 NAMDIATREAM project (EU NMP4-LA-2010-246479), the Programme Investissements d'Avenir under the program ANR-11-IDEX-0002-02, reference ANR-10-LABX-0037-NEXT".

Fig. 1: Core-shell nanorod HREM

Fig. 2: Core-shell nanorod STEM-EDS

Type of presentation: Poster

MS-1-P-2261 Chemical Analysis on the Nanometer Scale: Characterization of Copper Nanoparticles by Electron Energy Loss Spectroscopy and Energy Filtered Transmission Electron Microscopy

Schaumberg C. A.1, Wollgarten M.2, Rademann K.1
1Department of Chemistry, Humboldt-Universität zu Berlin, Berlin, Germany, 2Helmholtz-Zentrum Berlin für Materialien und Energie, Berlin, Germany
christian.schaumberg@chemie.hu-berlin.de

A major challenge of modern nanoscience is the need for a detailed knowledge of the chemical composition of novel materials on the nanometer scale. This challenge can be addressed by applying analytical methods to the transmission electron microscopy (TEM). In particular electron energy loss spectroscopy (EELS) opens a wide field of opportunities. Peaks at characteristic core edge energies in EEL spectra of a selected sample area provide information on the presence of certain chemical elements. Beyond that, the fine structure and the chemical shift of the observed core edges provide insights in the composition on an atomic level.[1]
Our work focuses on the characterization of copper nanoparticles generated by pulsed laser ablation of µm-sized powders in organic liquids.[2] The study of different copper precursors points to a reductive step during the synthesis of the copper nanoparticles. In order to investigate the formation mechanism in detail, a profound knowledge of the oxidation state of the copper atoms in the resulting particles is mandatory.
Oxidized copper shows distinct features, so called “white lines”, at the copper L2,3 edge in the EEL spectra. These white lines originate from energy losses through transitions of 2p electrons to empty 3d orbitals. As the 3d orbitals are completely filled for metallic copper the white lines are suppressed. Thus the occurrence of white lines can be used to determine the oxidation state of copper atoms.[3]
With this approach we can show, that the choice of the precursor not only determines the structural by also the chemical properties of the resulting copper nanoparticles. These findings are complemented by elemental maps obtained by energy filtered transmission electron microscopy (EFTEM). The variation of the precursor powder and the investigation of the resulting nanoparticles by EELS and EFTEM leads to a concept for the particle formation mechanism which will be presented.

References

[1] R. F. Egerton, Electron Energy-Loss Spectroscopy in the Electron Microscope, 3rd ed.; Springer: New York, Dordrecht, Heidelberg, London, 2011.
[2] C. A. Schaumberg, M. Wollgarten, K. Rademann, J. Phys. Chem., submitted.
[3] D. Shindo, K. Hiraga, A.-P. Tsai, A. Chiba, J. Electron Microsc., 42, 48-50 (1993).


Fig. 1: TEM images (left) and EEL spectra (right) of nanoparticles generated by laser ablation of CuO powder (top) and Cu3N powder (bottom). The TEM images show the filter entrance aperture used to record the EEL spectra. Thus, the EELS intensity originates solely from the depicted area.

Type of presentation: Poster

MS-1-P-2267 Consequences of gas dynamics on morphology and chemistry during Electron Beam Induced Deposition

Winkler R.1, Fowlkes J.2, Szkudlarek A.3, Melischnig A.1, Utke I.3, Rack P. D.2 4, Plank H.1 5
1Center for Electron Microscopy, Graz, Austria, 2Center for Nanophase Materials Sciences, Oak Ridge, USA, 3Laboratory for Mechanics of Materials and Nanostructures,Thun, Switzerland, 4Department of Materials Science and Engineering, Knoxville, USA, 5Institute for Electron Microscopy and Nanoanalysis, Graz, Austria
robert.winkler@felmi-zfe.at

Focused Electron Beam Induced Deposition (FEBID) is a versatile direct write tool for the fabrication of functional (3D) nanostructures. FEBID uses gaseous precursor which adsorb and diffuse on the surface where they get locally decomposed by a finely focused electron beam. Although many different application concepts have been demonstrates, final applicability depends strongly on predictable morphologies and defined deposit chemistries. Both demands require locally constant precursor coverage which leads to constant ratios between available precursor molecules and potentially dissociation electrons species which is denoted as working regime. What seems to be straightforward turns out to be very complicated when dimensions approaches the nanoscale where local working regimes are influenced by a number of variables like directional gas flux effects, deposit related barriers hindering ideal diffusion or geometrical shadowing effects, comprehensively discussed in this contribution. As starting point it will be demonstrated how patterning directions relates to the directional gas flux, caused by the geometrical arrangement of the gas injection system. It is found that volume growth rates (VGR) can vary by more than 50 % for different patterning orientations which significantly complicates the predictable deposit volumes (Fig. 2). Furthermore, it is shown how the chemistry changes along with the VGR which has strong implications on final functionalities. To demonstrate these effects in a comprehensive way, a new patterning strategy is introduced which visualize morphological and chemical effects within one deposit. Based on these experiments a model is derived which fully explains the observations taking directional adsorption, surface diffusion and local replenishment effects into account as well. The experiments are complemented by finite difference simulations and numeric calculations in well agreement and support the proposed dynamic model of laterally varying working regimes. In order to investigate the tunability of the regime situation, the accessible process parameters during deposition are systematically varied. It is demonstrated how constant working regimes can be established (Fig. 1) which provide both, predictable morphologies and laterally constant chemistries as indispensably required for potential applications. In summary the study demonstrates the nanoscale implications of molecular gas and surface dynamics on final deposit volumes and chemistries. Furthermore, it is also shown how stable conditions can be achieved by a careful setup of the deposition process which is essential for further steps toward industry related FEBID application.


We gratefully thank Prof. Ferdinand Hofer, Roland Schmied, Angelina Orthacker, Martina Dienstleder, Florian Kolb, Barbara Geier and Laura Resch and acknowledge the FFG for financial support.

Fig. 1: 3D AFM height images of FEBID structures fabricated with spiral out patterning strategy and constant electron doses. Unbalanced process parameters (1 ms dwell times) lead to disruption of the morphology (a) in contrast to balanced conditions b) (100 µs dwell times) as effect of indicated directional gas flux component.

Fig. 2: lateral variations of segment heights (red) and chemistry by means of C / Pt ratios (blue) for a disrupted deposit (Fig. 1a). Each patterning point has been patterned only once which demonstrates the strong implications of the directional gas flux effects.

Type of presentation: Poster

MS-1-P-2349 Atomic-scale Defects Leading to Lattice Strain in Single-crystal Ultrafine Gold Nanowires

Kundu P.1,3, Turner S.1, Van Aert S.1, Ravisankar N.2, Van Tendeloo G.1
1Electron Microscopy for Materials Research (EMAT), University of Antwerp, Groenenborgerlaan 171, 2020 Antwerp, Belgium, 2Materials Research Center, Indian Institute of Science, Bangalore 560012, India, 3Institute of Bioelectronics (PGI-8), Forschungszentrum Juelich, D-52425 Juelich, Germany
paro.124@gmail.com

Gold nanowires of molecular scale dimension are of fundamental as well as technological interest owing to their tunable electrical transport characteristics leading to ballistic conduction. This implies single electron sensitivity making them potentially active material for catalysis and molecule sensing. This demands a large scale production of the wires in pristine form for applicability and a detailed atomic structure study to interprete their properties different from the bulk. Although the chemical synthesis route has been reported and electrical transport studies have been carried out recently on the single crystal 2 nm gold wires of large aspect ratio (approx. 500 or more), the structural investigation is not done so far. HRTEM combined with image simulation and exit wave reconstruction can provide information on the local atomic structure, however, with aberration corrected microscopes and advanced analytical methods one can analyse the structure with picometer precision. This method is limited to atomically thin samples. Quantitative HAADF-STEM is a technique to analyse the structure of even few tens of nanometer thick samples and it allows us to determine atom positions in the lattice and determine elemental composition of the atomic columns. Aberration corrected electron microscopy, therefore, combined with advanced quantification methods is a state of the art technique to extract information atom-by-atom 1. Here we present our investigation on these ultrafine gold nanowires to determine their atomic structure by low dose aberration corrected high resolution (S)TEM. Quantification reveals patterned strain in the crystals which increases at the surface layer of atoms and that the wires are faceted with irregular atomic scale surface steps 2. These structural aspects can be related to their unique electrical features and makes them potential candidates for catalysis and sensorics. Besides, from the HRSTEM image, atom counts in the atomic columns in viewing direction is obtained and a 3D visualization of the wire atomic structure could also be deduced. Further, we looked into the atom dynamics due to interaction with the electron beam at higher dose which gives an insight to its mechanical behaviour and stability. Figure 1. provides an overview of the lattice strain and the atom counting analysis.


1 G. V. Tendeloo et al. Adv. Mater. 24, 5655-5675 (2012)
2 P. Kundu et al. ACS Nano 8, 599-606 (2014)


S.V.A and S.T. gratefully acknowledge financial support from the FWO. G.V.T. and P.K. acknowledge the ERC Grant N246791-COUNTATOMS. N.R. acknowledges financial support from the Department of Science and Technology (DST).

Fig. 1: (a) Aberration corrected HRTEM of 2 nm thin wire in [11 ̅0] zone (b) magnified view of the selected wire portion analyzed showing displacement of atomic columns (marked by arrows). (c) High resolution HAADF-STEM image (false color) of the wire analyzed for determining the atom counts in the columns in the [11 ̅0] zone direction as in (d).

Type of presentation: Poster

MS-1-P-2351 A Study on Formation and Thermal Stability of Au-silica Hybrid by Electron Tomography

Kundu P.1,3, Heidari H.1, Bals S.1, Ravishankar N.2, Van Tendeloo G.1
1Electron Microscopy for Materials Research (EMAT), University of Antwerp, Groenenborgerlaan 171, 2020 Antwerp, Belgium, 2Materials Research Center, Indian Institute of Science, Bangalore 560012, India, 3Institute of Bioelectronics, Forschungszentrum Jülich, D – 52425 Jülich, Germany
paro.124@gmail.com

SiO2 based metal nanoparticle hybrids form an important class of material which finds application in catalysis, sensors and biotechnology. Although several protocols exist for synthesizing these hybrids, the complete understanding of the morphology, composition and distribution of the heterounits in three dimensions is lacking. Conventional imaging techniques like SEM and TEM gives information on the size, external/surface morphology and partially on the shape of the nanostructure, but it can be misleading for detailed understanding of the internal distribution and structural composition of the hybrid with a 3D shape. However, these are important factors governing their functionalities like catalytic activity, stability, sensitivity, plasmonic behavior etc. Also, thermal stability of such hybrids are important to be investigated and only few studies are reported on that. Here we present a simple wet chemical route to obtain extremely stable Au nanoparticles (5 – 10 nm) decorated SiO2 spheres without using any external linkers. We investigated the composition by STEM-EDX and 3D ordering of the heterounits of the hybrid using STEM tomography. It reveals presence of Au nanoparticles exclusively on the surface of the SiO2 spheres and not inside the matrix. The same characterization method has been used for understanding the mechanism of formation of the hybrid, intermediate nanostructures resulted in course of reaction. This study reveals that the hybrid formation is mediated by self-assembling of Au nanoparticles due to the presence of oleyl amine and (3-mercaptopropyl) trimethoxysilane (MPTMS) on the Au surface and a formation mechanism of the hybrid is deduced. Thermal stability test is performed and change in morphology is studied by tomography which reveals an excellent stability of the structure up to 400oC beyond which the Au particles starts migrating from the surface of the SiO2 sphere into the matrix. However, no significant coarsening of the particles are observed. These kind of structural changes can have significant impact on physical properties or their functional behavior related to surface activity. This also implies low mobility of the Au particles on the SiO2 surface which is advantageous for several applications 1. The method being general could be used for making similar SiO2 based hybrid to stabilize metal nanoparticles.


1 P. Kundu et al. Angewandte Chemie Int. Ed. DOI: 10.1002/anie.201309288


Funding from the European Community’s Seventh Framework Program ERC grant N°246791 – COUNTATOMS, COLOURATOMS, as well as from the IAP 7/05 Programme initiated by the Belgian Science Policy Office is acknowledged. Funding from Department of Science and Technology (DST) is also acknowledged.

Fig. 1: A schematic description of formation of the Au-SiO2 hybrid, where the Au nanoparticles anchor only to the surface of SiO2, via Au nanoparticle self-assembly; however, on heating to higher temperature (400 oC) the Au particles migrate inside the matrix but not aggregate on the SiO2 surface. This confirms a good thermal stability of the hybrid.

Type of presentation: Poster

MS-1-P-2352 Growth and packaging of ultrathin Au nanowires for enhanced thermal stability: An in-situ TEM study

Kundu S.1, Ravishankar N.1
1Materials Research Centre, Indian Institute of Science, C.V. Raman Avenue, Bangalore 560012, India
subhodex@gmail.com

Ultrathin Au nanowires are potential candidate for catalysis, sensing, plasmonic and biological applications. Most of these applications require a clean interface for better performance. Fragility on polar solvent cleaning and hydrophobicity due to the associated linkers limit the use of the nanowires in their as-synthesized form. We have developed a strategy for growth of these nanowires directly on substrates (Figure 1) that imparts stability to the wires. The study on growth and mechanism of nanowire formation on substrates has been carried out using electron microscopy (SEM & TEM) and other techniques.
Poor thermal stability limits the use of these nanowires to low temperature applications only. Hence, for high temperature applications proper packaging of the nanowires is required. A simple wet-chemical method has been developed to coat these nanowires with mesoporous SiO2 (Figure-2) and TiO2 coatings. The SiO2 layer thickness could be controlled very easily by this method by varying the reaction time. Coating thickness of a few nanometers could be obtained. In-situ TEM thermal stability studies have been carried out on the SiO2 coated nanowires. Figure 3 shows the TEM images of the nanowires as the temperature is increased over a period of 4-5 hours. Bare nanowires had been drop-casted on the same grid for comparison. The non-coated nanowires (marked in red) break into nanoparticles at very low temperature as shown in the set of images. Coated nanowires became segmented at similar temperatures but the segments show remarkable stability at high temperature (5530C) for long times.


NR acknowledges Department of Science and Technology (DST) India for financial support. The electron microscopes are a part of the Advanced Facility for Microscopy and Microanalysis (AFMM) at the Indian Institute of Science.

Fig. 1: TEM image showing ultrathin Au nanowires grown on Carbon support.

Fig. 2: Au nanowires with a thin layer of SiO2 coated to enhance thermal stabilty. Inset shows a thicker coating of SiO2 on the nanowires.

Fig. 3: In-situ TEM heating experiment reveals that SiO2 coated Au nanowires are stable at a temperature of 5530C whereas the drop-casted bare Au nanowires break (marked in red).

Type of presentation: Poster

MS-1-P-2363 Mechanism of Au2Sx/CdS Nanorod Formation by Cation Exchange

Kundu S.1, Kundu P.2, Tendeloo G. V.2, Ravishankar N.1
1Materials Research Centre, Indian Institute of Science, C.V. Raman Avenue, Bangalore 560012, India, 2Electron Microscopy for Materials Research (EMAT), University of Antwerp, Groenenborgerlaan 171,B- 2020 Antwerp, Belgium
subhodex@gmail.com

Cation exchange is the process by which the cation in a compound is replaced by another cation from a suitable precursor. It is very difficult to replace any cation by Au, since the high electron affinity of Au leads to reduction of the precursor to form metallic Au rather than undergoing cation exchange. The competition between cation-exchange and reduction is not well understood. For the case of Au or other elements, one of the processes may be dominant over the other depending on the choice of system and the experimental conditions. Knowing the criterion and having a rational understanding of the process is essential for rational synthesis of heterostructures. In our study, we show that cation exchange is unexpectedly dominant over reduction for the case of CdS-Au.
Bright-field TEM imaging (Figure 1) reveals the presence of small, faceted particles of Au on the CdS nanorods. However, on careful observation it shows the formation of more particles under the electron beam. When the concentration of the Au precursor is low, most of the Au2Sx (x=1 & 3) formed as a result of cation-exchange is on the surface, which on exposure to the electron beam leads to the formation of faceted Au particles. In the case of a higher precursor concentration, the beam effects are highly accentuated as the Au2Sx is present across the depths of the sample which results in shortening of the nanorods; in some cases along with the formation of Au nanoparticles. Energy dispersive X-Ray mapping in STEM mode (Figure 2) clearly depicts the change taking place due to beam irradiation. The HAADF-STEM image in Figure 3a further shows three different regions of contrast. Careful investigation of the high magnification STEM images (Figure 3b) reveals the presence of the cubic Au2S phase which confirms that cation-exchange indeed takes place under the reaction conditions. Thermodynamic calculations have been carried out to understand the experimental observation that paves the way for better predictability of the viable product for various systems under different reaction conditions.


NR acknowledges the Department of Science and Technology (DST) India for financial support. PK and GVT acknowledge the ERC Advanced Grant COUNTATOMS.

Fig. 1: Bright field TEM image showing Au attached to CdS nanorods.

Fig. 2: More of such Au nanoparticles form under the electron beam as is evident from the HAADF-STEM image and the corresponding EDS map.

Fig. 3: (a) At high magnification we observe three different regions of contrast as marked by the blue dotted line. (b) Atomic resolution imaging clearly shows the Au2S and CdS domains.

Type of presentation: Poster

MS-1-P-2377 Recent advances in Catalysis Research using Electron Microscopy

Wagner J. B.1, Deiana D.1, Chorkendorff I.2, Stephens I.2, Hansen T. W.1
1Center for Electron Nanoscopy, Technical University of Denmark, DK-2800 Kgs. Lyngby, Denmark, 2Center for Individual Nanoparticle Functionality, Technical University of Denmark, DK-2800 Kgs. Lyngby, Denmark
jakob.wagner@cen.dtu.dk

Electron microscopy provides a highly versatile platform for the characterization of supported metal nanoparticles for heterogeneous catalysis. With both high spatial resolution as well as spectroscopic capabilities, the EM platform can characterize materials in detail. Recent developments include high solid angle EDX detectors, which can rapidly acquire high-resolution elemental maps, and micro electro-mechanical systems (MEMS) based heating holders that can heat samples at very high rates with only little spatial drift. With the addition of environmental capabilities, the microscope can even probe samples under reactive environments.
It is impractical to use all techniques and modification on a single instrument. Hence, in order to obtain the complete picture of catalyst samples, several platforms can be employed.
A recent trend in catalysis is the use of materials that have been engineered at an atomic level. In particular, Density Functional Theory (DFT) can be used to computationally screen for new materials. These are often multi-metal alloys, which add new functionality and can reduce the amount of precious metals. Such samples can be size selectively produced either by physical routes, e.g. time of flight mass selection or chemical synthesis e.g. micelle encapsulation. Whereas these approaches may not be technically applicable for large-scale synthesis, they provide a valuable route for gaining fundamental knowledge.
Here, we show findings from three different systems used in three different reactions. Namely Pt-Y for oxygen electroreduction to H2O, Pd-Hg for electrochemical synthesis of hydrogen peroxide and ruthenium based catalyst used for methanation [1-3]. With these examples, we illustrate two principle points of nanoparticle functionality: composition and shape.
In the case of the bimetallic catalysts, the elemental distribution in the nanoparticles is of fundamental interest: Do they form a core-shell system or do form an evenly distributed mixture/alloy? Using a high solid angle EDX detector, elemental maps can be efficiently collected and the elemental distribution monitored. Such verification is essential to understand the working principle of the catalyst.
Ruthenium nanoclusters can be used for methanation of carbon monoxide, a reaction used to clean up feed gas for e.g. proton exchange fuel cells (PEM). As-synthesized, the Ru particles assumed high surface-area raspberry-like shapes. However, after treatment under conditions relevant for the methanation reaction, the particles adopted more spherical shapes.
[1] F. Masini et al. J. Catal. 308 (2013) 282
[2] S. Siahrostami et al. Nature Materials 12 (2013) 1137
[3] A. Verdaguer-Casadevall et al. Nano Letters 13, dx.doi.org/10.1021/nl500037x


Fig. 1: a) STEM micrograph of a nanoparticle and b-d) corresponding Y, Pt and combined X-Ray elemental maps.

Fig. 2: a) STEM micrograph of a nanoparticle and b-d) corresponding Hg, Pd and combined X-Ray elemental maps.

Fig. 3: Ruthenium nanoparticle imaged under different conditions relevant for the methanation reaction. a) Room temperature, vacuum; b) 427°C, vacuum; c) 427°C, 230 Pa 1:10 CO/H2.

Type of presentation: Poster

MS-1-P-2392 Global-refinement electron exit wave reconstruction from focal series of ceria nanoparticles

Borisenko K. B.1, Young N. P.1, Kirkland A. I.1
1Department of Materials, University of Oxford, Parks Road, Oxford, OX1 3PH, UK
konstantin.borisenko@materials.ox.ac.uk

Nanoparticles play an increasingly important role in catalysis. At the nanoscale, stoichiometry and the thermodynamic stability of different crystallographic facets may be modified as compared to the bulk, which is one of the reasons for their increased importance. Understanding activity and selectivity of nanoparticle catalysts requires detailed understanding of catalytic reactions at the atomic scale. An important step towards this goal is to obtain accurate atomic structures of catalytic nanoparticles and especially of their surface and immediate subsurface regions.

Aberration-corrected high resolution transmission electron microscopy (HRTEM) is a well suited tool for studying atomic structures of such particles. Quantitative analysis of the electron exit wave obtained from the experimental series of images with different focus can in principle provide some additional information on the three-dimensional structure.

In the present work we test an enhanced approach to quantitative exit wave restoration from the focal series of HRTEM images obtained for ceria nanoparticles. The existing linear exit wave restoration codes are based on the assumption that the sample under investigation is a weak-phase object. This approximation applied to a general object can result in an incorrect restoration. A more general approach is to reconstruct the exit wave by minimising the sum of squared differences between the simulated and experimental images where both the amplitude and the phase of the exit wave can be restored accurately. We suggest using the exit wave reconstructed by the linear approach as an initial approximation for this more general reconstruction. The successful restoration using the suggested method is dependent on knowing accurate aberration parameters of the microscope, especially the focus range and the focal step, and accurate image alignment in the experimental focal series. These are not routinely available from the actual experimental conditions. The suggested approach employs refinement of the exit wave together with both the aberration parameters and the image alignment in a single refinement cycle. The resulting exit wave is compared with exit wave obtained by theoretical multislice simulations and with the exit wave obtained by linear restoration software. We also investigate the origin of the increased contrast at the edges of the nanoparticles seen in the reconstructed phase, examining whether it is a consequence of adsorbed light species or a manifestation of electrostatic surface potential.


Financial support from the European Union under the Seventh Framework Program under a contract for an Integrated Infrastructure Initiative (Ref 312483-ESTEEM2) is gratefully acknowledged.

Fig. 1: Global-refinement exit wave restoration algorithm implemented in the present study.

Fig. 2: Reconstructed electron exit wave amplitude a) and phase b) of ceria nanoparticles. Note the bright contrast at the edges of the nanoparticles in the phase image.

Type of presentation: Poster

MS-1-P-2497 TEM and STEM observations of Au/Fe2O3 catalysts

Akita T.1, Maeda Y.1, Kohyama M.1
1National Institute of Advanced Industrial Science and Technology (AIST)
t-akita@aist.go.jp

Gold exhibits characteristic catalytic properties when Au nano-particles are supported on the metal oxides [1,2]. It has been reported that catalytic properties depending on the kind of the metal oxide supports are observed for various catalytic reactions. For example, high catalytic activity is observed in the low temperature CO oxidation when TiO2 was used as support and high catalytic activity for water-gas-shift reaction at low temperature was observed when the CeO2 is used as support [3]. The origin of the catalytic properties of Au catalysts is not clear yet although it is suggested that the interface between the small Au particle and the metal oxide support act as active sites [4,5]. In order to clarify the relation between the fine structure and the catalytic properties at the Au-metal oxide interface, we have carried out the structure analyses on the small Au particle supported on TiO2, NiO and CeO2 with a transmission electron microscopy (TEM) and annular dark field scanning transmission electron microscopy (ADF-STEM) [6,7]. In this experiment, the basic structure of Au nano-particles supported on Fe2O3 was observed in atomic scale by HRTEM and STEM.

Au/Fe2O3 catalysts were prepared by solid grinding method using organogold complex [8] and deposition precipitation (DP) method. γ-Fe2O3 fine particle (Nanophase Tech. Corp.) which has spinel structure was used for support. The catalysts were calcined at 573K for 4 hours in air. The observations were carried out by using aberration corrected TEM/STEM (FEI Titan3 G2 60-300). Accelerating voltage for the observation was 300kV.

Figure 1 shows typical ADF-STEM images of Au/γ-Fe2O3 catalyst prepared by DP method. Small Au particles of approximately 2-10 nm in diameter are deposited on theγ-Fe2O3 support. Theγ-Fe2O3 support crystal exhibit polyhedral shape with low index facets such as {111}, {100}. Figure 2 shows profile-view HRTEM images of Au particles onγ-Fe2O3 (111). The incident electron beam direction was adjusted alongγ-Fe2O3 [1-10] zone axis. Gold particles tend to be deposited on theγ-Fe2O3 surface with the preferential orientation relationships of (111)[1-10]Au//(111)[1-10]γ-Fe2O3 or (111)[-110]Au // (111)[1-10]γ-Fe2O3 for theγ-Fe2O3 (111) surface. The high resolution STEM observations were also carried out for the Au/γ-Fe2O3 interface.

References

[1] M. Haruta et al., Chem. Lett., (1987) 405.

[2] M. Haruta, Catal. Today 36(1997)153.

[3] H. Sakurai et al., Appl. Catal. A: General 291 (2005)179.

[4] T. Fujitani et al., Angew. Chem. Int. Ed. 48(2009) 9515.

[5] T. Fujitani et al., Angew. Chem. Int. Ed. 50(2011)10144.

[6] T. Akita et al., Surf. Interface Anal.40, (2008)1760.

[7] T. Akita et al., J Mater Sci. 43(2008)3917.

[8] T.Ishida et al., Chem. Eur. J. 14 (2008) 8456.


The authors are grateful to Ms. F. Arai, Ms. C. Fukada and Ms. M. Makino for their assistance with sample preparation.

Fig. 1: FIG. 1. Typical ADF-STEM image of Au/γ-Fe2O3 catalyst.

Fig. 2: FIG. 2. HRTEM image of Au onγ-Fe2O3 substrate.

Type of presentation: Poster

MS-1-P-2533 Interconnection of Nanoparticles within 2D Superlattices of PbS/ Oleic Acid Thin Films

Simon P.1, Bahrig L.2, Baburin I. A.2, Formanek P.3, Röder F.4, Sickmann J.4, Lichte H.4, Hickey S. G.2, Eychmüller A.2, Kniep R.1, Rosseeva E.5
1] Max Planck Institute for Chemical Physics of Solids, Nöthnitzer Straße 40, 01187, Dresden,Germany, 2TU Dresden, Physical Chemistry, Bergstrasse 66b, D-01062 Dresden, Germany, 3Leibniz-Institut für Polymerforschung Dresden e.V., Hohe Straße 6, 01069 Dresden, Germany, 4Institute of Structure Physics, Triebenberg Laboratory for High-Resolution Electron Microscopy and Holography, Technical University of Dresden, Zum Triebenberg 50, 01328 Dresden Zaschendorf, Germany, 5University of Konstanz, Physical Chemistry, POB 714, D-78457 Konstanz, Germany
Paul.Simon@cpfs.mpg.de

Ensembles of nanoparticles possess collective properties that are dissimilar to those demonstrated by the individual particles and self-assembly has emerged as a powerful means by which the structure and properties of inorganic nanoparticle arrays can be manipulated.. In order to aid in the resolution of the keenly contested debate between proponents of the fibrillation model and those of the electrostatic forces interaction model the structure formed by monolayers of PbS colloidal nanocrystals was investigated using high-resolution spherical aberration corrected TEM, high-resolution electron holography and energy filtered TEM [1,2]. By employing this suite of techniques it could be observed that the truncated octahedrally shaped nanoparticles form 2D close-packed layers interconnected by organic fibrils of oleic acid which are partially mineralised by PbS. These bridges, whose diameters are between 0.3 and 2 nm, keep the face to face orientation of the nanoparticles fixed, thus preventing them from assuming an arbitrary orientation. The complex and textured structure of the monolayer assembly is caused by the habit of the truncated octahedral PbS nanoparticles bearing angles close to ideal values of 54° and 71° between their {100} and {111} faces. By means of electron holography, approximately 10-15 fibrillar interconnections between neighbouring particles in the as-prepared films have been observed. Each nanoparticle is surrounded by six other individuals. At least two or three organic “linkages” are formed between the particles and connect to a nearest neighbour. Most of the organic connections can be mineralised successively by PbS during careful annealing. By using this bottom-up technique access to length scales of sub-nanometer dimensions, presently not accessible to top-down techniques can be attained. This type of isolated but yet interconnected structure formed by the inorganic bridges, represents an ideal “isolated but connected” structure that preserves the effects of quantum confinement present within the individual nanoparticles whilst at the same time having the potential to provide high electron mobility throughout the extended structure.

[1] P. Simon, E. Rosseeva, I.A. Baburin, L. Liebscher, S.G. Hickey, R. Cardoso-Gil, A. Eychmüller, R. Kniep, W. Carrillo-Cabrera, Angew. Chem. Int. Ed. 2012, 51, 10776-10781.

[2] P. Simon, L. Bahrig, I.A. Baburin, P. Formanek, F. Röder, J. Sickmann, S.G. Hickey, A. Eychmüller, H. Lichte, R. Kniep, E. Rosseeva, Adv. Mater. 2014 DOI: 10.1002/adma.201305667


Fig. 1: 3D representation of the phase image retrieved from the electron hologram. Color code corresponds to 4 nm height from green to blue. The bridging organic fibrils appear yellow.

Fig. 2: 2D representation of phase image. The PbS nanoparticles and the interconnecting sub-nanometer oleic acid fibrils appear bright in the phase image.

Fig. 3: (a) Cs-corrected HR-TEM image of two nanoparticles interconnected by a PbS bridge. The PbS bridge (red arrow) has a diameter of 0.3 nm and a length of 1.5 nm. The periodicity along the bridge corresponds to 0.3 nm which is equivalent to the (200) lattice plane of PbS. (b) Digitally zoomed area.

Fig. 4: Idealized model of isolated but interconnected PbS nanoparticles.

Type of presentation: Poster

MS-1-P-2556 Atomic Resolution Characterization and Dynamics due to Beam Interaction of Ni base Nanoparticles for Energy Devices

Calderon H. A.1, Godinez-Salomon F.2, Solorza-Feria O.2, Specht P.4, Kisielowski C.3
1Dept. Física, ESFM-IPN, Zacatenco D.F. 07338, Mexico, 2Dept. Química, CINVESTAV, Mexico D.F., Mexico, 3JCAP and NCEM, LBNL, Berkeley, CA 94720, U.S.A., 4Dept. Mats. Sci. Eng., UCB, Berkeley, CA 94720
hcalder@esfm.ipn.mx

Ni base nanoparticles (NPs) are characterized under low dose conditions in TEM mode. These nanoparticles are mainly designed to act as catalysts in energy devices. Ni, NiO and Pt@NiO nanoparticles are investigated. Particularly the use of Nio@Pt for solar cells (artificial photosynthesis) is attractive, namely the hydrogen evolution center, while NiO has been tested as a catalyst for the oxygen evolution center. Consequently an atomic characterization of the involved nanocrystals is of particular importance. Here, transmission electron microscopy is used with the objective to determine nature, shape and atomic distribution of Pt for different loadings (0-16 at.%) on a Ni core basis. In all cases the electron dose rate has been kept in the range 20-150 e-2s in order to avoid surface rearrangement by interaction with the electron beam. The TEAM 05 (80 KeV) has been used together with focal series reconstruction (EWR) to recover both phase and amplitude images that provide information of the spacing and the chemical nature of the corresponding atomic columns. Two procedures have been used for synthesis of nanoparticles. One of them produces Ni and the other NiO-NPs. NiO NPs are then covered with different loadings of Pt in order to create incomplete core shell structures but with superior catalytic activity. Figure 1 shows phase images of Ni NPs, their size varies from 1 to 7 nm and can agglomerate most likely due to their magnetic characteristics. The dose rate used to acquire the experimental images is 30 e-2s. Figure 1b shows experimental images of NiO NPs acquired with a dose rate of 120 e-2s, their average size is around 1.5 nm. During processing Pt is deposited on NiO particles and a typical example is given in the phase images shown in Figs. 2a-b, the dose rate is around 55 e-/Å2s and the Pt coverage is nominally 8 at. %. The nanoparticles have mostly irregular shapes. There is a negligible particle transformation due to the weak interaction with the electron beam. These NPs are nevertheless susceptible to alteration in shape and structure as a consequence of electron beam sample interaction. An example is given in the phase images shown Figs. 3 a-c. In these cases, the dose rate has been increased from 55 e-2s (Fig. 3a) to 300 e-2S (Fig 3b) and 1400 e-2s (Fig. 3c). The particle under observation initially losses atoms that apparently redeposit on the carbon support and migrate (partially) to form a new crystal. The selected NP becomes bicrystalline at the end of this experiment that clearly shows the need to use a proper electron dosage for observation and the possible large influence of thermal effects. Phase images have been used for simulation in order to determine the Pt coverage.


CONACYT (FOINST. 75/2012, 129207 and 148304) and IPN (COFAA-SIP) are gratefully acknowledged for financial support.

Fig. 1: Figure 1. (a) Phase image of Ni nanoparticles and (b) Experimental image of NiO nanoparticles at a dose rate of 150 e-/A2s.

Fig. 2: Fig. 2. Phase images of NiO nanoparticles with a Pt coverage of 8 at.%. and taken with a dose rate of 55 e-/Å2s.

Fig. 3: Fig. 3. Phase images of NiO nanoparticle with an 8 at. % Pt coverage. (a) Dose rate of 55 e-2s. (b) Dose rate of 300 e-2s and (e) Dose rate of 1400 e-2s.

Type of presentation: Poster

MS-1-P-2562 Localization microscopy (SPDM) facilitates high precision control of lithographically produced nanostructures

Grab A. L.1, Hagmann M.2, Dahint R.1, Cremer C.2,3
1Angewandte physikalische Chemie, Im Neuenheimer Feld 253, 69120 Heidelberg, Germany, 2Kirchhoff-Institute for Physics, Im Neuenheimer Feld 227, 69120 Heidelberg, Germany , 3Institute of Molecular Biology, Ackermannweg 4, 55128 Mainz, Germany
martin.hagmann@kip.uni-heidelberg.de

In numerous fields, the development of innovative technologies requires a refinement and miniaturization of existing systems resulting in an increasing requirement for process friendly quality and dimension control for industry. Localization microscopy (SPDM) provides a precise control of nanostructures, which are indispensable for example for optronics, biosensing applications, manufacturing of electrical elements, biomedical applications, environmental issues, flow profiles in air or water and self-cleaning surfaces.
The principle of this "superresolution microscopy" technique is the use of "point-like" objects carrying different spectral signatures (e.g. fluorescent dyes different in absorption and/or emission spectra; fluorescent dyes with different life-times; time dependence of luminescence, reversible bleaching behaviour, etc.).
Using a lithographic approach, highly regular nanostructures have been generated and marked with Alexa 647 dyes. The spatial organization of the dyes on nanostructured surfaces consisting of interconnected cubes has been averagely localized down to 6 nm using localization microscopy. Herewith we illustrate two aspects: The application potential of localization microscopy as an integrated process for quality control in addition to the absolute spatial calibration of Spectral Precision Distance Microscopy (SPDM).
The findings will be important in the field of product control for industrial applications and long-term fluorescence imaging and calibration for most super-resolution fluorescence microscopes in general. As SPDM improves the optical resolution compared to standard fluorescence microscopy, structure dimensions and the excellent quality of the lithographical grating were resolved beyond the Abbe limit with high precision.


The authors gratefully acknowledge the sample fabrication by the Karlsruhe Nano Micro Facility, especially we like to express our gratitude to A. Nesterov-Müller. All authors thank S. Dithmar for financial support and our dear colleagues Gerrit Best, Sabrina Rossberger, Dr. Udo Birk, Florian Schock and Dr. Fanny Liu. We thank the Boehringer Ingelheim Foundation for generous support.

Type of presentation: Poster

MS-1-P-2567 Mechanical Behavior at the Nanoscale: the benefits of coupling of In situ TEM nano-compression and compression inside a Diamond Anvil Cell

ISSA I.1,2, Calvié E.1, Joly-Pottuz L.1, Rethore J.2, Amodeo J.1, Esnouf C.1, Chevalier J.1, Garnier V.1, Masenelli-Varlot K.1
11Université de Lyon, INSA-Lyon, CNRS, MATEIS, 69621 Villeurbanne, France, 22Université de Lyon, INSA-Lyon, CNRS, LaMCoS, 69621 Villeurbanne, France
inas.issa@insa-lyon.fr

Nanometer sized objects are attracting large attention nowadays due to their breakthrough mechanical properties such as high hardness, crack propagation resistance and high elastic limit in comparison to the bulk of their counterparts [1]. In situ TEM nanoindentation is a particularly well suited technique for the mechanical testing of nano-sized objects. Images of the deforming sample and force-displacement curves can simultaneously be acquired. The challenge remains in the identification of the material mechanical behavior, namely the constitutive law with the intrinsic parameters – Young modulus, yield strength, Poisson ratio – as well as the understanding of the deformation mechanism.In this study, we propose an innovative method for a complete mechanical analysis of nanoparticles in the size range [30 nm-300 nm]. This protocol consists in coupling of in situ TEM nano-compression tests of isolated nanoparticles, image analysis and mechanical simulations. After the experiments, the load–real displacements curves are measured by Digital Image Correlation. Then a constitutive law is obtained through an inverse Finite Elements simulation. The determination of a constitutive law includes the determination of the material intrinsic parameters such as Young modulus, Yield stress, hardening coefficient, and stress at fracture.In this presentation, the method will be presented through the analysis of transition alumina nanoparticles. It will be shown that such ceramic nanoparticles can undergo large plastic deformation, which is not observed in the bulk (Fig1). The parameters of the constitutive law will be discussed in the light of the literature, and especially the work from K. Zeng et al. [1]. They showed that the electron beam, during in situ TEM nano-compression tests of silica nanoparticles, creates structural and bonding defects throughout the entire sample and facilitates the plasticity of the nanoparticles.The deformation mechanisms will be investigated through performed compression experiments in a Diamond Anvil Cell, at room temperature and in the absence of electron beam. We will present the results obtained from HRTEM observations of thin foils extracted from samples compacted at various uniaxial pressures. We will show that plastic deformation occurs also in this case (Fig2). Moreover, the appearance of a nanoparticle preferential orientation will be evidenced (Fig3). This point will be discussed in function of the possible slip systems. Finally, we will demonstrate that such HRTEM analysis gives interesting pieces of information, which permit to better understand how the nanoparticles behave and deform during the in situ experiments.
[1] Kraft et al. Annual Review of Materials Research,2010.
[2] K. Zheng et al. Nature Communications, 2010.


S. Le Floch, D. Machon . Institute ILM, Université Lyon 1 (compaction nano-powder DAC)

Fig. 1: TEM in situ nano-compression force-displacement curve. The simulations using DIC-FE (red) or the analytical method (blue). A good agreement is found for both simulations with the experiment, especially for the DIC-FE method which takes into account the plastic regime, contrary to the analytical method which is valid only in the elastic domain.

Fig. 2: (Left) TEM image revealing the plastic deformation of transition alumina nanoparticle compacted in a DAC at 5 GPa uniaxial pressure. (Right) TEM image of compacted transition alumina in DAC at 20 GPa. It reveals the oriented crystallographic texture with respect to the compression axis.

Fig. 3: (a) HRTEM image of the FIB thin foil of a zone of contact between two alumina nanoparticles. (b) Fourier Transform of the dotted zone of the deformed particle.

Type of presentation: Poster

MS-1-P-2568 In situ TEM Nano-Compression and Mechanical Analysis of MgO

Issa I.1,2, Amodeo J.1, Joly-Pottuz L.1, Réthoré J.2, Esnouf C.1, Garnier J.1, Morthomas J.1, Masenelli-Varlot K.1
1Université de Lyon, INSA-Lyon, CNRS, MATEIS, 69621 Villeurbanne, France, 2Université de Lyon, INSA-Lyon, CNRS, LaMCoS, 69621 Villeurbanne, France
inas.issa@insa-lyon.fr

Nanometer-sized objects are attracting large attention nowadays due to their breakthrough mechanical properties such as high hardness, crack propagation resistance and high elastic limit in comparison of the bulk state of the studied material [1].
Moreover, these nano-objects exhibit large plastic deformation under high load; this was not expected for certain materials and especially for ceramics. Large numbers of studies nowadays are dedicated to plastic deformation of Metals at the nano-scale, and few are reported on ceramics [2, 3].
The origin of this plastic deformation is still not very well defined. Mechanisms proposed are size dependent, and link this behavior to dislocations nucleation at surfaces and slipping on certain planes depending on the crystal orientation with respect to the solicitation direction. Another mechanism proposed is the mechanical twining via full dislocations dissociations into partial Shockley dislocations that glide on a slipping plane (the denser) of the crystal.

A protocol consisting of in situ TEM nano-compression tests of isolated nanoparticles coupled with data processing by Finite Elements and Molecular Dynamics simulations has been developed [2], and applied to the study of spherical alumina nanoparticles. Identification of deformation mechanisms remains quite difficult since the orientation of the nanoparticle on the substrate prior to compression is not controlled.
In this study, we will present in situ TEM nano-compression experiments on MgO nanocubes. The main advantage of studying such nanocubes lies in the fact that their crystallographic orientation with respect to the indenter tip is fully known. It will be shown that MgO can undergo large plastic deformation, more than 50%, without any fracture. Then, we will propose a mechanical behavior law from the analysis of the images and curves followed by Finite Elements simulation. Finally, deformation mechanisms will be identified from the comparison between the contrasts in the images and Molecular Dynamics simulations [4].

[1] Kraft et al. Annual Review of Materials Research (2010) 40:293-317
[2] Calvie et al. Journal of the European Ceramic Society (2012) 32:2067-71
[3] Korte et al. Acta Materialia (2011) 59:7241-54
[4] Amodeo et al. Acta Materialia (2011) 59:2291-2301


Fig. 1: 100 nm edge size, MgO Nanocube before compression in situ in TEM

Fig. 2: 100 nm edge size, MgO Nanocube After compression in situ in TEM

Fig. 3: Stress-strain curve Obtained from Load-Real displacements curve of the nanocube compressed in situ

Type of presentation: Poster

MS-1-P-2577 Monodisperse embedded nanoparticles derived from an atomic metal-dispersed precursor of layered double hydroxide for architectured carbon nanotube formation

Tian G.1, Zhao M.1, Zhang B.2, Zhang Q.1, Zhang W.3, 4, Huang J.1, Chen T.1, Qian W.1, Su D.2, 3, Wei F.1
1Beijing Key Laboratory of Green Chemical Reaction Engineering and Technology, Department of Chemical Engineering, Tsinghua University, Beijing, China, 2Shenyang National Laboratory for Materials Science, Institute of Metal Research, Chinese Academy of Sciences, Shenyang, China, 3Department of Inorganic Chemistry, Fritz Haber Institute of the Max Planck Society, Berlin, Germany, 4Department of Energy Conversion and Storage, Technical University of Denmark, Roskilde, Denmark
tian-gl10@mails.tsinghua.edu.cn

Monodisperse metal nanoparticles (NPs) with high activity and selectivity are among the most important catalytic materials. However, the intrinsic process to obtain well-dispersed metal NPs with tunable high density (ranging from 1013 to 1016 m-2) and thermal stability is not yet well understood. Herein, the preparation of metal NPs with tunable areal density from layered double hydroxide (LDH) precursors in which the metal cations were pre-dispersed at an atomic scale was explored. Large quantities of mesopores induced by the Kirkendall effect were formed on the as-calcined layered double oxide (LDO) flakes. The O atoms bonded with Fe3+ cations were easily to be extracted at a temperature higher than 750 oC, which greatly increased the mobility of Fe. Consequently, coalescence of the reduced Fe atoms into large NPs enhanced the Kirkendall effect, leading to the formation of monodisperse embedded Fe NPs on the porous LDO flakes. The flake morphology of LDHs was well preserved, and the areal density of Fe NPs on the LDO flakes can be well controlled through adjusting the Fe content in the LDH precursor. With higher Fe loading, larger Fe NPs with higher areal density were available. When the areal density was increased from 0.039 to 0.55, and to 2.1 × 1015 m-2, the Fe NPs embedded on the LDO flakes exhibited good catalytic performance for the growth of entangled carbon nanotubes (CNTs), aligned CNTs, and double helical CNTs, respectively. This work provides not only new insights on the chemical evolution of monodisperse NPs from an atomic metal-dispersed precursor, but also a general route to obtain tunable NPs as heterogeneous catalysts for chemical and material production.

References:
1. G. L. Tian, M. Q. Zhao, B. S. Zhang, Q. Zhang, W. Zhang, J. Q. Huang, T. C. Chen, W. Z. Qian, D. S. Su and F. Wei, J. Mater. Chem. A, 2014, 2, 1686–1696


The work was supported by the Foundation for the China National Program (No. 2011CB932602) and Natural Scientific Foundation of China (No. 21306102). Bingsen Zhang is supported by the IMR SYNL-T.S. Keˆ Research Fellowship. Bingsen Zhang thanks the financial support provided by the China Postdoctoral Science Foundation (2012M520652).

Fig. 1: STEM image of Fe distributed on LDO flakes.

Fig. 2: (a) Entangled CNTs grown on LDH-I, (b) aligned CNTs grown on LDH-III, and (c) double helical aligned CNTs grown on LDH-V. (d) The phase diagram of CNTs grown on flat/flake substrates with different catalyst densities and sizes.[1]

Type of presentation: Poster

MS-1-P-2623 Effect of the amount of dopant and the synthesis method on structure and morphology of nanocrystalline Ce1−xRExO2−y

Mendiuk O.1, Kepinski L.1
1Institute of Low Temperature and Structure Research, PAS, Wrocław, Poland
o.mendiuk@int.pan.wroc.pl

Nanocrystalline pure or doped ceria is an important material widely used in various fields of technology, including optics, microelectronics and catalysis. Doping of ceria with transition metal ions enhances its property and improves the thermal stability of nanocrystalline ceria against sintering. It has been established that catalytic activity of ceria nanoparticles depends strongly on their morphology: nanoparticles with cube or rod morphology, exposing {1 0 0} planes at the surface, are desirable for catalytic reactions of CO and soot combustion [1,2].
In this work mixed Ce1−xLnxO2−y (Ln=Gd, Er) oxides were synthesized by the hydrothermal treatment [3,4]. Two modifications of the hydrothermal treatment – classical and microwave assisted – were applied. The effect of the amount of dopant and the synthesis method on the phase composition and morphology of the of lanthanide oxides was studied by SEM-EDS, EBSD, TEM, XRD and Raman spectroscopy.
By classical hydrothermal treatment, for low doping level, nanocubes of the mixed Ce-Ln oxide with fluorite structure and bimodal size distribution (small 5-20 nm and much bigger 50-80 nm) were formed (Fig.1), while at higher doping (x > 0.3 ) rod-like particles of Ln hydroxide were also observed. Using of microwave radiation enabled the synthesis of the nanocubes of the mixed oxides at significantly shorter time, but the resulting materials is different: over broad range of Ln contents (0.05 <x< 0.5) particles with nanorod and nanocube morphology were obtained (Fig.2). TEM show that smallest particles with low doping level, which could not be characterized by SEM, contains mostly regular cube shape particles, though there is a fraction of small particles having rounded corners (Fig.3). SAED pattern contain sharp rings that can be assigned to fluorite structure of ceria. Particle size distribution is very broad and bimodal.
EBSD combined with EDS was used to analyze the structure and composition of unusual, large oxide nanocubes (50 – 80 nm) appearing in the samples (Fig.4). It appeared that the nanocubes of the mixed Ce-Ln oxide have fluorite type structure of CeO2 and are single crystals but not aggregates of smaller crystallites.
[1] X.W. Liu, et.al., J. Am. Chem. Soc. 131 (2009) 3140–3141;
[2] K.B. Zhou, X. Wang, X.M. Sun, Q. Peng, Y.D. Li, J. Catal. 229 (2005) 206–212;
[3] H.X. Mai, L.D. Sun, Y.W. Zhang, R. Si, W. Feng, H.P. Zhang, H.C. Liu, C.H. Yan, J. Phys. Chem. B 109 (2005) 24380–24385.;
[4] Z. Wang, Q. Wang, Y. Liao, G. Shen, X. Gong, N. Han, H. Liu, Y. Chen, ChemPhysChem., 12 (2011) 2763–2770


The authors thank Mrs. E. Bukowska for XRD measurements and Mr. M. Ptak for recording Raman spectra

Fig. 1: SEM image from Ce0.95Er0.05O2−y (classical hydrothermal treatment)

Fig. 2: SEM image from Ce0.95Er0.05O2−y (microwave assisted hydrothermal treatment)

Fig. 3: TEM image and SAED pattern from Ce0.95Er0.05O2−y (classical hydrothermal treatment)

Fig. 4: EBSD indexed pattern from single nanocube of Ce0.95Er0.05O2−y

Type of presentation: Poster

MS-1-P-2625 Fractal growth of porous Cu2S nano-crystals

Zhu G. Q.1, Wang C. H.1, Shi L.1, Lu W.1, Zhang J. P.1
1Suzhou Institute of Nano-Tech and Nano-Bionics, Chinese Academy of Sciences, China 215125
jpzhang2008@sinano.ac.cn

In this work we report on the structure characteristics of Cu2S nano-crystals and the fractal features in crystal growth.

The Cu2S crystals were produced by using a carbon-coated TEM copper grid added with a few drops of dispersed Graphene loaded with sulfur nano-particles in ethanol. When the solution is dried, a variety of well-constructed tine crystals were observed near the copper bars in microscope, which were mostly dendrites as shown in Figure 1a, similar to the observation by Q Han[1]. The diffraction patterns obtained from different dendrites all showed a 6-fold symmetry, as shown in Fig.1b, which can be indexed with hexagonal Cu2S[2]. No other phases or amorphous copper sulfide, as reported in [1], were observed.

Actually the crystallized dendrites exhibit a porous feature since they are composed of numerous nano-crystals in size of few nanometers, as indicated in dark-field scanning TEM images, an example presented in Figure2a. An interesting question is the diffraction pattern from a large area of a dendrite having a number of branches, see in Fig.1a, show a simple [001] pattern that means all nano-crystals are so well oriented that not only along the C-axis, but also the atomic arrangement of those c-planes are aligned in 3o of rotation with respect to each other, see the Fig.1b. That implies hundreds or even thousands of copper sulfide nano-particles bonded together porously could behave as a single crystal, rather than a randomly arranged one.

Of the well-oriented nano-crystals the produced copper-sulfide crystals present interesting self-similarity morphologies, the dendrites like a leaf, a fern, or even a mountain top, some of them We calculated the fractal dimensions using the box-counting method [3] on the nano-crystals and the Matlab codes designed by San Pedro [4]. The fractal dimension of a typicl Cu2S-dentride, as shown in Figure 3a is 1.8623, while the ideal value is 2.

[References]

[1] Qiaofeng Han, Shanshan Sun, Jiansheng Li and XinWang, Nanotechnology 22 (2011) 155607 .

[2] Cava. R.J., Reidinger. F., Wuensch. B.J., Solid State Ionics, 5 (1981) 501.

[3] G. Hartvigsen, The Analysis of Leaf Shape Using Fractal Geometry. The American Biology Teacher. 62 (2000) 664.

[4] S. San Pedro, Fractal Dimensions of Leaf Shapes, Math 614-Sp2009 Web site: http://www.math.tamu.edu/~mpilant/math614/StudentFinalProjects/SanPedro_Final.pdf


[Acknowledgement]

This project is supported by National Basic Research Program of China (2010CB934700) and National Natural Science Foundation of China (Grant No. 21210004).

Fig. 1: Figure 1. (a) Dendrites observed on Cu-bars of thin carbon film coated copper grids; (b) the corresponding electron diffraction pattern from the selected area circled in (a) indicating a [001] oriented Cu2S and the C-plane of different branches are well-aligned within 3o in rotation respectively.

Fig. 2: Figure 2. An enlarged image from a portion of a dendrite obtained in dark field scanning TEM mode (Z-contrast imaging) exhibits a porous structure composed of nano- Cu2S-crystals (bright dots) in size of 5±1 nm and different size of holes (black dots).

Fig. 3: Figure 3. Cu2 Nano-crystal growth produces a variety of fragmented self-similar shapes similar to leafs, ferns, or other plants, a typical example shown in (a). After the fractal dimension analysis, the grayscale image (a) became a binary image (b) with a threshold of 190, estimated from the slope of the test-fit line as show in (c).

Type of presentation: Poster

MS-1-P-2628 In situ deposited nanocarbon for TEM characterization of zeolite supported metal catalysts

Chu Y.1, Zhang B. S.2, Zhang Q.1, Wang Y.1, Su D. S.2, Wei F.1
1Beijing Key Laboratory of Green Chemical Reaction Engineering and Technology, Department of Chemical Engineering, Tsinghua University, Beijing, China, 2Shenyang National Laboratory for Materials Science, Institute of Metal Research, Chinese Academy of Sciences, Shenyang, China
zhang-qiang@mails.tsinghua.edu.cn

Zeolite supported metal catalysts are widely used while the electron beam sensitive zeolite makes the characterization of the catalysts using electron microscope (EM) difficult. In this contribution, a sacrificial-zeolite specimen preparation (SZSP) technique is developed for the EM analysis of the catalyst. The metal particles are transferred from the zeolite support to the deposited nanocarbon generated in the metal catalyzed hydrocarbon reaction. SAPO-34 zeolite with Al2O3 binder supported Pt catalyst is employed as the model catalyst. Pt catalyzed propane dehydrogenation reaction is carried out to deposit the nanocarbon overlayer which the Pt particles are transferred to as the new support for EM observation. The original catalyst, the deposited nanocarbon and the Pt particles on the new support are characterized by scanning electron microscope (SEM), transmission electron microscope (TEM-EDXS), thermogravimetry/differential thermal analysis (TG-DTA), Raman spectrometry, scanning transmission electron microscope (STEM-EDXS). The coke deposited on SAPO-34 and Al2O3 are of different morphologies and structures. The as-observed distribution of Pt particles on the new support suggests enrichment of Pt on SAPO-34. The shape and size of the Pt particles as well as the strong Pt-SAPO-34 interaction are directly observed. The shape and size of the Pt particles as well as the mechanism of SMSI between Pt and the original support are directly observed. This offers a novel route to monitor the metal size and the interaction between the metal and support, which shed a light on the mystery science of heterogeneous catalyst and provide new insights on the relationship among the structure, active site, and reactivity.


We thank Ling Hu and Tongwei Wu for their help with the transmission electron microscope and the H2-chemisorption experiment. This work was supported by National Basic Research Program of China (973 Program, 2011CB932602), Research Fund for the Doctoral Program of Higher Education of China (No. 20100002110022) and the IMR SYNL-T.S. Kê Research Fellowship.

Type of presentation: Poster

MS-1-P-2659 Magnesium Oxide Nanoparticles Studied by Electron Microscopy

Gärtnerová V.1, Remiášová J.1, Jäger A.1
1Institute of Physics AS CR - Prague 8 (Czech Republic)
gartner@fzu.cz

Magnesium oxide can be used in a wide range of applications covering, for instance, a catalyst in organic chemistry, an adsorbent for a variety of toxic substances, and as a refractory material. There are many routes for preparation of MgO particles but the smallest crystallite size is usually obtained via sol-gel techniques. Here, we present an innovative method for production of magnesium oxide nanoparticles and their microstructure characterization by scanning and transmission electron microscopy (SEM and TEM). MgO can be prepared via a reaction between magnesium (Mg) and methanol (CH3OH) that can be described as

Mg + 2CH3OH → Mg(OCH3)2 + H2            (1)

where the final products are magnesium methoxide Mg(OCH3)2 and hydrogen H2. This reaction, however, practically does not occur at ambient temperatures and must be accelerated either by catalyst or by heating at higher pressures in reflux apparatus. The most common catalyst used is iodine. Although iodine is essential for nutrition, due to its toxicity in elemental form, higher price and problematic manipulation, this element introduces an obstacle for use. Very recently, it was shown that the reaction (1) can be significantly accelerated by Zn in solid solution of Mg. Final product Mg(OCH3)2 is a valuable precursor for production of nanocrystalline MgO (particle size ~5 nm) by simple thermal decomposition (400°C/2h), see Fig. 1.

Beside SEM and TEM, we employed a set of additional experimental techniques such as, mass spectroscopy combined with thermogravimetry, differential scanning calorimetry (DSC) and BET surface area analysis for thorough characterization of MgO nanoparticles prepared by thermal decomposition of Mg(OCH3)2.


Financial support offered by COST MP1103, MEYS LD13069 and LM2011026 is appreciated.

Fig. 1: Fig.1: SEM image of MgO powder (left) and high resolution TEM image of MgO particles (right).

Type of presentation: Poster

MS-1-P-2662 Structural Characterization of Inclined GaN Nanowires Grown in r-plane Sapphire by MBE

Lotsari A.1, Dimitrakopulos G. P.1, Kehagias T.1, Adikimenakis A.2, Komninou P.1, Georgakilas A.2
1Physics Department, Aristotle University of Thessaloniki, GR-541 24, Thessaloniki, Greece, 2Microelectronics Research Group (MRG), IESL, FORTH, P.O. Box 1385, 71110 Heraklion Crete, Greece; and Physics Department, University of Crete, Heraklion Crete, Greece
komnhnoy@auth.gr

We present a structural characterization of inclined GaN nanowires (NWs) on r-plane sapphire grown by plasma-assisted molecular beam epitaxy (PAMBE) under nitrogen rich conditions and after excessive substrate nitridation. The NW size and density were dependent on the nitridation conditions. Photoluminescence measurements of these NWs showed excellent crystal quality and strong emission even at room temperature.
The NWs were grown along the c-axis and subtended a 61o angle to the r-plane sapphire substrate as shown in the TEM image of Fig.1, where the two growth variants are also observed. A rough and discontinuous nonpolar a-plane GaN thin film was formed between the NWs. By combining TEM observations in cross-section and plan view geometries, the crystallographic model of the NWs was constructed. CBED was employed in order to identify the polarity of the NWs. Using HRTEM, the growth origin of the NWs was elucidated.
It was found that the sapphire nitridation pre-treatment enhances substrate roughness forming a stepped surface which provides facets for the nucleation of semipolar nanocrystals [1,2]. These nanocrystals evolve into NWs under N-rich rich growth conditions. Moiré fringes observed close to the interface confirm the presence of such interfacial areas. Analysis of the Moiré fringes along with Bragg filtering of the HRTEM images were performed for the identification of the NW nucleation sites. Geometrical phase analysis (GPA) was also employed at the points of emanation of the NWs since different phases like AlN formation or sapphire protrusions could promote the initial nucleation of the NWs. It was found that the NWs were either grown on semipolar GaN nanocrystals (as shown in Figures 2 and 3) or directly on sapphire steps. No other phase was identified at the growth origin of the NWs.
HRTEM image analysis was also performed at the grain boundaries between the nonpolar a-plane GaN matrix and the NWs in order to identify the accommodation mechanism between the two orientations. The orientation relationship between the two corresponds to a 90o <1210> rotation which ensures high coincident symmetry. Energetically favourable grain boundaries comprised flat terraces with disconnections being introduced between them in order to accommodate the misfit [3].

[1] J. Smalc-Koziorowska et al., Appl. Phys. Lett. 93, 021910 (2008)
[2] J. Smalc-Koziorowska et al., J. Appl. Phys. 107, 073525 (2010)
[3] J. Kioseoglou et al., J. Appl. Phys. 111, 033507 (2012)


This research has been co-financed by the European Union (European Social Fund - ESF) and Greek national funds through the Operational Program "Education and Lifelong Learning" of the National Strategic Reference Framework (NSRF) - Research Funding Program: THALES: Reinforcement of the interdisciplinary and/or inter-institutional research and innovation

Fig. 1: CTEM image of the inclined GaN NWs viewed along the [1101]Al2O3 zone axis. Between the NWs, a rough thin film of nonpolar a-plane GaN is formed

Fig. 2: HRTEM image of a NW grown between two nonpolar a-plane GaN crystals (n-GaN). In the point of emanation of the NW a small semipolar crystallite (s-GaN) is identified

Fig. 3: HRTEM image showing the nucleation site of a thin NW. The base of the NW is surrounded by semipolar GaN

Type of presentation: Poster

MS-1-P-2670 The Reactivity and Structural Dynamics of Supported Metal Nano-Clusters using Electron Microscopy, in situ X-ray Spectroscopy, and Electronic-Structure Theory and Simulations

Yang J. C.1, Johnson D. D.3, Nuzzo R. G.2, Frenkel A. I.4, Ciston J.6, Stach E. A.5, Bonifacio C. S.1, Rehr J.7, Long L.1
1University of Pittsburgh, Pittsburgh, PA, USA, 2University of Illinois at Urbana-Champaign, Urbana, IL, USA, 3The Ames Laboratory (US Department of Energy), Ames, IA, USA, 4Yeshiva University, New York, NY, USA, 5Center for Functional Nanomaterials, Brookhaven National Laboratory, Upton, NY, USA, 6National Center of Electron Microscopy, Lawrence Berkeley National Laboratory, Berkeley, CA, USA, 7University of Washington, Seattle, WA, USA
judyyang@pitt.edu

Heterogeneous catalysis, which impacts the worldwide economy and sustainability due to its ubiquitous role in energy production, depends sensitively on the nano-sized 3-dimensional structural habits of nanoparticles (NPs) and their physicochemical structural sensitivity to the environment. Very small metal clusters can exhibit patterns of reactivity and catalytic activity that are dramatically distinct, and sometimes completely opposite, than behaviors seen with larger clusters. It therefore remains a significant need in research to fundamentally understand and predict the local structure and stability of catalytic materials that can be specifically tailored by design and optimized for an application in technology. Our focus is on the development of integrated characterization and modeling tools and their applications appropriate for carrying out detailed studies on metallic nanoscale clusters comprised of a few to as many as 100 metal atoms. Two state of the art methodologies, synchrotron X-ray absorption fine-structure (XAFS) and quantitative scanning transmission electron microscopy (STEM) methodologies are used and specially designed for determining the 3D structure and structural habits, both individually and as an ensemble, critical for understanding metallic nanoclusters. The experimental work is integrated with theoretical calculations. It is now clear that the structural dynamics of small metallic clusters is actually quite complex. For example, we have shown that the structures of Pt NPs may be both ordered and disordered (Fig 1), depending on its size, support and adsorbates where theoretical simulations predicted and corraborated all of the experimental data (Fig 2). While bulk amorphous Pt is unstable, its existence in NPs is a manifestation of their mesoscopic nature. To bridge the theory-experiment gap, we are producing model Pt/γ-Al2O3 systems using oxidation of NiAl(110) to form a thin film of single crystal γ-Al2O3. To bridge the complexity gap, we are developing an universal environmental cell that is compatible currently with synchrotron XAFS and environmental TEM.


Research was supported by Office of Basic Energy Sciences of U.S. Department of Energy (DE-FG02-03ER15476); ETEM was performed at Center for Functional Nanomaterials, Brookhaven National Laboratory (DE-AC02-98CH10886); statistical studies of NP visibility at NCEM (DE-AC02-05CH11231); and computational support by Ames Laboratory (DE-AC02-07CH11358), operated for DOE by Iowa State University.

Fig. 1: FS-HRTEM images and histograms of Pt NP structures on gamma-Al2O3:(a-c) Disordered NPs < 2nm and (d) ordered 1.2 nm NP, magnified in (e) and its FFT in (f). Histogram of ordered and disordered NPs and on g-Al2O3 (g). Fraction of ordered NPs vs. size (h) for g-Al2O3, with a transition zone of 1.1-2.5 nm.

Fig. 2: Relative DFT energy change for Pt37 in different structural motifs and chemical environments. Electron gain (loss) in yellow (red)] of the lowest-energy structures on C and γ-Al2O3 with(out) H are indicated. Spheres of dark (light) blue show the Pt (H) atoms, and magenta show support atoms.

Type of presentation: Poster

MS-1-P-2709 Unexpected hexagonal CeAlO3 phase formation at low temperatures depending on the hydrogen purity.

Malecka M. A.1, Kepinski L.1
1Institute of Low Temperature and Structure Research, Polish Academy of Sciences, P.O. Box 1410, 50-950 Wrocław 2, Poland
m.malecka@int.pan.wroc.pl

Solid state reaction between highly dispersed CeO2 and alumina provides to crystalline CeAlO3 formation. In normal conditions cerium aluminate crystallizes in the cubic crystal system and its structure can be described in space group Pm-3m (221). This presentation reports results of studies on formation and structure of CeAlO3 in CeO2-Al2O3 system (Ce:Al = 1:10). Samples were prepared by simple impregnation technique from aqueous solution of cerium nitrate. As-prepared samples were dried and pre-heated at 550 oC in static air. Next, CeO2-Al2O3 samples were calcined at temperatures from 850 to 1100 oC in hydrogen flow. Morphological and structural changes upon heat treatment in the reducing atmosphere at evaluated temperatures were studied by HRTEM, XRD and XPS methods. Unexpected, hexagonal phase of CeAlO3 was observed on XRD patterns (see fig. 1) for samples after thermal treatment at lower temperatures (~850 oC). Increasing of heating temperature up to 1000-1100 oC provided to conventional cubic CeAlO3 crystallization (see fig. 1). Moreover, new hexagonal phases of CeAlO3 have been found on HRTEM and SAED images recorded for studied samples (see insets in fig. 1). Depending on the hydrogen purity (concentration of residual oxygen), thereby Ce3+/Ce4+ concentration in cerium aluminate structure, two kind of hexagonal CeAlO3 phases could be formed. It would seem that in case heating CeO2-Al2O3 system in hydrogen flow at low purity, part of Ce4+-ions has not been reduced. As it was found on XRD patterns, formed hexagonal phases are differed in a and b lattice parameters, where c parameter remains stable. Additional, clearly visible differ between samples heated at 850 oC in hydrogen flow (pure and oxygen polluted) was in color of both samples. First of them was gray and second one was dirty-yellow, what could be the proof for presence of Ce4+ ions in sample.


This work was financially supported by the National Science Centre Poland (grant UMO-2011/01/B/ST5/06386). The authors thank Mrs. Z. Mazurkiewicz for valuable help with preparation of the samples and Mrs. E. Bukowska for XRD work.

Fig. 1:  XRD patterns and HRTEM images (in sets) obtained for samples heated at 850 oC in hydrogen flow (A) pure, (B) oxygen polluted and (C) at 1100 oC.

Type of presentation: Poster

MS-1-P-2716 On the structures of sequentially deposited Ru-Au bimetallic catalysts for green chemistry

Wang D.1, Villa A.2, Kotula P. G.3, Prati L.2, Kübel C.1
1Institute of Nanotechnology, and Karlsruhe Nano Micro Facility, Karlsruhe Institute of Technology, Hermann-von-Helmholtz Platz 1, 76344, Eggenstein-Leopoldshafen, Germany, 2Dipartimento di Chimica, Università di Milano, via Golgi 19, I-20133 Milano, Italy, 3Materials Characterization Department, Sandia National Laboratories, Albuquerque, NM 87185-0886, USA
di.wang@kit.edu

Aerobic catalytic oxidation has been appreciated in view of its application in green chemistry in contrast to non-catalytic methods. By controlling the synthesis methodology, the structures of nanocatalysts can be designed towards specific activity and selectivity for the reaction of interest. Recent developments in analytic electron microscopy techniques have enabled the structure and chemistry of individual catalyst particles to be resolved at sub-nanometer resolution. In this presentation, we will focus on two activated carbon (AC) supported Ru-Au bimetallic catalysts synthesized by sequential deposition following a two-step procedure [1] to correlate the elemental distributions and surface decoration with the activity and selectivity of the catalysts.

For Ru@(Au/AC), Ru(III) was reduced with H2 at 80 °C in the presence of preformed Au/AC. Au@(Ru/AC) was prepared by depositing PVA stabilized Au nanoparticles onto a commercial Ru/AC. The catalysts were examined in an image aberration corrected FEI Titan 80-300 electron microscope with conventional EDX detector as well as in a probe aberration corrected Titan 80-200 with in column Super-X EDX detector.

Spectrum imaging with probe corrected STEM and Super-X EDX detector offers high spatial resolution and the opportunity to resolve the components by multivariate analysis before changing the structure during electron beam irradiation. In Fig. 1a, components containing Au and Ru are displayed, forming an Au core-Ru shell structure in the case of Ru@(Au/AC). The high resolution HAADF STEM image in Fig. 1b shows two Au particles with Ru clusters (lower intensity) situated on their surface. In the case of Au@(Ru/AC), we obtained a more inhomogeneous distribution, with the presence of small Ru particles and larger bimetallic particles, as seen in Fig. 2b. Interestingly, the bimetallic particles, as shown e.g. in Fig. 2a, is also composed of an Au core and Ru shell. Both catalysts were tested in oxidation of n-octanol in toluene and oxidation of glycerol in water, respectively. Ru@(Au/AC) shows almost no activity in the former case but is highly active for the latter; while Au@(Ru/AC) behaves in the contrary, being very active for oxidation of n-octanol but shows only limited activity in glycerol oxidation. The TEM results and catalytic tests therefore suggests that Ru is the main active phase in oxidation of aliphatic alcohols and the addition of Au has a detrimental effect on the Ru particles on it. But this Au core-Ru shell structure leads to distinctly enhaced activity in oxidation of water soluable and highly hydrophilic polyols [2].

Reference
[1] D. Wang, A. Villa, F. Porta, D. Su, L. Prati, Chem. Commun. (2006) 1956.
[2] L. Prati, F. Porta, D. Wang, A. Villa, Catal. Sci. Technol. 1 (2011) 1624.


Fig. 1: a) Component maps by multivariate analysis, with red for Au and green for Ru, and b) HAADF STEM image of bimetallic particles showing Ru situated on Au core for the Ru@(Au/AC) catalyst.

Fig. 2: a) Components maps by multivariate analysis, with red for Au and green for Ru, and b) HAADF STEM image of segregated small Ru particles together with big bimetallic particles for the Au@(Ru/AC) catalyst.

Type of presentation: Poster

MS-1-P-2739 Substitutional Gold Doping in ZnO Mesocrystals as Outstanding Catalyst for CO Oxidation

Liu M. H.1, Chu M. W.2, Chen Y. W.3, Kuo J. R.3, Mou C. Y.1
1Department of Chemistry and Center of Condensed Matter Science, National Taiwan University, 2Center of Condensed Matter Science and Center for Microscopy and Nano Analysis, National Taiwan University, 3Institute of Atomic and Molecular Sciences, Academia Sinica
mhliu0811@gmail.com

Zinc oxide is a wide band-gap, n-type semiconductor and also an important material in understanding fundamental optical physics. However, only few research groups utilized ZnO material as a support to do the catalysis researches because of its inertness. Herein, we prepare a new type of ZnO material, namely twin-brush ZnO (TB-ZnO) mesocrystals, which were as a novel support for gold. Gold nanoparticles were deposited on TB-ZnO by means of a modified DP method, forming Au/TB-ZnO catalyst. The catalytic activity of Au/TB-ZnO in CO oxidation was examined from -50 oC to 30 oC (Fig. 1), showing an extraordinary performance in comparison with conventional Au/ZnO catalysts.
To unravel the origin of the outstanding catalysis, the catalyst was characterized using aberration-corrected canning transmission electron microscope (Cs-corrected STEM), high-resolution transmission electron microscopy (HRTEM), X-ray absorption spectroscopy (XAS) and FTIR study of adsorbed CO. Through the analysis of Cs-corrected STEM (Fig. 2) and the refined data from EXAFS (not shown here), it is definitely realized that a number of zinc sites of ZnO sub-lattice were substituted by gold. This unexpected phenomenon was also supported by theoretical calculation. With DFT calculations (4×4×3 ZnO model), it is found that when Zn vacancies exist in the TB-ZnO support, gold atoms can not only diffuse into ZnO but tend to aggregate instead of random dispersion (white arrows in Fig. 2). When the catalyst was treated at 200 oC and further underwent a series of CO reactions, the substituted gold would segregate from interior, resulting in the formation of 2 nm AuNPs in size (Fig. 3). For the analysis of CO adsorption on AuNPs, there are three types of gold species on the support surface, including Au0, Au+ and Au3+. It is significant that the substitution of gold into the ZnO lattice was observed for the first time and further contributes an extraordinary activity in CO oxidation. Consequently, through several examination results, it can be realized that outstanding activity is apparently originated from high active gold atoms and ca. 2 nm gold nanoparticles. This synthetic approach of Au/TB-ZnO can open up a new opportunity to design an excellent catalyst with a finely-controlled particle size.


Fig. 1: CO conversion of Au/TB-ZnO catalysts as a function of low-temperature from -50 oC to 30 oC. The activity revealed at -10 oC was 0.13 molCO•(molAu•s)-1.

Fig. 2: Aberration-corrected HAADF-STEM image of O2-pretreated Au/TB-ZnO catalyst at 200 oC. (The insert is the Fourier transformed result from the corresponding image.)

Fig. 3: The HRTEM image from ultramicrotome slice of Au/TB-ZnO with O2 pretreatment underwent a series of CO oxidation reactions ramping from -20 oC to 90 oC.

Type of presentation: Poster

MS-1-P-2816 Hydrothermal synthesis of twin-based multiply branched rutile-type TiO2

Jordan V.1, Podlogar M.1, Rečnik A.1
1Department for Nanostructured Materials, Jožef Stefan Institute, Ljubljana, Slovenia
vanja.jordan@ijs.si

Production of nanostructures, especially 3D branched architectures with controlled morphology and size, has been the focus of research in the recent years. One of such materials that can be grown in complex branched morphologies is rutile (TiO2). While being applicable in variety of applications,1 rutile is known to be prone to twinning, which can be exploited as a basic structural element for growing branched structures.2 Branched structures of rutile-type TiO2 are obtained in an acidic medium with the use of metalorganic Ti-precursor. Several nucleation mechanisms have been proposed in the literature, however the true nature of branching is yet to be explained.3-5 Several synthesis routes have been tested to study twinning of rutile. Following the synthesis pathway suggested by Tomita et al. (2006),3 1st generation of twinning was obtained. Briefly, titanium powder was dissolved in H2O2 and NH3, to produce titanium oxyhydroxide. This was followed by a ligand exchange reaction with glycolic acid, to form Ti-glycolato complex, which was hydrothermally treated at 200 °C for 1-24 hours to obtain nanocrystalline rutile. 2nd generation of twins was obtained by a subsequent hydrothermal treatment of already existing twins by the addition of Ti-butoxide in strongly acidic medium. Another approach, following Zhou et al. (2011),4 yielded 2nd generation of twins in a single synthesis step. Two generations of twinning were obtained through hydrothermal synthesis in acidic medium. As precursor, Ti-butoxide was used and dissolved in 7-10M HCl aqueous solution. The syntheses were conducted at different temperatures and processing times. Morphology and composition of the products were characterized by SEM and TEM. The first route yielded clusters of twinned rutile crystals (Fig. 1). Electron diffraction study of twin relations indicated that the products are composed of (101) and (301) twins, with the characteristic angles of 114° and 55°, respectively.2 The second synthesis route led to formation of complex-branched structures, with abundant twinning and other types of intergrowths (Fig. 2a). Unlike in the first synthesis route, the rutile crystals here are composed of numerous parallel rutile fibers lined along the crystallographic c-axis (Fig. 2b and 2c). In crystals that are oriented along the c-axis inherent porosity can be observed, which might be a consequence of imperfect fiber alignment. Further, the presence of anatase phase, as suggested by several authors,5 could not be confirmed, nevertheless some unidentified reflections that could correspond to this TiO2 phase are observed in electron diffraction patterns (Fig. 2d). Pores and imperfect alignment of the fibers indicate the possible mechanism of rutile formation and branching.


1. C. Cheng, H.J. Fan, Nano Today 7 (2012) 327-343.
2. N. Daneu, H. Schmidt, A. Rečnik, W. Mader, Am. Mineral. 92 (2007) 1789-1799.
3. K. Tomita, V. Petrykin, M. Kobayashi, et al., Angewandte Chemie 45 (2006) 2378-2381.
4. W. Zhou, X. Liu, J. Cui, D. Liu, J. Li, , et al., CrystEngComm 13 (2011) 4557-4563.
5. D. Li, F. Soberanis, J. Fu, et al., Crystal growth & Design 13 (2013) 422-428.

Fig. 1: (a) {101} and {301} twins of rutile obtained from Ti-glycolato complex, (b) TEM image of (301) twin. Fig. 2: (a) Rutiles synthesized from Ti-butoxide. (b) Fibrous rutiles coinciding in twin-type orientations. (c) Close-up of rutile roughly aligned fibers. (d) Single rutile fiber in [001] projection. Diffraction rings mainly correspond to rutile.

Type of presentation: Poster

MS-1-P-2820 Growth mechanism of Ag catalysed InAs nanobelts grown in MOCVD

Xu H.1,3, Gao Q.2, Tan H. H.2, Jagadish C.2, Zou X.3, Zou J.1,4
1Materials Engineering, University of Queensland, 2Department of Electronic Materials Engineering, Australian National University, 3Department of Materials and Environmental Chemistry, Stockholm University, 4Centre for Microscopy and Microanalysis, University of Queensland
h.xu5@uq.edu.au

Zinc-blende structured Au catalysed epitaxial III-V semiconductor nanowires prefer to grow along the [111]B direction featuring approximately hexagonal cross sections with side-wall facets dominated by six {112} planes or six {110} planes. In contrast, only a hand-full of studies in non-Au catalysed 1-D nanostructure growth have been reported up to date. In this study, by using Ag as catalysts, we demonstrate to grow 1-D InAs nanobelts grown along <112>B directions. Through detailed electron microscopy characterizations, the growth mechanism of these InAs nanobelts is explored.

Commercially available 40nm Ag nanoparticles were used to grow epitaxial 1D InAs nanostructures on GaAs (111)B substrate in a MOCVD reactor. The growth was carried out using trimethylindium and arsine as the group III and group V precursors, respectively. A growth temperature of 500°C and a V-III ratio of 2.9 were selected as the key growth parameters.

Fig. 1a is an overview SEM image and shows the general morphology of as-grown Ag-catalyzed 1-D nanostructures. From the enlarged SEM image (Fig. 1b), inclined nanostructures show belt-like morphology. Fig. 1c is a side-view SEM image of a typical nanobelt, from which the inclined angle of the nanobelt is measured as ~70° (when the electron beam is parallel to a <110> direction and perpendicular to the inclined nanobelts), so that its axial direction can be crystallographically determined to be along <112>B directions. It is of interest to note that "steps" can be found on the top facets of the nanobelt, whereas the bottom surface is relatively smooth. Fig. 1d and e are SEM images of a typical nanobelt viewed from top-view and edge-on view, respectively. As can be seen from the top view (Fig. 1d), the nanobelt is tapered. When the nanobelt is viewed along its axial direction (Fig. 1e, ~20° tilt), the sidewall facets of the nanobelt shows a rectangular shape with two short {110} facets and two long facets developed from {111} planes. The top facet contains many surface "steps" consistent with other SEM observations.

By using a number of TEM techniques, the structure, sidewall facets, chemical composition, planner defects and nanobelt/catalyst interface of the as-grown nanobelts are investigated in detail, from which a growth schematic of Ag-catalyzed InAs nanobelt is proposed and shown in Fig. 2. During Ag catalyzed 1-D nanostructure growth, the catalyst promotes two processes: (1) collects group III growth material and (2) transports the growth material to the growth front of the nanostructure and nucleates the growth. It is believed that the formation of the “steps” and the planer defects is closely related to the In concentration in the catalyst during the nanobelt growth.


The Australian National Fabrication Facility and Australian Microscopy and Microanalysis Research Facility are gratefully acknowledged. The Wenner-Gren Foundation and 3DEM NATUR project at MMK, Stockholm University are also gratefully acknowledged.

Fig. 1: (a) SEM overview of the Ag-catalysed InAs nanostructures. (b) SEM image of a typical nanobelt. (c) Side-view of a typical nanobelt. (d) and (e) a typical nanowire viewed when the beam is perpendicular to the substrate and parallel to the nanobelt axial direction.

Fig. 2: Schematic illustration of the “step” and defect formation. Perturbations in In concentration introduces new facets during the nanobelt growth.

Type of presentation: Poster

MS-1-P-2892 Microstructure Study of W1-xMox O3 0.33H2O for Tunable Wavelength Absorption

Arzola-Rubio A.1, Ornelas C.1, Antúnez-Flores W.1, Collins-Martínez V.1, Paraguay-Delgado F.1
1Centro de Investigación en Materiales Avanzados S. C., CIMAV Miguel de Cervantes 120, Chihuahua, Chih. México. CP 31109.Ar
alejandro.arzola@cimav.edu.mx

Hydrated tungsten oxide WO3 0.33H2O has been studied extensively due to its electronic and optoelectronic properties, it has has an enormous potential application ranging from condensed-matter physics to solid-state chemistry [1], such as photo-electrochemical energy conversion, gas sensors, photocatalysis, lithium-ion batteries, solar cells [2]. Tremendous effort has been dedicated to the synthesis, solid solution mechanism and property investigation of W1-xMox O3 0.33H2O over the past years. This material showed improved electrochromic, gas sensing, catalytic, lithium ion transport, and photocatalytic properties [3] when compared with their single oxide WO3 and MoO3. Recently, Zhou et al. were capable of modulate the band gaps of the W1-xMox O3 0.33H2O materials with different Mo/W ratio values [4]. We synthesized a series of W1-xMoxO3 0.33H2O nano/microstructures with controlled stoichiometry (x = 0, 0.25, 0.50, 0.75). With gradual increase of Mo content, we narrowed the band gap from 2.61 to 2.10 eV. This result is better than Zhou et al. but in our case, we use friendly to the environment chemical precursors such as ammonium heptamolybdate and ammonium metatungstate instead of metal powders.
Figure 1a shows a SEM image for orthorhombic WO3 • 0.33H2O, the particles have an average length and wide of 100 and 50 nm respectively. Figure 1b is a bright field TEM image and inset is the SAED pattern for the W75Mo25O3 0.33H2O compound, the diffraction spots were to the orthorhombic structure and it has [3,0,-1] zone axis, which is a single crystal. Figure 1c corresponds to TEM image of solid solution W50Mo50O3 0.33H2O and its corresponding SAED pattern which indexed to orthorhombic structure too, with [2,-1,0] zone axis, this pattern was from a particle labeled with Z1. In the case of this compound, there is different size of particles which measurements are 160nm length and 80nm width. Figure 1d shows SEM image of the compound W25Mo75O3 • 0.33H2O. It can be notice hexagonal flake-like particles with lengths of ∼150nm and widths of ∼70nm.
We were able to measure and characterize our compounds with advanced microscope techniques such as SEM and TEM getting information about structure by SAED patterns where we were capable to index all spots and determine the zone axis and figure out all crystalline path growths of these type materials.

References
1. Zheng HD, Ou JZ, Strano MS, Kaner RB, Mitchell A, Kalantarzadeh K (2011) Adv Funct Mater 21:2175.
2. Turyan I, Krasovec UO, Orel B, Saraidorov T, Reisfeld R, Mandler D (2000) Adv Mater 12:330.
3. Baeck, S. H.; Jaramillo, T. F.; Jeong, D. H.; McFarland, E. W. Chem. Commun. 2004, 390–391.
4. Zhou, L.; Zhu, J.; Yu, M.; Huang, X.; Li, Z.; Wang, Y.; Yu, C. J. Phys. Chem. C 2010, 114, 20947−20954.


Fig. 1: FIG. 1 SEM, TEM and SAED images of W1-xMoxO3 • 0.33H2O (x = 0, 0.25, 0.50, 0.75) nano/microstructures

Type of presentation: Poster

MS-1-P-2896 Microstructural and Electrical Characterization of Langmuir-Blodgett Films of Ultrathin Semiconductor Nanoheterowires

Ashokkumar A. E.1, Li H.1, Hayat A.1, Dalui A.2, Jafri H.4, Sarma D. D.3, Acharya S.2, Leifer K.1
1Department of Engineering, Applied Materials Science, Uppsala University, Sweden, 2Centre for Advanced Materials, Indian Association for the Cultivation of Science,Kolkata, India, 3Solid State and Structural Chemistry Unit, Indian Institute of Sciences Bangalore, India, 4Electrical Engineering Department, Mirpur University of Science and Technology, Pakistan
anumolea@gmail.com

Semiconductor heterostructures with suitable band alignment which can promote charge separation is interesting for various applications in electronics and optoelectronics. Chemical synthesis methods are developed for obtaining complex nanostructures including semiconductor hybrids of few nanometers sizes leading to quantum confinement and resultant unique properties. The assembly of such structures on substrates as thin films can facilitate its characterization and further applications. In this work, heterostructures consisting of ZnS rods and CdS dots with sequential alignment synthesized by wet chemical synthesis is investigated in this direction. Ultrathin superlattice nanowires of ZnS-CdS were developed into monolayer thin films on SiO2/Si substrates using Langmuir-Blodgett (LB) technique. The various aspects of the assembly are investigated. The morphology of the film and the orientation of the nanowires in the LB film were studied by microscopy techniques including Scanning Electron Microscopy, Transmission Electron Microscopy and Atomic Force Microscopy. Due to the sub 2 nm size of the nanowires, Transmission Electron Microscope is mandatory to observe the alignment of the individual nanowires and also to observe the components of the hetero- nanowire. Transfer of the Langmuir-Blodgett film on to Cu grids facilitates the observation of the film under TEM.
The electrical characterization of such sub 2-nm wires is challenging as the fabrication of electrical contacts is nontrivial. A nanoplatform where the electrodes with a spacing of ~ 50 nm fabricated using lithography and Focussed Ion Beam technique was developed for this purpose. The LB film was deposited onto the substrate with such prefabricated electrodes. The electrical and optical properties of the LB films are presented.
Thus the present work investigates the role of microscopy techniques in characterizing LB films of nanomaterials and also attempts to bridge the gap between wet chemical synthesis of semiconductor nanowires and their device fabrication.


The Swedish Foundation for International Cooperation in Research and Higher Education (STINT) is acknowledged for research grant.

Fig. 1: SEM image of Langmuir-Blodgett film of ZnS-CdS heterowires

Fig. 2: AFM image of Langmuir-Blodgett film of ZnS-CdS on SiO2/Si substrate

Fig. 3: TEM bright field and high resolution images of the ZnS-CdS Langmuir-Blodgett film transferred to Cu grid

Type of presentation: Poster

MS-1-P-2899 Self-assembled Supraballs by Spherical Confinement

Wang D.1, de Nijs B.1, Dussi S.1, Smallenburg F.1, Filion L.1, Pietra F.2, Meeldijk J. D.3, van Dijk-Moes R.2, van Huis M.1, Vanmaekelbergh D.2, Imhof A.1, Dijkstra M.1, van Blaaderen A.1
1Soft Condensed Matter, Debye Institute for Nanomaterials Science, Utrecht University, Utrecht, The Netherlands, 2Condensed Matter and Interfaces, Debye Institute for Nanomaterials Science, Utrecht University, Utrecht, The Netherlands, 3Electron Microscopy Group, Utrecht University, Utrecht, The Netherlands
d.wang@uu.nl

Colloidal crystalline supraparticles, which are self-assembled from size- and morphology- controlled nanoparticles, can exhibit many different interesting meta-materials properties, while still having a size in the colloidal domain and thus the possibility with a second self-assembly step to form other interesting structures. An example is colloidal crystal lattices with Bragg-reflections for visible light. In research aimed at making colloidal crystalline supraparticles by having monodisperse spherical nano- (and micron-sized) colloids crystallize in slowly evaporating oil emulsion droplets, we discovered icosahedral symmetry in the resulting dried colloidal crystals for particle number less than ~100,000. Subsequent computer simulations confirmed the icosahedral symmetry even in the absence of any attractions and thus are entropically favored over the face-centered-cubic (FCC) crystal structure that is stable in the bulk.1

We also extended the spherical confinement method to a binary particle system and an anisotropic rod-like particle system. For instance, 6.2 nm Au nanocrystals and 22.0 nm FexO/CoFe2O4 nanocrystals were used to synthesize binary supraballs. CdSe/CdS quantum rods can also self-assemble into supraballs. By tuning the concentration of the nanocrystals, supraballs with different structures and sizes can be obtained. After freeze drying, the structure of the supraballs was studied with high-angle annular dark-field scanning transmission electron microscopy (HAADF STEM) and secondary electron scanning transmission electron microscopy (SE STEM). Work is in progress to study the structures of the more complex supraballs by HAADF STEM tomography.

Reference

(1) Bart de Nijs et al. submitted.


Fig. 1: SE STEM image of supraballs with rhombicosidodecahedron structure made from 6.0 nm FexO/CoFe2O4. (Scale bar 50 nm)

Fig. 2: HAADF STEM image of binary supraballs made from 6.2 nm Au and 22.0 nm FexO/CoFe2O4. (Scale bar 50 nm)

Fig. 3: SE STEM image of binary supraballs made from 6.2 nm Au and 22.0 nm FexO/CoFe2O4. (Scale bar 50 nm)

Fig. 4: HAADF STEM image of supraballs made from CdSe/CdS quantum rods. (L=53.2 nm, D=4.1 nm) (Scale bar 50 nm)

Type of presentation: Poster

MS-1-P-2901 On the correct grain size characterization of nanometric polycrystalline materials

Pinto A. L.1, Gama G. R.1, Silva A. M.1
1Centro Brasileiro de Pesquisas Físicas (CBPF) 1
pinto@cbpf.br

There has been an increasing interest in polycrystalline materials with nanometric grain size. These materials may be produced through thin film technology1, electrodeposition or severe plastic deformation. The grain size is the most basic microstructural parameter that is usually described in a sample. Some properties may be predicted directly from the grain size, but as to nanomaterials, they have called for the revision of some concepts2. One of the most common ways of characterizing the grain size at a TEM is through dark field images. As the diffraction contrast obtained in bright field does not allow a proper description of the microstructure, the objective aperture is used for illuminating grains within a determined range of crystallographic orientations given by the size of the objective aperture used. In this work we have used this approach and then we have compared to the microstructures obtained through nanodiffraction mapping. Cu thin films were deposited over an oxidized Si(100) substrate through DC magnetron sputtering with a 5mTorr Ar plasma. The substrates were positioned at 10cm from a Cu target. Four different currents were used for the deposition: 30, 75, 100 and 125mA. Depositon times were calibrated to obtain 100nm thin films. TEM analysis was performed at a Jeol 2100F equipment with Nanomegas ASTAR nanodiffraction mapping system. Dark field images were obtained from each sample, which resulted in the grain mean size of 89, 22, 20 and 15nm respectively for the 30, 75, 100 and 125mA samples. Fig. 1 presents a bright field image and a dark field image from the 100 mA sample. The grain size estimation through dark field images or even from bright field images may lead to serious misinterpretation of the results. Using a 30μm objective aperture the crystal orientation deviation is much greater than the 15o commonly used for defining a grain; on the other side using a 5μm aperture almost only one grain is measured per picture. Using nanodiffraction mapping we have obtained 7, 16, 15 and 17nm as grain mean size, respectively. Fig. 2 presents an orientation map and an Index map from 100mA sample at which it is possible to notice the distinct (100) and (110) textures. This is the reason for the poor distinction between neighbor grains. It is also clear there are some larger grains that seem to have grown at the expense of their smaller neighbors. These larger grains have a different orientation from the predominant texture. The use of nanodiffraction mapping at the TEM not only gives a better evaluation of the grain size but also gives much more information about the microstructure.

References

1 – Hodge, A.M. et al. Materials Science and Engineering A 429 (2006) 272–276.

2 – MEYERS, M. A. et al. JOM, April (2006) 41-48.


The authors thank to CAPES, CNPQ and FINEP for financial support of this work.

Fig. 1: Fig. 1 – (a) Bright field image and (b) dark field image from sample 100 mA obtained with 5 μm objective aperture.

Fig. 2: Fig. 2 – (a) orientation map ((001) – red, (101) – green and (111) – blue color reference), (b) reliability map and index map from 100 mA sample.

Type of presentation: Poster

MS-1-P-2977 qHRTEM analysis of the (211)B In(Ga)As QDs/GaAs heterostructure

Florini N.1, Kioseoglou J.1, Dimitrakopulos G. P.1, Walther T.2, Hatzopoulos Z.3, Pelekanos N. T.3, Kehagias T.1
1Physics Department, Aristotle University of Thessaloniki, GR-54124 Thessaloniki, Greece, 2Department of Electronic and Electrical Engineering, University of Sheffield, Mappin St, Sheffield S1 3JD, UK, 3Materials Science & Technology and Physics Departments, University of Crete and IESL/FORTH, GR-71003 Heraklion, Greece
kehagias@auth.gr

The nanostructural properties and strain state of In(Ga)As quantum dots (QDs), embedded in GaAs (211)B by plasma-assisted molecular beam epitaxy (PAMBE), such as shape and dimensions, existence of associated dislocations, thickness of the wetting layer and the possibility of interdiffusion or segregation phenomena in the QDs, were investigated by high-resolution and scanning transmission electron microscopy (HRTEM-STEM) methods. HRTEM imaging showed that the wetting layer thickness does not exceed 2 monolayers. Moreover, the embedded QDs are not associated with any linear defects, suggesting fully strained and optically active nanostructures. However, the shape and dimensions of QDs cannot be precisely extracted, due to the dark strain contrast surrounding the QDs. Conversely, STEM annular bright-field (ABF) imaging revealed that In(Ga)As QDs are elongated along the [-111] direction [Fig. 1(a)].

Quantitative measurements of the local strain in the QDs from HRTEM images have been performed, by the geometrical phase analysis (GPA). GPA is used to determine the strain field in a HRTEM image with respect to a reference region. The GPA lattice strain eg = (ααGaAs)/αGaAsα being the in-plane or the out-of-plane strained values of In(Ga)As QDs, is defined relative to the unstrained GaAs underlayer corresponding values, which was taken as reference. GPA measurements using a g/2 mask, showed that the in-plane strain approximates 0, implying a fully registered heterostructure at the interface, as also illustrated in the HRTEM images of two individual QDs in Figs. 1(b) and (c). Assuming a biaxial strain state of the QDs, the corresponding GPA strain surface plots of two QDs along the growth direction are also shown. The quantitative analysis of the InxGa(1-x)As QDs on GaAs resulted in chemical composition maps of the investigated QDs. The In content was found to vary from 52% at the base of the QDs to almost 100% at the apex area [Figs. 1(b) and (c)], implying possible Ga segregation in the initial stages of QD growth and formation of an InGaAs alloy. However, since the QDs are entirely embedded in GaAs, possible influence from the matrix cannot be excluded.

Moreover, the samples were analyzed by energy dispersive X-ray (EDX) spectroscopy in order to estimate the chemical composition on the InAs QDs in comparison to the quantitative GPA measurements.


Research co-financed by the European Union (European Social Fund–ESF) and Greek national funds through the Operational Program "Education and Lifelong Learning" of the National Strategic Reference Framework (NSRF)–Research Funding Program: THALES, project NANOPHOS.

Fig. 1: (a) ABF STEM image depicting the embedded InAs QDs. (b) & (c) HRTEM images of embedded InAs QD along the [0-11] zon eaxis and the corresponding GPA strain surface plots along the growth direction.

Type of presentation: Poster

MS-1-P-3045 Atomic-scale characterization of light-emitting diodes based on ordered InGaN nanocolumns

Torres-Pardo A.1, 2, Bengoechea-Encabo A.3, Albert S.3, Sánchez-García M. A.3, López-Romero D.3, Gacevic Z.3, Calleja E.3, González-Calbet J. M.1, 4
1Dept.Química Inorgánica, Facultad de Químicas, Universidad Complutense de Madrid, 28040, Madrid, Spain, 2CEI Campus Moncloa, UCM-UPM, Madrid, Spain, 3ISOM-Dept. Ingeniería Electrónica, ETSIT, Universida Politécnica, 28040 Madrid, Spain, 4Centro Nacional de Microscopía Electrónica CNME, 28040 Madrid, Spain
jgcalbet@ucm.es

The potential of InGaN alloys to generate light emission in the UV to IR range makes them an ideal choice for light emitting diodes (LED) covering the whole visible range and beyond. LEDs based on self-assembled nanocolumns (NCs) with InGaN/GaN disks constitute an alternative to conventional LED planar devices which major limitation is a strong reduction in efficiency at high current injection [1]. However, the efficiency and reliability of LEDs based on self-assembled NCs are hindered by a strong dispersion of electrical characteristics among individual nanoLED. Polychromatic emission derives from an inhomogeneous distribution of indium concentration, changes in the NCs geometry and the inherent tendency of InGaN alloys to develop composition fluctuations as a function of the polarity of the growth crystallographic planes [2]. The recent development of selective area growth of NCs by molecular beam epitaxy has allowed the achieving of highly homogeneous and controllable GaN/InGaN NCs with improved crystalline quality and higher control over the indium distribution [3].

In this work, we present results on the characterization of blue, green and yellow LEDs based on ordered NCs with InGaN active layers (figure 1). The detailed structural characterization of the nanostructures has been performed by scanning transmission electron microscopy (STEM) carried out on an aberration-corrected JEOL-JEMARM200 microscope [4]. High crystal quality of the NCs is set by the analysis of atomically-resolved high angle annular dark field (HAADF) images, while the polarity determination of the semiconductor NCs is followed by locating the nitrogen atomic columns in annular bright field (ABF) images (figure 2). The indium distribution within the InGaN disks is studied by EDS elemental mapping, confirming homogeneity of the InGaN layers. The optical response is evaluated from the analysis of electroluminescence spectra.

[1] E. Kioupakis, P. Rinke, K. T. Delaney, C. G. Van de Walle, Appl. Phys. Lett. 98, (2011), 161107.

[2] A. L. Bavencove, G. Tourbot, J. Garcia, Y. Desieres, P. Gilet, F. Levy, B. Andre, B. Gayral, B. Daudin, and L. S. Dang, Nanotechnology,22, (2011), 345705.

[3] S. Albert, A. Bengoechea-Encabo, M. A. Sanchez-Garcia, X. Kong, A. Trampert, E. Calleja, Nanotechnology 24, (2013), 175303.

[4] Y. Li, L. Zhang, A. Torres-Pardo, J.M. González-Calbet, Y. Ma, P. Oleynikov, O.Terasaki, S. Asahina, M. Shima, D. Cha, L. Zhao, K. Takanabe, J. Kubota, K. Domen, Nature Communications, 4, (2013), 2566.


Authors acknowledge financial support by the Spanish projects MAT2011-23068 and CSD2009-00013. Research by A.T.P. has been supported by PICATA postdoctoral fellowship CEI Moncloa.

Fig. 1: (a) Cross-sectional SEM images of a representative sample. (b) Top view SEM picture. (c) Low magnification TEM image of GaN/InGaN nanocolumn.

Fig. 2: (a) Schematic draw of [010] GaN structure with wurzite-type structure. (b) Atomically resolved High Angle annular Dark Field (HAADF) image and (c) corresponding Annular Bright Field (ABF) image revealing the Ga and N atomic columns on the wurzite-type structure.

Type of presentation: Poster

MS-1-P-3074 In-situ monitoring gas-solid reaction in nanoparticles by a static nanoreactor in the TEM

Wu M.1, Shen C.1, Zandbergen H.1
1Kavli Institute of NanoScience, HREM, Delft University of Technology, Delft, The Netherlands
m.y.wu@tudelft.nl

There is a big interest in realizing TEM experiments beyond the ~ 20 mbar pressure regime that can be achieved by ETEM. This is possible using a nanoreactor concept, in which the gas is enclosed along the beam direction by two very thin membranes of for instance SiN (1). With this approach Yokosawa (2) showed that pressures up to 4.5 bar are obtainable. An obvious question in this approach is what kind of resolution can be obtained, given that the resolution limit is now no longer set by the electron microscopes, provided these are equipped with aberration correctors. In the approaches reported by Creemer and Yokosawa gas tubes embedded in the TEM holder were used to lead gas to the nanoreactor and this requires a sophisticated gas supply system. Such a system of gas inlet and outlet and gas regulation could be hazardous for both the TEM and the TEM-user if there exists a leakage.
Here we present a new type of gas holder, which we will call a static gas holder, which has also two windows but no dynamic gas supply system (figure 1). Instead, the holder has a separable tip, which contains an airtight chamber that can store gas with volume of 1.5 to 10 cubic millimeter. Gas is loaded in or pumped out through a valve in the tip (see figure 1). Similar to the dynamic nanoreactor, it consists of two silicon chips, which have a low stress 400 nm thick SiN membrane of for instance 400 um x 400 um. One of the two membranes contains a Pt heater spiral and both of the membranes contain with 5-20 small thin SiN membranes “windows” with thickness of 10-20 nm. The windows of the top and bottom chip have to be aligned to be overlapping, such that a sample on top of one of the windows can be investigated by transmission electron microscopy. In this system, the temperature can be changed within a second over for instance 100°C with low specimen drift. Since the gas volume is very small, no harm to the gun part of the TEM when there is a sudden release of all gas inside the nanoreactor and the tip. An obstacle for high-resolution imaging can be the contamination in the system, which can originate from sample, chips, gases, O-rings etc. We demonstrate that when contamination is minimized, the resolution of the system can reach the resolution limitation of the microscope at gas pressures of e.g. oxygen of at least 0.6 bar (figure 2 and figure 3).

1. J.F. Creemer, S. Helveg, G.H. Hoveling, S. Ullmann, A.M. Molenbroek, P.M. Sarro, H.W. Zandbergen, Ultramicroscopy 108 (2008) 993– 998.
2. Tadahiro Yokosawa, Tuncay Alan, Gregory Pandraud, Bernard Dam, and Henny Zandbergen Ultramicroscopy 112(2012)47–52.


This work is supported by ERC NEMinTEM Project 267922.

Fig. 1:  (a) Image of the static gas holder and an enlarged view of the tip part. (b) Cross sectional sketch of the holder tip disassembled and assembled onto the holder.

Fig. 2: (a) PdOx nanoparticles in O2 with pressure of 0.645 bar at 500 °C. (b) FFT of image (a), the white circle is 1 Å. The triangle indicates a diffraction spot with a d-spacing of 0.88 Å.

Fig. 3: Pd nanoparticles in H2 with a pressure of 0.52 bar at 200 °C. (b) FFT of image (a), the white circle is 1 Å. The triangle indicates a diffraction spot with a d-spacing of 0.85 Å.

Type of presentation: Poster

MS-1-P-3055 Atomic resolution imaging of SiOx quantum dots in diamond by TEM/STEM

Sung K. T.1, Wei L. L.1, Chiu K. A.1, Chang L.1
1Department of Materials Science and Engineering, National Chiao Tung University, Hsinchu, Taiwan
lichang@cc.nctu.edu.tw

Si-based quantum dots (QDs) have received intensive studied in the past decade because of their light emitting properties. Most Si-based QDs are fabricated to be embedded in silicon dioxide. Here we show that amorphous SiOx QDs can be embedded within single crystalline diamond. A SiOx film was deposited on a 2 mm (111) single crystal diamond substrate by sputtering. The sample was then treated with hydrogen plasma to form QDs, followed by microwave plasma chemical vapor deposition (CVD) of homoepitaxial diamond. Cross-sectional TEM specimens in diamond <110> orientation were prepared by focused ion beam. TEM/STEM observation was carried out in a JEOL ARM200F microscope with STEM annular dark field (ADF) image resolution of ~ 0.8 Å.

Figures 1(a) and (b) show typical STEM BF and ADF images, respectively, in which Si-based QDs are seen in dark and bright contrast. The size of the QDs is ~ 2-6 nm. The QDs covered with CVD diamond can be observed in the HRTEM image (Fig. 2 ). It can be seen that lattice fringes continuously cross over the areas of the QDs, illustrating diamond homoepitaxy. The atomically resolved STEM-ADF image in Fig. 3 shows a QD (~ 2 nm size) in very bright contrast superimposed with diamond atomic columns, indicating that the QD is amorphous and embedded in single crystalline diamond. Furthermore, only Si, O, and C peaks are detected in x-ray EDS measurements on those QDs, suggesting that the QDs are SiOx (x ~ 0.6). Also, photoluminescence measurements show light emitting wavelength at a peak > 520 nm.


The work was supported by National Science Council, Taiwan, R.O.C. under Contract No. 101-2221-E-009-049-MY3.

Fig. 1: (a) STEM-BF and (b) STEM-ADF images obtained from an interfacial region.

Fig. 2: HRTEM image showing diamond lattice fringes.

Fig. 3: STEM-ADF image in atomic resolution showing SiOx QDs in diamond.

Type of presentation: Poster

MS-1-P-3065 Atomic surface diffusion on Pt nanoparticles quantified by high-resolution transmission electron microscopy

Schneider S.1,2, Surrey A.1,2, Pohl D.1, Schultz L.1,2, Rellinghaus B.1
1IFW Dresden, Institute for Metallic Materials, Dresden, Germany, 2TU Dresden, Institute of Condensed Matter Physics, Dresden, Germany
sebastian.schneider@ifw-dresden.de

The continuing development of aberration-corrected high-resolution transmission electron microscopy (HRTEM) led to the possibility to study the structure of specimens at the atomic scale in great detail and with highest precision [1]. Besides the determination of these static structural information, dynamic changes of the specimen, e.g., due to the impact of the imaging electron beam, are often observed in (scanning) TEM [2], in particular when working at high doses that are frequently mandatory to investigate the structural and chemical properties. Such dynamic phenomena which are mostly related to unwanted radiation damage [3], may also occur spontaneously without the electron beam and are most easily observable on surfaces. It is, however, still an open question whether the observed atomic surface diffusion, which is the main underlying mechanism of most of these processes, can be solely ascribed to the physical properties of the material, or if and to which extent it is rather promoted by the impact of the imaging electron beam.
Recently, a HRTEM method was introduced that allows for the quantitative estimation of the surface diffusion coefficient, and it was shown that this method can be used to quantify the surface self-diffusion on Au nanoparticles [4]. The method is based on the analysis of temporal fluctuations in the occupancy of surface atomic columns.
Thus in the present study, the motion of atoms at the surfaces of Pt nanoparticles is characterized by means of aberration-corrected HRTEM with the resolution of individual atomic columns. Fig. 1 shows six out of 31 example images of a Pt nanoparticle on a holey amorphous carbon film supported by a copper grid. It can be seen that during the time delay of 0.8 s between the acquisition of two images, some atom columns are emptied while other neighboring columns are filled. Even though atoms are indistinguishable and only a two-dimensional projection of the three dimensional diffusion can be registered in the TEM, the applied method is capable of a quantitative estimation of the diffusion coefficient from the temporal sequence of HRTEM images.
The coefficient of the surface self-diffusion of Pt as derived with this novel approach turns out to be in very good agreement with the results of both experimental [5] and theoretical [6] studies.

[1] Urban, K.W., Science 321 (2008), p. 506 - 510.
[2] Bals, S. et al., Nature Communications 3 (2012), 897.
[3] Egerton, R. et al., Micron 35 (2004), p. 399 - 409.
[4] Surrey, A. et al., Nano Letters 12 (2012), p. 6071 - 6077.
[5] Bassett, D., Webber, P., Surface Science 70 (1978), p. 520 - 531.
[6] Feibelman, P.J., Physical Review Letters 81 (1998), p. 168 - 171.


Fig. 1: Temporal series of images of a Pt particle. The delay between two subsequent images is 0.8 s. “o”/”e” = occupied/empty atomic column. From left to right, the culumn occupation changes as follows: turquois: o, o, o, o, e, e; white: o, o, e, e, e, e; green: e, e, o, o, e, e; blue: o, o, e, e, o, e; yellow: o, o, o, o, o, e; red: e, o, e, o, e, o.

Type of presentation: Poster

MS-1-P-3070 Analytical TEM study of Au-Ag bimetallic catalysts prepared by solid grinding method

Akita T.1, Maeda Y.1, Kohyama M.1
1Research Institute for Ubiquitous Energy Devices, National Institute of Advanced Industrial Science and Technology (AIST)
t-akita@aist.go.jp

Gold catalysts exhibit high catalytic activity for low temperature CO oxidation [1]. It has been reported that active site is interface between Au and metal oxide supports [2]. Activation of oxygen at the interface seems to be important step for the CO oxidation reaction. While CO molecules are adsorbed on Au surface, the oxygen seems to be activated at the interface between Au and metal oxide support. Details of mechanism of whole reaction is not clarified yet. On the other hand, addition of Ag also improves the activity of Au catalysts [3,4]. It is significant to investigate the mechanism of improvement of the activity by adding Ag in order to study the general mechanism of oxygen activation for low temperature oxidation. In this study, Au-Ag bimetallic catalyst was prepared by solid grinding method [5], and fine structure of Au-Ag bimetallic catalysts was investigated by aberration corrected TEM/STEM. The distribution of Au and Ag in one nanoparticle was also investigated by EDS.

Au-Ag bimetallic catalysts was prepared by simultaneous solid grinding method using Me2Au (acac: acetylacetonate) and Ag (acac). Au and Ag precursor and supports was physically mixed for 20min in Ar atmosphere. Subsequently, the catalysts were calcined at 300°C for 4 hours in air. Catalytic activity was measured by using a fixed bed reactor and a standard gas containing 1 vol.% CO in air. The structure of Au-Ag bimetallic catalysts was observed by aberration corrected TEM/STEM (FEI Titan3 G2 60-300). EDS measurement was carried out by high sensitive EDS system, Super-X (Bruker) equipped with 4-silicon drift detectors (SDD).

Figure 1 indicates catalytic activity for CO oxidation of Au/SiO2 and Au-Ag/SiO2 catalysts. It is clearly confirmed that the catalytic activity of Au/SiO2 is improved by addition of Ag. This effect is prominent for inert supports such as SiO2 and Al2O3. Figure 2 shows ADF-STEM image of Au-Ag/SiO2 catalyst. Nanoparticles with the diameter of 2-10nm are well dispersed on the SiO2 support by solid grinding method. Elemental maps by EDS were carried out and both Au and Ag signal was detected from most nanoparticles. This is indicating that the Au-Ag bimetallic nanoparticles are formed by simple mixing of individual organic complexes of Au and Ag by solid grinding method.

References

1. M. Haruta, T. Kobayashi, H. Sano , N.Yamada, Chem. Lett. (1987) 405.

2. T. Fujitani, I. Nakamura, Angew. Chem. Int. Edit. 50 (2011) 10144.

3. A.Q. Wang, J.H. Liu, S.D. Lin, T.S. Lin, C.Y. Mou, J. Catal. 233 (2005) 186.

4. Y. Iizuka, T. Miyamae, T. Miura, M. Okumura, M. Daté, M. Haruta, J. Catal. 262 (2009) 280.

5. T. Ishida, M. Nagaoka, T. Akita, M. Haruta, Chem. Eur. J. 14 (2008) 8456.


The authors are grateful to Ms. F. Arai, Ms. C. Fukada and Ms. M. Makino for helpful work on the preparation of catalysts and measurements of catalytic activity.

Fig. 1: Catalytic activity for CO oxidation of Au/SiO2 and Au-Ag/SiO2 catalysts. SV: 20000mLh-1g-cat.-1

Fig. 2: ADF-STEM image of Au-Ag/SiO2 catalyst.

Type of presentation: Poster

MS-1-P-3091 Structural properties of GaN/AlN/GaN core-double shell nanowires

Koukoula T.1, Kehagias T.1, Kioseoglou J.1, Eftychis S.2, Kruse J.2, Georgakilas A.2, Karakostas T.1, Komninou P.1
1Physics Department, Aristotle University of Thessaloniki, GR-54124 Thessaloniki, Greece, 2Microelectronics Research Group, Physics Department, University of Crete, P.O. Box 2208, GR-71003 Heraklion, and IESL/FORTH, P.O. Box 1385, GR-71110 Heraklion, Crete, Greece
komnhnoy@auth.gr

Core-shell nanowires (NWs) are an ingenious discovery of interfacial nanoengineering that comprises structural characteristics, which cannot be reproduced by any kind of epitaxial growth. This stems from the fact that core-shell NWs tolerate a larger lattice mismatch without interfacial defects than other heterostructures, thus offering a wider area of band-gaps with several potential optoelectronic applications. It is established that wurtzite NWs grown along the polar direction are bounded by the {10-10} m-planes. As a result, core-shell NWs comprise non-polar {10-10} interfaces between the core and the shell.
The structural properties of core-double shell (GaN/AlN/GaN) NWs, grown by plasma-assisted molecular beam epitaxy (PAMBE) on Si(111), were explored by transmission electron microscopy (TEM) methods. GaN NWs were grown on a thin AlN nucleation layer (3 nm) for 3h, with intermediate AlN spacers (10-15 nm thick) deposited at 1h and 2h growth time [Fig. 1(a)]. High-resolution TEM (HRTEM) imaging revealed the core-double shell morphology with an AlN shell of 0.7-1 nm thick, and a GaN shell varying from 1.6 to 2.7 nm [Fig. 1(b)]. Line profiles of HRTEM images along the growth axis showed that this particular configuration imposes the c-lattice constant of the AlN shell to be adapted to the c-lattice constant of the GaN core. Therefore, a full elastic accommodation of the AlN on GaN is established, considering the absence of misfit dislocations (MDs) from the interface.
The strain state of NWs was evaluated by geometrical phase analysis (GPA). A gradual relaxation of the AlN spacers was observed from the GaN/AlN interfaces towards the center of the spacer for the a-lattice parameter as illustrated in Fig. 1(c), without the presence of MDs. The corresponding FFT is shown in Fig. 1(e). The GPA strain map of the AlN spacer/GaN along the growth direction and the corresponding line profile are shown in Figs. 1(d&f). Regarding the AlN shell, the GPA lattice strain along the growth direction was estimated near zero, verifying the HRTEM observations on the full lattice registration. Considering the very small thickness of the shell, the average in-plane lattice constant approximates pseudomorphic growth. This implies that the AlN shell deviates from the biaxial strain state. Moreover, the GaN shell exhibits the relaxed lattice constant values in both in-plane and out-of-plane directions. It seems that the core-shell configuration of the NWs induces strain fields, which may be exploited in band-gap engineering.


This research has been co-financed by the European Union (European Social Fund - ESF) and Greek national funds through the Operational Program "Education and Lifelong Learning" of the National Strategic Reference Framework (NSRF) - Research Funding Program: THALES: Reinforcement of the interdisciplinary and/or inter-institutional research and innovation.

Fig. 1: (a)TEM image illustrating the NWs morphology along with their schematic model. Black arrows denote the AlN spacers (b) HRTEM image depicting the double shell-core configuration (c) In-plane GPA phase image of the AlN/GaN and (e) the corresponding FFT (d)&(f) GPA strain map of the AlN/GaN along the growth direction and the corresponding line profile

Type of presentation: Poster

MS-1-P-3092 Surface dependent structure of GaN nanowires spontaneously grown on Si

Koukoula T.1, Kehagias T.1, Kioseoglou J.1, Eftychis S.2, Kruse J.2, Georgakilas A.2, Komninou P.1
1Physics Department, Aristotle University of Thessaloniki, GR-54124 Thessaloniki, Greece, 2Microelectronics Research Group, IESL, FORTH, P.O. Box 1385, GR-71110 Heraklion, and Physics Department, University of Crete, P.O. Box 2208, GR-71003 Heraklion, Greece
tkouk@auth.gr

Self-assembled GaN nanowires (NWs) were grown by plasma-assisted molecular beam epitaxy (PAMBE) on Si(111) substrates. Treatment of the substrate surface is critical for NWs growth, as well as their morphological features and crystal quality, and hence their optical properties. To this end, a transmission electron microscopy (TEM) study was performed, to compare the spontaneous nucleation of GaN NWs, when they grow on bare Si, with or without nitridation of the surface, and when they grow on top of an AlN nucleation layer (NL) of varying thickness.
Direct GaN growth on bare Si without nitridation resulted in a high density of tilted GaN NWs, grown on a thin amorphous SixNy layer due to the inevitable interaction of the active N species with the Si surface. NW tilting is attributed to the roughness of the SixNy layer, following the roughness of a stepped Si surface. Indeed, high-resolution TEM (HRTEM) images revealed that tilted NWs nucleated on SixNy just over Si surface steps. When the surface was intentionally nitridated, prior to NWs growth, axial alignment of NWs substantially improved, owing to the formation of a thicker SixNy at the GaN/Si interface, which accommodated any Si surface steps (Fig. 1). Moreover, wurtzite GaN crystalline remnants detected on SixNy might have functioned as potential NW nuclei at the onset of NW growth. Besides the axial inclination of GaN NWs from the growth direction, plan-view observations showed, occasionally, a ~3o in-plane rotation between GaN and Si. We constructed the interfacial atomistic models of a GaN NW epitaxially grown on Si, and a NW where the (0001) planes of GaN were rotated about 3o relative to the (111) planes of Si. In both cases, the GaN/Si superlattice unit cell exhibits a hexagonal shape.
In order to optimize the Si surface treatment, an AlN NL with thickness ranging from 2 nm to 20 nm was used, either as-grown or annealed under active N, i.e., nitridated. In contrast to the previous cases, the amorphous SixNy layer was eliminated from the interface allowing improved alignment and crystal quality of the GaN NWs. When using a 2 nm thick AlN nitridated NL, GaN islands appeared along with GaN NWs (Fig. 2). Conversely, a compact faceted GaN layer with sparse GaN NWs was observed over a 20 nm thick AlN NL. The latter also emerged when the 2 nm AlN NL was not nitridated, however in this case a significantly higher density of NWs was observed (Fig. 3). Therefore, at the initial stages of NL growth, AlN forms 3D islands, which during annealing evolve into a compact 2D AlN NL affecting the morphology of the NWs and the GaN faceted domains.


Research co-financed by the European Union (European Social Fund – ESF) and Greek national funds through the Operational Program "Education and Lifelong Learning" of the National Strategic Reference Framework (NSRF) - Research Funding Program: THALES, project “NanoWire”.

Fig. 1: (a) TEM image, near the [11-20]GaN zone axis (z.a.), showing the morphology of GaN NWs grown on nitridated Si. (b) HRTEM image of the GaN/Si interface, along the [ 11-20]GaN z.a., with GaN crystalline remnants (black arrows) on top of the amorphous SixNy layer.

Fig. 2: (a) TEM image, off the [11-20]GaN z.a., depicting the improved alignment of NWs, when they grow on top of 2 nm AlN nitridated NL. (b) HRTEM image of the GaN/AlN/Si interface, along the [11-20]GaN z. a., showing the solid AlN NL.

Fig. 3: (a) TEM image, near the [11-20]GaN z.a., depicting a compact faceted GaN layer when the 2 nm AlN NL was not nitridated. (b) HRTEM image, along the [11-20]GaN z.a., showing the formation of AlN 3D islands.

Type of presentation: Poster

MS-1-P-3101 Analytical characterization of bimetallic gold-ruthenium catalysts supported on ceria zirconia mixed oxides

Chinchilla L. E.1, Olmos C.1, Blanco G.1, Kurttepeli M.1, Bals S.1, Van Tendeloo G.2, Villa A.2, Prati L.2, Calvino J. J.3, Chen X.1, Hungría A. B.1
1Dpto. de C.M.I.M.Q.I.,Universidad de Cádiz, Spain, 2EMAT,University of Antwerp, Belgium, 3Dipartimento di Chimica,Universita’ degli Studi di Milano, Italy
ana.hungria@uca.es

In the present work, HAADF STEM and XEDS studies were performed in a TITAN 80-50 and a JEOL 2010F microscopes on a series of supported AuRu bimetallic catalysts. In microscopy studies of such complex systems, conclusions are usually drawn about the composition of the particles having analyzed a limited number of them, but rarely estimations are done to know whether the set of chosen particles are representative of the catalyst at a macroscopic level. The results have highlighted the difficulty to characterize the composition of bimetallic systems with particles of varying sizes and metal contents. Despite the technical issues related to STEM mode accompanied with XEDS of very small particles, at least 80 particles of each catalyst were analyzed (Fig 1). All catalysts presented a large fraction of monometallic particles, and a variable amount of particles displaying XEDS signals corresponding to the two metals. XEDS analyses showed that the particles within the small size range were predominantly Ru rich, larger particles were found to be Au and also entities containing both Au and Ru were detected. Based on the models reported by Van Hardeveld et al. [Surf. Sci. 1969, (2), 189] for fcc (Au) and hcp (Ru) crystallites, the precise relationship between the particle size and the total amounts of atoms was calculated. We estimated the number of Au and Ru atoms in each particle present in the particle size-composition diagrams (Fig 1). For 1:2AuRuCZ, instead of the atomic ratio 1Au:2.4Ru measured by ICP, a ratio 1Au:0.25Ru was estimated. The discrepancy may result from the lack of representativeness of microanalysis of individual particles, which caused underestimations of the amount of smaller particles, richer in Ru than the larger (> 5 nm) particles. This hypothesis is supported by the results of XPS analysis showing a relationship 1Au:14Ru corresponding to a presence on the surface of the catalyst of a large number small particles of Ru versus a more aggregated state of Au, given that studies by HAADF-STEM electron tomography have excluded encapsulation phenomena of the metallic phase on the support. An unbiased confirmation of the presence and weight of each particle type can be performed by high resolution XEDS maps (Fig 2), where an aggregate containing a large number of small particles of Ru and an only Au particle of about 30 nm can be seen. In summary, bimetallic AuRu catalysts can be rather non-uniform and can show variation in particle size and composition, due to limited miscibility of the metal components. Catalytic tests showed that bimetallic catalysts were more active than pure Au and Ru catalysts for octanol oxidation, suggesting that Au-Ru interaction, albeit limited, increase the specific activity with respect to the pure components.


Funding from Junta de Andalucía (FQM-3994) and EU FP7 Program (Grant Agreement 312483-ESTEEM2) is gratefully acknowledged. X.C. and A.B.H. thank Ramon y Cajal Program. L.E.C. thank Armand Béché

Fig. 1: Particle size-composition diagrams and relative contribution by particles at a given metal composition for AuRuCZ catalysts

Fig. 2: HAADF-STEM image and XEDS elemental distribution maps (Ce-L, Ru-K, Zr-L and Au-L) recorded on the 1:2 AuRuCZ catalysts. Also is presented an overlay map of the Au, Ru and Ce chemical distribution

Type of presentation: Poster

MS-1-P-3116 Preparation and characterization of Zn (II) complex, [Zn(bipy)2(C6H5)2CHCO2)](ClO4)(bipy), -SLNs formulation

Dikmen G.1, Kani I.2
1Eskişehir Osmangazi University of Eskişehir, Central Research Laboratory, 2Anadolu University of Eskişehir, Department of Chemistry, Eskişehir, Turkey
gokhandikmen@anadolu.edu.tr

Metal organic frameworks offer diverse chemistry as metal-medicine. In these complexes, the metal serves to coordinate the organic ligands. The direct use of metal complexes sometimes is restricted due to lethal side effects. To overcome their disadvantages, solid lipid nanoparticles (SLNs) have been introduced as an alternative drug delivery systems. They carry anticancer compounds with different physiochemical characteristics, higher drug stability, improved pharmacokinetics and controlled drug release. In this study, we synthesized bimetallic Zn(II) complex, [Zn(bipy)2(C6H5)2CHCO2)](ClO4)(bipy), with the reaction of bifunctional 2,2′-bipyridine (bipy) and diphenyl acetic acid (C6H5)2CHCO2H). SLNs formulation prepared by hot homogenization methods and characterized by Zeta Sizer, NMR (Nuclear Magnetic Resonance) and SEM (Scanning Electron Microscopy). In conclusion, SLNs of Zn(II) complex have good stability at -16.5 mV and average particle size around 230 nm. We determined chemical structure as 3D by using XRD. In addition, NMR spectras were carried out tween 80, complex and complex loaded SLN formulations and these spectra compared with each other. According to NMR spectra, both difference in chemical shifts and new peaks were not observed for complex loaded SLN and plasebo SLN. Moreover,  the particle size of Zn (II) complex-SLN formulations was also supported by using SEM. In general, the Zn (II) complex-SLN formulations were spherical shape and uniform in particle size. 

Keywords: Zn, Solid lipid nanoparticles (SLN), Scanning Electron Microscopy (SEM), XRD.


Fig. 1: Figure 1. Sem photo of Zn(II) complex-SLNs formulation.

Type of presentation: Poster

MS-1-P-3118 Characterization and spectral measurement of light emission from individual Au nanoparticles using scanning tunnelling microscopy

Nepijko S. A.1, Chernenkaya A.1, Medjanik K.1, Chernov S. V.1, Yarmak A. V.2, Odnodvorets L. V.2, Schulze W.3, Schönhense G.1
1Institute of Physics, University of Mainz, Staudingerweg 7, 55128 Mainz, Germany, 2Sumy State University, Rimsky-Korsakov Str. 2, 40007 Sumy, C.I.S./Ukraine, 3Fritz Haber Institute of the Max Planck Society, Faradayweg 4-6, 14195 Berlin, Germany
nepijko@uni-mainz.de

The light emission spectra of individual Au nanoparticles induced by a scanning tunnelling microscope (STM) have been investigated. At the same time the nanoparticles were characterized by STM measurements. Two-dimensional ensembles of tunnel-coupled Au nanoparticles were prepared by thermal evaporation onto a native oxide silicon wafer in ultrahigh vacuum (10-10 mbar). We present the experimental evidence of photon emission from Au nanoparticles excited with W tip in the tunnel (tip voltage lower than 5 V) and field emission modes (tip voltage higher than 5 V). In the first case the tunnel current has inelastic component that is used for the tip-induced plasmon excitation. The photon emission that corresponds to them is characterized by a maximum at 1.62 eV. In the second case the photon emission spectrum is more complicated. The photon emission spectrum for Au nanoparticle obtained after subtraction of photon emission from the substrate (the native oxide silicon wafer) is characterized by peaks at 2.22 and 1.45 eV connected with the Mie plasmon and the density of unoccupied states above the Fermi level, relatively. The low-energy peak at 1.45 eV has not been discussed in literature. It was more pronounced then in other publications most likely due to more blunt W tip in our experiment and consequently larger applied voltage (the Au nanoparticle size was a few nanometers in all cases). The use of an STM in the field emission mode with the light signal detection allows implementing of low-energy electron-photon spectroscopy (inverse photoemission spectroscopy).


Type of presentation: Poster

MS-1-P-3142 A complete TEM study of microstructural changes within bifunctional refining catalysts at different stages of their preparation

Moldovan S.1, Grillet N.2, 1, Florea I.3, 1, Danilina N.4, Minoux D.4, Ersen O.1
1Institut de Physique et Chimie des Materiaux de Strasbourg, UMR 7504 CNRS, 23 rue du Loess, 67034 Strasbourg, France, 2Institut Materiaux Microelectronique Nanosciences de Provence, UMR 7334 CNRS, Campus de St Jérôme, 13397 Marseille, France, 3Ecole Polytechnique, Route de Saclay, 91128 Palaiseau, France, 4Total Petrochemicals Research Feluy, Zoning Industriel, Zone C, 7181 Feluy, Belgium
simona.moldovan@ipcms.unistra.fr

Owing to their exceptional properties the bifunctional catalysts are widely employed for processing of heavy oils and are generally composed of an active cracking matrix, a binder and an active (de)hydrogenation function. The actual study focuses on the morphological evolution of bifunctional refining catalysts consisting of a zeolite, an alumina binder and metal sulfides. Three general steps schematized in Figure 1 are considered for the preparation of the catalysts: the extrusion of the zeolite by using alumina as binder, impregnation with Ni and Mo salts and sulfidation to obtain the active metal sulfide phase. The morphology of the zeolite grains is originally marked by spherical mesopores and channels, as well as micropores. The extrusion with alumina does not change fundamentally the zeolite 3D porous structure, but fixes the alumina grains on the zeolite crystals surfaces. The needle-like shaped geometry of the alumina grains contribute to the built-up of a novel porous structure on the zeolite grain surface, without a considerable modification of the accessibility to the zeolite porous structure. The impregnation with Ni and Mo salts leads to a new nanometric architecture, such that one identifies large and small nanoparticles (NPs) within both the zeolite grain and the embedding alumina matrix. A combined analysis evidenced that on the zeolite grains, the small NPs (mean size: 5 nm) are Ni-dominated, whilst the bigger NPs are Mo-rich. One cannot exclude that Mo can appear as small particles and/or as atomic clusters. The zeolite grain morphology and pores size delimit the NPs location: the small NPs will access the micropores, whereas the big ones will be placed exclusively on the mesopores and the channels rims. Though the Mo was predominantly found in the large particles, one cannot exclude that Mo-rich NPs, atoms and/or atomic clusters would penetrate the zeolite grain anchored to the micropores rims. Most of particles deposited on the alumina grains are in Mo, but small amounts of Ni have been also identified. The sulfidation treatment is expected to induce at a structural level the formation of metal sulfide slabs. Indeed, they have been identified mostly on the alumina grains, but also within the zeolite pores inner rims. These finding backups our observations on the presence of Mo NPs in the zeolite pore after impregnation. The electron tomography stressed the slabs location as against the zeolite pores: the metal sulfide slabs do not locate only on the inner mesopores, but also decorate the micropores rims. New 3D core-shell structures were identified in the inner zeolite mesopores: the core is most probably the Ni sulfide, whereas the shell is constituted by MoS2 slabs disposed around the core but not always in contact with it.


Fig. 1: Figure 1. General steps followed for preparation of bifunctional refining catalysts.

Fig. 2: Figure 2. Micrographs acquired under different modes showing the evolution of the zeolite with the preparation steps: A. Slice and model of the zeolite; B. EFTEM image and slice of the extruded zeolite; C. Slices of the TEM and STEM volumes of the impregnated specimen; D. STEM-HAADF micrograph and section corresponding to sulfided specimen.

Type of presentation: Poster

MS-1-P-3159 Wurtzite ZnO/Zinc Blende ZnS Coaxial Heterojunctions and Hollow Zinc Blende ZnS Nanotubes: Synthesis, Structural Characterization and Optical Property

Huang X.1,2, Willinger M. G.2, Fan H.1, Xie Z. L.2, Wang L.1, Hoffmann A. K.2, Lee C. S.3, Meng X. M.1
1Key Laboratory of Photochemical Conversion and Optoelectronic Materials, Technical Institute of Physics and Chemistry, Chinese Academy of Sciences, Beijing, 100190, P. R. China, 2Department of Inorganic Chemistry, Fritz Haber Institute of the Max Planck Society, Faradayweg 4-6, 14195 Berlin, Germany, 3Center of Super-Diamond and Advanced Films (COSDAF) & Department of Physics and Materials Science, City University of Hong Kong, Hong Kong SAR, P. R. China
xinghuang@fhi-berlin.mpg.de

Synthesis of ZnO/ZnS heterostructures in thermodynamic conditions generally results in the wurtzite (WZ) structure of the ZnS component because its WZ phase is thermodynamically more stable than its zinc blende (ZB) phase. In this report, we demonstrate for the first time the preparation of ZnO/ZnS coaxial nanocables composed of single crystalline ZB structured ZnS epitaxially grown on the WZ ZnO nanorod via a two-step thermal evaporation method. The deposition temperature is believed to play a crucial role in determining the crystalline phase of ZnS. Through a systematical structural analysis, the ZnO core and the ZnS shell are found to have an orientation relation of (0002) WZ ZnO//(002) ZB ZnS and [01-10] WZ ZnO//[2-20] ZB ZnS. Observation of the coaxial nanocables in cross-section reveals the formation of voids between the ZnO core and ZnS shell during the coating process, which is probably associated with the nanoscale Kirkendall effect known to result in porosity. Furthermore, by immersing the ZnO/ZnS nanocable heterojunctions in an acetic acid solution to etch away the inner ZnO cores, hexagonal shaped ZnS nanotubes orientated along the [001] direction of ZB structure were also achieved for the first time. Finally, the optical property of the hollow ZnS tubes was investigated. It was found that the tubes can give a strong green emission which may originate from some self-activated centers, vacancy states, interstitial states or structural defects. However, for those tubes with residual ZnO located on tops, they showed much lower emission intensity due to the type-II band alignment of ZnO/ZnS heterojunction that can efficiently decrease the recombination of the electron-hole pairs in both ZnO and ZnS. Our study gives some insights on the controlled fabrication of 1D semiconductors with desired morphology, structure and composition and the as-synthesized WZ ZnO/ZB ZnS coaxial nanocables and ZB ZnS nanotubes provide ideal candidates for the study of optoelectronics of II-VI semiconductors at the nanoscale.


We thank the financial support from the “Strategic Priority Research Program" of the Chinese Academy of Sciences (XDA09040203).

Fig. 1:  (a) TEM image of the nanocables’ cross-sections; (b) HRTEM image and (c) reconstructed structure from (b), cyan: ZnO, red: ZnS; (d) FFT of (b); (e-f) Enlarged HRTEM image recorded from regions of i and ii of (b); (g) HAADF image as well as elemental mapping ; (h-i) EDX data corresponding to spots A and B, respectively.

Type of presentation: Poster

MS-1-P-3199 Self-assembled catalyst promotion by overgrowth of layered ZnO in industrial Cu/ZnO/Al2O3 catalysts

Lunkenbein T.1, Schumann J.1, Behrens M.1, Schlögl R.1, Willinger M. G.1
1Fritz-Haber-Institute of the Max-Planck-Society
lunkenbein@fhi-berlin.mpg.de

Methanol is one of the most important petrochemical molecules. It is considered as a prospective sustainable synthetic fuel obtained by the catalytic hydrogenation of CO2.[1] Industrial relevant catalysts are mainly composed of Cu (>50 mol%)/ZnO in combination with a structural promoter, such as Al2O3 (<10 %).[2] The presence of ZnO drastically increases the intrinsic activity of Cu-based catalysts. This Cu-ZnO synergy can be explained by the appearance of strong metal support interaction (SMSI) upon reduction in hydrogen.[3] The nature of the SMSI effect is versatile and can be expressed by electronic or morphological changes. In the former case an electron transfer from the support to the metal can occur, whereas for the latter situation a migration of the partially reduced oxide over the metal particle arises.
For model systems this migration has already been proposed in one of the early studies of Cu-ZnO synergy[4] and has recently been identified as metastable graphitic ZnO by IR measurements.[5]
Here we present experimental evidence for the presence of this metastable graphitic ZnO overlayer on Cu nanoparticles in an industrially relevant Cu/ZnO/Al2O3 catalyst for methanol synthesis. Direct structural imaging and elemental mapping in the transmission electron microscope show the formation of a layered ZnO overgrowth during reduction (Fig 1).
The results demonstrate a step further towards a complete understanding of the synergistic effects in Cu-ZnO based catalyst for methanol synthesis.


References

[1] G. A. Olah, Angew. Chem. Int. Ed. 2005, 44, 2636-2639.

[2] M. Behrens, S. Zander, P. Kurr, N. Jacobsen, J. Senker, G. Koch, T. Ressler, R. W. Fischer, R. Schlögl, J. Am. Chem. Soc. 2013, 135, 6061-6068.

[3] a) S. J. Tauster, S. C. Fung, R. T. K. Baker, J. A. Horsley, Science 1981, 211, 1121-1125; b) S. J. Tauster, S. C. Fung, R. L. Garten, J. Am. Chem. Soc. 1978, 100, 170-175.

[4] J. C. Frost, Nature 1988, 334, 577-580.

[5] V. Schott, H. Oberhofer, A. Birkner, M. Xu, Y. Wang, M. Muhler, K. Reuter, C. Wöll, Angew. Chem. Int. Ed. 2013, 52, 11925-11929.


Fig. 1: A) Representative TEM micrograph of the reduced catalyst that shows the graphitic overlayer. B) The corresponding energy filtered TEM map of the O K edge indicates a core-shell structure. C) Scanning TEM image of Cu/ZnO/Al2O3. The inset denotes electron energy loss spectra of the Cu L2,3 and Zn L2,3 edge of one single Cu nanoparticle (see ROI).

Type of presentation: Poster

MS-1-P-3206 In-situ electrical characterization of single ZnO nanowires

Fontana M.1, 2, Chiodoni A.1, Bejtka K.1, Jasmin A.1, 2, Porro S.1, Pirri C. F.1, 2
1Istituto Italiano di Tecnologia, Center for Space Human Robotics, Torino, Italy, 2Dipartimento di Scienza Applicata e Tecnologia, Politecnico di Torino, Torino, Italy
marco.fontana@iit.it

Among the different ZnO nanostructures, nanowires (NWs) are of particular interest for applications: they are envisioned as possible future building blocks for nanoelectronics [1] due to their well-defined geometry and possibly enhanced electronic transport properties, coupled with size-dependent piezoelectric response [2]. It is of great importance to understand how the electronic transport behavior of the NWs is influenced by morphology and crystalline structure on one side, and by the particular approach adopted for the implementation of electrical contacts at the nanoscale on the other side. Electrical measurements on single NWs have been reported in literature, with resistivity values ranging over many orders of magnitude ([3,4]), depending on the synthesis techniques, morphology, crystalline structure and defects, type of electrical nanocontacts, ambient conditions and experimental set-up. Therefore, work still needs to be done in order to gain fundamental understanding of the electrical properties of single ZnO NWs and their relationship with synthesis and structure.
In this work we report on the structural, morphological and electrical characterizations of ZnO NWs by means of Field Emission Scanning Electron Microscopy (FESEM), Transmission Electron Microscopy (TEM) and two-probe I-V measurements performed in-situ by the dual-beam FIB-FESEM system. The ZnO NWs were grown by low-pressure chemical vapor deposition (LPCVD) on Si wafers. The morphology of the samples was initially characterized by FESEM: the aspect ratio and the homogeneity of the cross-section along the whole NWs were evaluated. High-resolution TEM characterization shows that the NWs are wurtzite single crystals and they are oriented along the [001] crystalline direction. In order to perform electrical measurements, the NWs were detached from the Si substrates, ultrasonically dispersed in ethanol and then transferred onto SiO2 substrates for the subsequent preparation of metallic contacts. FIB-induced deposition of Pt pads from a gas precursor (methylcyclopentadienyl-trimethyl platinum) was carried out and two-probe measurements were performed in the dual-beam chamber by direct contact of two micromanipulators on the deposited Pt contacts. In the case of NWs with suitable length (> 4 µm), the deposition of more than two contacts was considered to perform transmission line measurements in order to gain information about the contact resistance.

1. Z.L. Wang, Material Science and Engineering R 64 (2009), 33-71
2. H.D. Espinosa et al., Advanced Materials 24 (2012), 4656-4675
3. E. Schlenker et al., Nanotechnology 19 (2008), 365707
4. Jr-Hau He et al., Nanoscale 4 (2012), 3399


Type of presentation: Poster

MS-1-P-3233 New understanding of the atomic structure and nonstoichiometry of Sc-doped TiO2 photocatalysts by using aberration corrected STEM ARM200CF microscope and EELS spectroscopy

Bakardjieva S.1, Klie R.2, Paulauskas T.2, Philips P.2, Bezdička P.1, Šubrt J.1, Janoš P.3, Gartnerova v.4, Jager A.4
1IIC AS CR v.v.i., 250 68 Rez, Czech Republic, 2UIC Chicago, 845 W Taylor Street, 60607 Chicago, Illinois, USA, 3UJEP, Králova Výšina 3132/7, Ústí nad Labem, 400 96, Czech Republic, 4IP ASCR v.v.i., Na Slovance 1999/2, 182 21 Praha, Czech Republic
snejana@iic.cas.cz

The studies on scandium doped TiO2 photocatalyst nanocrystallites are still limited in the iterature. We select Sc-doped process as an effective method to control the surface structure of TiO2 by generation of defects into TiO2 lattice. Characterization of structures at atomic resolution was performed by aberration-corrected JEM-ARM200CF microscope allows for 68 pm spatial resolution. Z-contrast STEM images were collected using HAADF detector. EELS was acquire atomic-scale maps of the chemical composition and assess the local bonding and Ti valence. This is the first example of a Ti-Sc-O system demonstrated that the substitution of Sc for Ti results in changes in photocatalytic activity due to the preferred occupancy of Sc atoms and its effects on the anatase lattice. The changes of both the lattice parameters and surface morphology are related to the chemical bonding between Ti and Sc cations and oxygen atoms, and species formed at the surface with respect to oxygen deficiencies.The highest activity was observed in TiO2-Sc 4.18 at.%. The photocatalytic activity of the Sc doped TiO2 strongly depends on Sc concentration and particle size of both the dopant ion and TiO2 matrix. A general view of the Sc-doped anatase is shown in Fig. 1. The well-known spherical morphology of TiO2 is established. Detailed atomic scale analysis performed with STEM-HAADF detector shows composites with core-shell structure (Fig.2a). High resolution Z-contrast images demonstrates that the scandium ion dispersed into TiO2 nanoparticle and its concentration in the shell regions is higher and formed scandium-rich regions where Sc3+ replace Ti3+ and mixed oxygen vacancy generated composite (Ti3+xTi4+1-x)O2-x is developed. Therefore, the Sc ion can be caused defects by introducing oxygen vacancies in the lattice of TiO2. Two different crystallographic regions can be recognized based on the intensity and the atomic column symmetry. The border region has much larger intensity. This contrast is explained mainly by the effect of the strain field presented in the TiO2@Sc interface (Fig.2b,c and Fig. 3.). HAADF find out that the lattice parameters of the doped TiO2 samples begin to be larger than pure TiO2. We suppose that during nano-crystal growth, the method by which solutions of reacting components are mixed and the intensity of their stirring can influence the precipitation and the physical characteristics of the product. The precipitation of Sc-doped TiO2 results from mixing of two liquids on microscale level. The mixture consists of entirely segregated parts with different local concentration. In such an way a molecular diffusion occurs to form crystals where the inner TiO2 core has a different composition from the outer shell


The authors thank Project EnviMod UJEP for fanantial support and Dr. Stengl team for synthesis of samples.

Fig. 1: HRTM and HAADF images of Sc doped TiO2

Fig. 2: Atomi scale HAADF images of TiO2@Sc core-shell and interfaces

Fig. 3: ABF images and EELS of grain boundaries

Type of presentation: Poster

MS-1-P-3270 Oxidation study of silver nanoparticles

BAZAN-DIAZ L. S.1, HERRERA-BECERRA R.2
1PCeIM, UNAM, 2INSTITUTO DE FISICA, UNAM
bazanlulu@gmail.com

During the last few years has been a great interest to develop new functional nanomaterials for diverse applications, including areas such as semiconductors, optical devices, photo-electrochemistry, among others [1]. It is well known that the formation of nanoparticles is highly sensitive to different the reaction conditions. Particularly, the control of shape and size of the nanoparticles are parameters of major importance for the different applications where they are used. One of the main physicochemical parameters that control the final size and shape of nanoparticles is the pH [2], resulting in the formation of Ag nanoparticles of different morphologies depending on precursors concentration at different pH values.

Silver nanoparticles (AgNP) have been subject of several studies due to their diverse applications; however, it has been observed that when AgNP are synthesized in aqueous solutions they might suffer oxidation that could derive in loss of activity[3].

In this work, we studied the oxidation of AgNP obtained in aqueous solutions, by employing a so-called "green" method for the production of nanoparticles, with especial focus on the characterization of forms and structures of AgNP. The method employed here is at room temperature, using synthetic tannins as principal reducing reagents and AgNO3 as precursor in aqueous solution. During the reactions, the pH was modified by using NaOH solutions at low concentrations, obtaining different sizes and shapes of AgNP. The mixtures were subjected to ultrasonic treatment, centrifugation and drying. This straightforward method has proven its effectiveness in the reduction of metal nanoparticles on earlier work by our group [4].

It was observed that exist a critical stable size, after it AgNP have tendency to oxidize. Advanced analytical electron microscopy characterization will be employed to determinate the critical size and the final structure, shape, atomic arrangement and chemical composition of the AgNP produced by this method.

References:

[1] L.A.Botello et al., Ingenierías, Octubre – Diciembre 2007, Vol. X, N°. 37
[2] Fang Liao, Zhou Feng, Xing-QiHu, Ionics 17, (2011)81-86.
[3] Chung-Ming Li, I.M. Robertson, M.L. Jenkins, J.L. Hutchison, R.C. Doole, Micron 36 (2005)9-15.
[4] Herrera-Becerra, R. et al; 2010. Appl. Phys. A, 100 (2), 453-459.


The authors acknowledge the financial support from DGAPA with grant PAPPIT IN105112 and from the Graduate Program PCeIM, UNAM.Our gratitude to Roberto Hernández and Cristina Zorrilla for the technical support given.Authors thank access to the facilities of the Kleberg Advanced Microscopy Center and the Research Centers in Minority Institutions (RCMI) at The University of Texas at San Antonio, USA.

Fig. 1:
Type of presentation: Poster

MS-1-P-3275 In-situ Electron beam induced transformation of boehmite to gamma alumina

Ozkaya D.1
1Johnson Matthey Technology Centre, Reading UK
ozkayad@matthey.com

Gamma alumina is widely used as a support material for industrial catalysts such as refining petrochemicals or Fischer-Tropsch processes. Boehmite, Aluminum oxihydroxide (AlOOH), is the precursor to Gamma-alumina (γ-Al2O3) in the manufacturing process[1].The transformation from boehmite to gamma alumina is critical to the understanding of the structure of gamma alumina, thus has been widely studied [2]. The removal of water from the structure creates a metastable gamma phase with a structure that gets classified as tetragonal or cubic depending on the choice of unit cell [2]. The changes to the structure and surface area during this transformation are still not very well understood. One of the problems with studying boehmite using transmission electron microscopy (TEM) is the electron beam sensitivity of boehmite. Boehmite transforms to gamma alumina if observed at ambient temperatures in the TEM even under low beam dose conditions. Here, Cryo-TEM has been used to slow down the transformation under the beam in-situ as the sample de-hydrates and forms gamma alumina. The main aim of the study was to observe changes in the shape and form of the nanocrystalline material after it has transformed to gamma alumina. The properties of gamma alumina is what determines the diffusion rates and sintering behaviour of Platinum group metals (PGM’s)in most of the current auto catalysts. The study was carried out using a Gatan Cryo holder at -184C in a FEI Tecnai F20 Twin lens configuration S/TEM. The boehmite sample was dispersed on a holey carbon coated Cu-grid and put into the cryo-holder and the microscope before cooling. Some ice formation was observed at some parts of the sample as boehmite itself contains some excess water. A sequence of images was acquired as function of time together with diffraction patterns to confirm the structural transformation. A sequence of diffraction patterns diffraction patterns of boehmite and gamma alumina are shown in Figure 1 where the rings formed have been identified using JEMS software. The difference is quite stark in the way the pattern goes as the boehmite is orthorhombic and gamma alumina is pseudo cubic or tetragonal structure. The characterisation of the diffractions patterns shown together with the calculated spacings on the diffraction patterns in figure 1.

Figure 2 shows the sequence of in-situ images from boehmite to gamma alumina, the only driving force for the transformation being the electron beam. Images were acquired in sequence together with the diffraction patterns at regular intervals. The images shows that the shape of the crystallites remain the same after the transformations has taken place. However there is a slight, in shrinkage in the size of the particles around 10-15% after the transformation.


I would like to thank Stephen Spratt from Johnson Matthey Technology Centre for help during cryo experiments

Fig. 1: Figure 1. Series of diffraction patterns at different time intervals showing the transformation from boehmite to gamma alumina under the electron beam. The calculated diffraction patterns from boehmite and gamma alumina are shown for comparison

Fig. 2: Figure 2. Series of in-situ images taken at various time intervals from the same area that diffraction patterns were taken in figure. Transformation from boehmite to gamma alumina transformation does not cause drastic change in the crystal shape but a 10-15% reduction in size.

Type of presentation: Poster

MS-1-P-3278 TEM Characterization of Core/Shell (FexOy/Au) nanoparticles

Mendoza Cruz R.1, Herrera Becerra R.2
1PCeIM-Insititute of Physics, UNAM, 2Institute of Physics, UNAM.
rumec21@gmail.com

In the last years it has been much interest on the synthesis of bimetallic or core/shell nanoparticles since these systems present novel properties different from the monometallic particles of the same elements. Even more, the properties can be tuned by varying the way of mixing the elements. In systems where a magnetic core as iron oxide is covered by a noble metal shell (Ag, Au or Pt), the magnetic properties and the efficient adsorption of molecules and optical properties of the shell can be exploited for numerous applications in fields as biomedicine, biotechnology or catalysis. Also, a metallic shell can modify the surface of the magnetic cores and provide them of extra stability against aggregation or oxidation. However, to achieve novel characteristics, control of the size and thickness of the shell must be reached, since all of these could generate loss of their intrinsic properties.

Many reports on the synthesis of iron oxide core and gold shell can be found and several methods have been reported to control the nanoparticle size and the phase. Nevertheless, few works have been developed on the structure and shape of the magnetic seeds, and the final structure and morphology of this core/shell system. The study of the chemical composition, internal structure, size and morphology is of great interest given that the properties of these systems depend on these parameters. For instance, the interphase effect between the two materials can yield to an enhancement of the catalytic activity or in the magnetic moment of the total particle [1,2].

In this work, a tannin-assisted method for the synthesis of core/shell nanoparticles (iron oxide/gold) is carried out. This method has been effective for the production of iron oxide particles with a narrow size distribution [3]. The final structure and morphology are determined by high resolution transmission electron microscopy. This work attempt to contribute to the understanding of all these factors have on the novel properties of oxide core/metal shell nanoparticles.

References:

[1] L. Guczi, et al. Topics Catal. (2006) Vol. 39, Nos. 3-4, 137-142.
[2] S. Banerjee, et al. J. Appl. Phys. (2011) 109, 123902.
[3] Herrera-Becerra, R et al. Appl Phys A (2010) 100, 453-459.


The authors acknowledge the financial support from DGAPA with grant PAPPIT IN105112 and from the Graduate Program PCeIM, UNAM.
Our gratitude to Roberto Hernández and Cristina Zorrilla for the technical support given.
Authors thank access to the facilities of the Kleberg Advanced Microscopy Center and the Research Centers in Minority Institutions (RCMI) at The University of Texas at San Antonio, USA.

Type of presentation: Poster

MS-1-P-3293 The InAs Islands Morphology and Structure Dependence With the FIB Ion Dose Used to Induce Its Localized Growth.

Ribeiro Andrade R.1,6, Miquita D. R.1, Vasconcelos T. L.2, Kawabata R.4,6, Malachias A.5, Pires M. P.3,6, Souza P. L.4,6, Rodrigues W. N.5,6
1Centro de Microscopia, UFMG, Belo Horizonte, MG, Brazil, 2Divisão de Metrologia de Materiais, INMETRO, Duque de Caxias, Brazil, 3Instituto de Física, UFRJ,Rio de Janeiro, Brazil, 4LabSem/CETUC, PUC-Rio, Rio de Janeiro, Brazil, 5Departamento de Física, ICEx, UFMG, Belo Horizonte, MG, Brazil, 6DISSE – Instituto Nacional de Ciência e Tecnologia de Nanodispositivos Semicondutores, CNPq/MCT, Brazil
rodriban@gmail.com

Semiconductor quantum dots have optoelectronic properties that attract and justify the efforts of many research teams around the world. However the production of high efficiency optoelectronic devices depends on a fine control of the morphology, density, size, shape, uniformity and chemical composition of these nanostructures. Several techniques to induce a localized growth of QDs can be combined with modern growth techniques in an attempt to control the precursor’s nucleation on the substrate surface. One of techniques employed in this process is the Focused Ion Beam Microscopy (FIB) [1,2]. In this technique a gallium ion beam is used to create localized surface defects that become nucleation sites for the localized growth of quantum dots. In this work we used this technique to create an array of holes on InP(001) surface to serve as diffusion barrier increasing the nucleation rate located during InAs grown by Metal Organic Vapor Epitaxy. The doses used were 3.7x1015, 5.6x1015 and 1.3x1016 Ga+/cm2. Islands were grown for two sub-monolayer coverages, occurring mostly in clusters in the inner surfaces of the FIB produced cavities. For low doses templates the nanostructures are mainly coherent. For high doses the islands are mostly incoherent and numerous. A simple model correlating the surface potential of the template with the net adatom flow to the cavities is presented. Two regimes were identified, coarsening and coalescence when low doses were applied, and incoherent growth when high doses were used.


This work was supported by Programa Institutos Nacionais de Ciência e Tecnologia, CNPq/MCT, CAPES, FAPEMIG and FAPERJ. We would like to thank the UFMG Microscopy Center as well the Brazilian National Light Synchrotron Laboratory for the technical and financial support.

Fig. 1: Figure (a) shows a schematic diagram of the adatoms diffusion process into the cavities and islands growth. (b) Analysis of islands density on an InP modified substrate as a function of ion dose, indicating the presence of two distinct growth regimes.

Type of presentation: Poster

MS-1-P-3304 Long-range chemical orders in gold-palladium nanoalloys studied by aberration-corrected TEM

Nelayah J.1, Nguyen N. T.1, Alloyeau D.1, Wang G.1, Ricolleau C.1
1 Laboratoire Matériaux et Phénomènes Quantiques, Université Paris Diderot/CNRS, UMR 7162, Bâtiment Condorcet, 4 rue Elsa Morante, 75205 Paris Cedex 13, France
jaysen.nelayah@univ-paris-diderot.fr

Gold-palladium nanoparticles (Au-Pd NPs) are of great industrial and scientific interests as they are promising catalysts for many oxidation and hydrogenation reactions. Catalytic activity of Au-Pd NPs can be fine-tuned by controlling the ratio and arrangement of the surface atoms at their surface. As in the bulk phase, Au and Pd atoms can be mixed in any proportion in nanoalloys to form chemically-disordered face-centered cubic (fcc) structures. Aside from the fcc structures, electron diffraction studies of epitaxially-grown Au-Pd thin films have also hinted at the existence of long-range chemical orders of L10 type (around composition AuPd) and of L12 types (around compositions Au3Pd and AuPd3). However, no long-range chemical-order has ever been observed in Au-Pd NPs. In this contribution, we report the first observation of long-range chemical orders in Au-Pd NPs using an aberration-corrected transmission electron microscope (TEM).

Monodispersed Au-Pd NPs under 10 nm were grown by pulsed laser deposition in vacuum on either freshly cleaved NaCl(001) single crystals or amorphous carbon films of standard TEM grids. Material transfer to the support and precise control of particle composition was achieved by alternately ablating two ultra pure Au and Pd targets. The NPs were annealed in vacuum at temperatures superior to 400°C. Epitaxially-grown NPs were transferred to standard TEM grids prior to annealing. The structure of the as-grown and annealed NPs was studied by high-resolution imaging on a JEM-ARM 200F TEM.

Fig. 1 shows a corrected TEM image (UHRTEM) of an epitaxially-grown Au-Pd NP in closed [001] orientation. Due to the rapid kinetics of the deposition process, the structure of all as-grown NPs was in a non-equilibrium state disordered fcc type. Transition to the equilibrium structure occurred during high temperature annealing. Extensive UHRTEM studies of Au-Pd NPs of various compositions annealed in different conditions (support, temperature, time) showed that both L10 and L12 chemical order are stable in the Au-rich NPs at temperatures as high as 600°C. UHRTEM images of a L10- and a L12- chemically-ordered NPs in closed [001] orientation are shown in Fig. 2 and 3 respectively. Besides chemical ordering, particle ripening also occurred during thermal annealing. Single particles X-ray analysis in TEM scanning mode demonstrates clearly that crystal growth proceeds through Ostwald ripening phenomenon. Due to the difference in evaporation rate of Au and Pd atoms from the NPs during this ripening process, the chemical composition of the NPs became size-dependent after annealing processes with a systematic enrichment in Pd as particle size increases.


The authors acknowledge the support of the French Agence Nationale de la Recherche (ANR) under reference ANR-11-BS10-009. We are also grateful to Région Ile-de-France for convention SESAME E1845 for the support of the JEOL ARM 200F electron microscope installed at Paris Diderot University.

Fig. 1: UHRTEM image of an as-grown Au-Pd NP composed of 62 % at Au viewed close to the [001] direction. The fast fourier transform of the image is displayed in insert. It shows that the structure of the NP is of disordered fcc type.

Fig. 2: UHRTEM image of a L1o chemically-ordered NP obtained after annealing the sample in Fig 1 at 600°C in vacuum during 10h. Chemical ordered is evidenced through the appearance of superlattice structures in the fast-fourier transform of the image.

Fig. 3: Image of a L12 chemically-ordered NP obtained after annealing at 500°C in vacuum during 8h a sample with composition comparable that of Fig 1.

Type of presentation: Poster

MS-1-P-3307 Ex-situ and In-situ Analysis of MoVTeNb Oxide by An Aberration-Corrected Scanning Transmission Electron Microscope

Xu P.1, Sanchez-Sanchez M.2, Browning N. D.3, Lercher J. A.2,3
1Department of Chemical Engineering and Materials Science, University of California, Davis, Davis, USA, 2Department of Chemistry and Catalysis Research Center, Technische Universitat Munchen, 85748 Munich, Germany, 3Physical Sciences Laboratory, Pacific Northwest National Laboratory, 902 Battelle Boulevard, Richland, WA 99352, USA
amy.pinghongxu@gmail.com

Short chain olefins, especially ethylene and propylene are in very high demand worldwide making them important industrial raw materials. Two-phase MoVTeNb oxide, has shown great promise in oxidative dehydrogenation (ODH) of converting inexpensive light alkanes into these olefins. In this work, we report an identification of the nature of crystalline termination of the catalytically relevant M1 phase at atomic scale using an aberration-corrected scanning transmission electron microscope (STEM). Figure 1 shows a typical M1 phase particle viewed in the crystal growth direction (the [001] orientation). It is clearly observed that at standard conditions no amorphous overlayer is formed on the external surface of the particles. Based on the analysis of over 50 particles, it is shown that the lateral surfaces of these rods are faceted and the most preferential lateral facets have been determined to be {010}, {120} and {210} (as shown in Figure 1 B-D for STEM images and Figure 1 E-G for crystallographic models). These results indicate that morphology of the M1 phase particles might have a large impact on catalytic activity of this system. Additionally, electron energy loss spectroscopy and energy dispersive X-ray spectroscopy has been applied to determine the vanadium distribution on the surface, which is believed to play an important role in catalytic performance of the M1 phase in ODH.

Reaction conditions affect the surfaces of the M1 phase, which makes it essential to perform the in situ experimental observations for a fundamental understanding of the catalytic performance of this material. Direct imaging of structural changes in the M1 phase was performed under an oxidative atmosphere in an environmental TEM. Tellurium units were observed to disappear from hexagonal channels when heated to 350 °C in 10 mbar oxygen/argon (23%/77%), while the crystal framework remained unaffected. Further in-situ experiments using a much higher partial pressure (above 1 bar) are underway to investigate the relationship of structural changes to catalytic performance of this system under realistic reaction conditions.


This work was supported in part by the United States Department of Energy (DOE) Grant No. DE-3-BDOE797 through the University of California, Davis, the Laboratory Directed Research and Development Program (LDRD): Chemical Imaging Initiative at Pacific Northwest National Laboratory (PNNL), and the Environmental Molecular Sciences Laboratory (EMSL), a national scientific user facility at PNNL.

Fig. 1: (A) Unprocessed STEM image showing crystalline and faceted surface of M1 phase. (B-D) Magnified view of rectangular areas 1-3 in (A), showing different configurations of facets {010}, {120} and {210}, respectively. (E-G) Rendering of the crystallographic model of the three facets shown in (B-D).

Type of presentation: Poster

MS-1-P-3369 Transformation of CdSe/Cu3P/CdSe Nanocrystal Heterostructures upon Thermal Annealing

Falqui A.1,2, Genovese A.1, Casu A.1, De Trizio L.1, Manna L.1
1Nanochemistry, Istituto Italiano di Tecnologia, Via Morego 30, 16163 Genova, Italy, 2Biological and Environmental Sciences and Engineering Division, King Abdullah University of Science and Technology (KAUST) – Thuwal 23955-6900, Kingdom of Saudi Arabia
andrea.falqui@gmail.com

Thermal annealing in conjunction with electron irradiation can also be exploited to create new nanostructures. One key point of nanocrystal (NC) heterostructures tranformation is that the relative thermal or irradiation stability/reactivity of a certain material domain in the heterostructure can be much different from that of the same domain if it were the only constituent of the NC. This is due to its proximity to domains of other materials that can elicit concomitant processes such as diffusion of chemicals, alloying, charge transfer and others. We have studied the in-situ phase transformation of CdSe/Cu3P/CdSe NCs. Starting from hexagonal Cu3P nano-platelets, and working at higher temperatures, we have obtained sandwich-shaped nanoparticles consisting of one top and one bottom layer of CdSe encasing each original Cu3P platelets according an epitaxial-relationship (Figure 1 a-d). When these sandwich-like heterostructures were annealed under high vacuum (pressure »10-4 Pa) up to 450°C, sublimation of P and Cd species with concomitant interdiffusion of Cu and Se species were observed by in-situ high resolution TEM (HRTEM) and energy filtered TEM (EFTEM) analyses. These processes transformed the pristine sandwiches, triggering the complete evolution of each original heterostructure into single NCs that were thoroughly characterized by fcc lattice, with d-spacings compatible with those of fcc Cu2Se (Figure 1 e-f). Under the same conditions the single domains, i.e the pristine (uncoated) Cu3P platelets and isolated CdSe NCs were stable, no transformation occurred. Therefore, the thermal instability of these heterostructures under high vacuum might be explained by the preferential diffusion of Cu species from Cu3P cores into CdSe domains through epitaxial related interfaces. The Cu interdiffusion triggered sublimation of Cd, as well as out-diffusion of P species and the partial dissolution of NCs together with the overall transformation of the sandwiches into Cu2Se single NCs (Figure 2). Therefore, one further development in this direction could be that of creating plasmonic micro- and nanostructures “on demand”, for example, by annealing with a laser individual or groups of CdSe/Cu3P/CdSe NCs deposited on a substrate [1].

[1]. De Trizio et al. ACS Nano, 5, (2013), 3997.


The authors acknowledge financial support from European Union through the FP7 starting ERC grant NANO-ARCH (contract no. 240111). We thank Marijn van Huis and Anil Yalcin for many useful discussions.

Fig. 1: a) CdSe/Cu3P/CdSe view. b) HAADF STEM image and EDX line profile c) HRTEM image of epitaxial interfaces. d) FFT patterns of CdSe (blue) and Cu3P (red) showing lattice similarities; habit schema. e) HRTEM (RT @ 300°C) of single NC showing CdSe/Cu3P/CdSe structure. f) HRTEM (400°C) of the same NC completely converted into fcc Cu2Se.

Fig. 2: Elastic filtered zero loss images (a-c) and the corresponding EFTEM elemental maps (d-f) of several NCs observed at room temperature (RT), 400° and 450°C displaying Cu diffusion. Elemental maps of P (132 eV, green,) and Cu (931 eV, red). Zero loss image at 450°C shows sublimation morphologies. Scale bars 50 nm.

Type of presentation: Poster

MS-1-P-3378 Formation of NiSi2 nanoparticles epitaxially embedded in silicon nanowires

Panciera F.1,2, Chou Y. C.2,3,4, Reuter M. C.2, Hofmann S.1, Ross F. M.2
1Department of Engineering, University of Cambridge, Cambridge, UK, 2IBM Research Division, T. J. Watson Research Center, Yorktown Heights, NY, USA, 3Department of Electrophysics, National Chiao Tung University, Hsinchu city, Taiwan , 4Center for Functional Nanomaterials, Brookhaven National Laboratory, Upton, NY, USA
fp301@cam.ac.uk

Silicon nanowires (SiNWs) have a variety of applications including field-effect transistors (FETs), energy storage, and chemical and biological sensors. These applications generally require electrical contact to be made with the nanowires. For this reason, metal silicides formed by solid state reactions with Si nanowires have attracted great attention in the last few years. However, silicides also have prospects in applications such as spintronics, sensors and energy storage devices at the nano-scale.

We developed a novel method to synthetize silicon/silicide heterostructures by introducing silicide nanoparticles inside silicon nanowires. We found that the Au droplets, which originally catalysed the formation of the SiNWs themselves, can also act as a catalyst for the Ni-Si reaction. This catalytic reaction leads to immediate formation of NiSi2, even at low temperature. To demonstrate this process, we first synthesise SiNWs with a radius between 20 and 50 nm and the usual growth direction, (111). Synthesis took place in an ultrahigh vacuum transmission electron microscope using the vapour-liquid-solid method with Au as the catalyst and disilane as the precursor gas. After the wires were formed, a Ni layer was then deposited by electron beam evaporation at room temperature, without breaking the vacuum. During annealing, as shown in Figure 1, an octahedral silicide particle formed within the liquid AuSi droplet. Measurements of crystal structure allow this to be identified as the NiSi2 phase. This silicide nanoparticle first formed and moved around in the liquid droplet and only later on became attached to the AuSi/Si(111) interface, forming an epitaxial contact (Figure 1a). Once the NW growth was restarted by flowing the disilane precursor gas, the silicide particle became incorporated into the NW forming a Si-silicide heterostructure in which a nanoscale silicide region is embedded epitaxially within the nanowire (Figure 1b, d). By repeating the process, we can embed multiple silicide particles within a single nanowire. We will present video-rate imaging allowing measurements of the nucleation, growth and incorporation of the silicide (Figure 1c) and discuss the mechanism and kinetics. We will then show that the silicide growth mechanism we have demonstrated here for Ni deposited on SiNW applies to other systems as well and can be exploited to introduce nanosized particles in various kinds of nanowires. We will finally discuss the potential applications of these new heterostructures for future nanowire-based devices.


Fig. 1: (a) TEM images recorded at 400°C after deposition of 1nm of Ni onto a Si nanowire. (b) TEM images recorded at 500°C and 1×10-5 Torr Si2H6. The incorporation of the silicide inside the Si matrix is visible. (c) The plot of length vs. time for the nanowire in (b). (d)  A higher magnification image of a NiSi2 particle showing its octahedral shape.

Type of presentation: Poster

MS-1-P-3404 Lanthanum Nanoparticles: Synthesis and Structural Properties

Schabes-Retchkiman P. S.1, Romero-Ibarrra Josue E.1
1Instituto de Fisica, UNAM, Mexico, DF, Mexico
pabloschabes@yahoo.com.mx

In the past decade, a lot of attention has been given for the development of novel strategies for the synthesis of different kinds of nano-objects. Most of the current strategies usually use physical or chemical principles to develop nano-objects with multiple applications. One such system is the Lanthanum (La) system, with notorious catalytic properties.
Lanthanum has diverse properties derived from the structure of the atom: óptical, magnétic, catalytic, superconductors, etc. (in particular, new nanometric scale properties; Reactivity, cathodos in SOFC, etc.) These type of metals are usually synthesized as oxides (La2O3) In this work we have synthesized metallic Lanthanum (Lazero) and La2O3,by bio-reduction with plants.

In these work we show that through a bioreduction method [1] we have been able to produce small Lanthanum nanoparticles, (2-5 nm) and La nanorods, using two different reductor-plants i.e. alfalafa and Callistemon citrinus,  which is interesting because it constitutes a pest. In essence the reaction that seems to take place in both plants is the following acting the water-soluble tannins remaining from the plant treatment:
 

The controlling parameter is the pH of the solution in which the reactions take place [1].

Characterization: Electron microscopy characterization included transmission electron microscopy (TEM), high resolution transmission electron microscopy (HRTEM) characterization, EELS (electron energy-loss spectroscopy) HAADF (High Angle-Annular Dark Field). The samples were prepared by placing a drop of the solution on carbon-coated copper grids and allowed to dry. Electron microscopy was performed in a JEOL JEM-2010F FasTem Microscope at IFUNAM, equipped with EDS and EELS analysis. High resolution images were obtained under many different conditions and the images analyzed by obtaining digital diffractograms by FFT (Fast Fourier Transform)

The synthesized La0 nanoparticles with average diameter of 3 nm can be observed in figure 1. The formation of the nanophases can be appreciated in the HAADF images, particles corresponded to metallic Lanthanum. Some of the particles were oxides as well, some clustered. Rods were also observed. The main difference between the two synthesis was that with the callistemon the particle density was smaller than for alfalfa, which might be due to the breaking of cell walls having very different energies.


Conclusion: The pH of the solution constitutes a parameter of easy control, that defines the size and morphology of the LaN (particles or rods or wires) .

For pH 13, nanorods of metallic character were obtained.

As a final thought, we have to remark that these synthesis methods are simple and self-sustainable.


We are indebted to the personnel at LACMIF, IFUNAM, particularly Mr Roberto Hernandez.

Fig. 1: HAADF of Lanthanum nanoparticles pH10

Fig. 2: Typical morphologies observed for pH10.

Fig. 3: Metallic La particles

Fig. 4: Metallic La, Larger particle.

Type of presentation: Poster

MS-1-P-3412 Electron Beam Nanosculpting and Controlled Morphological Transformation of Kirkendall Oxide Nanochannels

Molina-Luna L.1, Abdel-Aziz A. A.2, 4, Buffière M.3, Tessier P. Y.4, Du K.5, Choi C. H.5, Kleebe H. J.1, Konstantinidis S.2, Bittencourt C.2, Snyders R.2, 6
1Department of Material- and Geosciences, Technische Universität Darmstadt, Germany, 2Chimie des Interactions Plasma-Surface (ChIPS), CIRMAP, Research Institute for Materials Science and Engineering, University of Mons, Belgium, 3imec, Heverlee,Belgium, 4Institut des Matériaux Jean Rouxel, Université de Nantes, France, 5Department of Mechanical Engineering, Stevens Institute of Technology, Hoboken, New Jersey, United States, 6Materia Nova Research Center, Mons, Belgium
molina@geo.tu-darmstadt.de

The nanomanipulation of metal nanoparticles inside oxide nano-tubes, synthesized by means of the Kirkendall effect, is demonstrated. In this strategy, a focused electron beam, extracted from a transmission electron microscope source, is used to site-selectively heat the oxide material in order to generate and steer a metal ion diffusion flux inside the nanochannels. The metal ion flux generated inside the tube is a consequence of the reduction of the oxide phase occurring upon exposure to the e-beam. We further show that the directional migration of the metal ions inside the nanotubes can be achieved by locally tuning the chemistry and the morphology of the channel at the nanoscale. This allows sculpting organized metal nanoparticles inside the nanotubes with various sizes, shapes, and periodicities. This strategy is based on the control of the thermally activated local diffusion of Cu ions inside the nanotube using an e-beam extracted from a TEM source. The migration of Cu ions was found to be governed by a surface diffusion mechanism occurring on the innerwalls of the nanotube. This nanomanipulation technique is very promising since it enables creating unique nanostructures that, at present, cannot be produced by an alternative classical synthesis route. Aditionally, temperature dependent in-situ TEM experiments were carried out yielding a controlled morphological transformation of the Kirkendall oxide nanochannels.


References

El Mel A-A, Molina-Luna L, Buffière M, Tessier P Y, Du K, Choi C-H, Kleebe H-J, Konstantinidis S, Bittencourt C, Snyders R. Electron Beam Nanosculpting of Kirkendall Oxide Nanochannels. ACS Nano, 2014, 8 (2), pp 1854–1861.


This work was funded in part by the Directorate of Research in Wallonia, under the scope of the ERA-NET MATERA Programme and by the COST Action MP0901. The French Community of Belgium is acknowledged through the “Cold Plasma” project. The TEM´s employed for this work were partially funded by the German Research Foundation (DFG).

Fig. 1: Morphological evolution of an oxide nanotube upon exposure to an electron beam. (a) TEM image of the as-grown oxide nanotube. (b-e) formation of Cu nanoparticles inside the oxide nanotube upon exposure of several regions to the e-beam for different subsequent shots. The time of each shot was 2 s. Scale bar:100 nm.

Type of presentation: Poster

MS-1-P-3414 Novel M1/M2 heterostructure in Mo-V-M-Ta (M = Te or Sb) complex oxide catalyst revealed by aberration corrected HAADF STEM

He Q.1, Woo J.2, Guliants V. V.2, Borisevich A.1
1Materials Science and Technology Division, Oak Ridge National Laboratory, Oak Ridge, Tennessee, USA, 2School of Energy, Environment, Biological and Medical Engineering, University of Cincinnati, Cincinnati, Ohio, USA
heq1@ornl.gov

Direct interpretability and atomic number (Z) sensitivity of HAADF combined with sub-Å resolution makes it the techniques of choice for uncovering atomic-scale underpinnings of materials behavior[1]. For Mo-V-M-Ta (M = Te or Sb) catalysts, HAADF has revealed local cation distribution[2], structure and chemistry of the crystal defects[3], polar domain structure[4], surface[5] and other features that helped to understand their high performance in propane (amm)oxidation[6]. Previous works showed that the M1 phase (Fig 1a) is the active phase, while another phase in the system, namely M2 (Fig 1b), has a distinctive synergistic effect with the M1 by improving the selectivity. Aiming to understand and further exploit this synergistic effect, we studied a series of M2 phase catalysts by HAADF imaging in an aberration corrected STEM.

Our results show that the M2 phase in Mo-V-Te-Ta system, unlike other reported M2 phases[4], has unusual microstructure: the side planes of the M2 phase crystals are decorated with pentagons (e.g. Mo6O21), the building blocks for M1 phase (Figure 2). These pentagons appear to contain heavy cations (e.g. Ta) in higher concentration than the bulk of the material, showing brighter HAADF contrast in the center. Two types of the interface between the pentagon units and the hexagon units in the M2 phase matrix are identified. Interestingly, monolayer pentagon coverage is dominantly found associated with the type I interface (Figs. 2b,c), where multilayer coverage is always found with the type II interface (Figs. 2d,e). In a M2 phase in a related Mo-V-Sb-Ta system, surface pentagon layers were also observed. In some areas, surface pentagon layers serve as seeds for the M1 phase attached to M2 surface. To our best knowledge, this is the first example of the intergrowth of these two very distinct structures.

Our observations suggest that the pentagon layers come from a self-assembly process of the preformed pentagon units during the synthesis.[7] The presence of Ta, a sub-group V element, is found crucial for this intergrowth. Comparing samples with and without Ta in composition clearly suggests that Ta stabilized the pentagon units during synthesis. Correlation of the microstructure with the catalytic performance will also be discussed. We believe that this work will pave the way for development of novel catalysts with M1/M2 heterostructures that can improve the catalytic properties by maximizing the synergistic effect.

[1] Pennycook, SJ et al., Philos. Trans. Roy. Soc. A(2009)
[2] Yu, JJ et al., Catal. Commun.(2012)
[3] Pyrz, WD et al., Chem. Mater.(2010)
[4] Zhu, YH et al., Chem. Mater.(2012)
[5] Zhu, Y et al., Angew. Chem. Int. Ed.(2012)
[6] Shiju, NR et al., Appl. Catal., A(2009)
[7] Sadakane, M et al., Angew. Chem. Int. Ed.(2009)


The MSE Division, US DOE; through a user project in ORNL’s CNMS, sponsored by the SUF Division, Office of BES, US DOE; The CSGB Division, Office of BES, US DOE.

Fig. 1: Polyhedral models of the ab planes of (a) the M1 and (b) the M2 phases in molybdenum vanadate oxide catalysts (e.g. MoVTeTa Oxide). Pentagon units, heptagonal channels, hexagonal channels and unit cells are highlighted in blue, yellow, red and black respectively. Sparsely occupied Te sites in heptagonal channels are omitted for clarity.

Fig. 2: (a) A HAADF image of the pentagon decorated M2 phase particle. (b,c)&(d,e) magnified views and proposed models of Types I (purple) and II (green) interfaces between the pentagon units (blue) and the hexagon units (red), respectively. Their representative units are highlighted in black. Heptagonal channels are highlighted in yellow.

Type of presentation: Poster

MS-1-P-3415 Synthesis and Electron Microscopy Characterization of Non-stoichiometric Tin Oxide Structures

Orlandi M. O.1, Suman P. H.1, Barbosa M. S.1, Longo E.1, Varela J. A.1
1Institute of Chemistry, Sao Paulo State University, Brazil
orlandi@iq.unesp.br

Nanomaterials have attracted the attention of researchers in the recent past due to their interesting properties. Moreover, there is a wide applicability of these materials in several areas of knowledge, for example, chemicals sensors, solar cells and microelectronic devices. However, in order to use these materials in devices it is very important to have a deep morphological and structural characterization of materials and electron microscopy techniques are important tools for these characterizations.

In this work, we used a carbothermal evaporation method to obtain tin oxide structures. By controlling the oxygen amount inside the furnace it was possible to grow structures in different oxidation states; i.e. SnO2, Sn3O4 and SnO. The samples were characterized by scanning and transmission electron microscopy, X-ray diffraction and electrical measurements. SEM characterization showed that SnO material is composed by nanobelts with spheres at their extremities and micro discs (which can be separated by sedimentation) while Sn3O4 and SnO2 materials are composed of nanobelts. TEM and SAD characterization showed that all synthesized materials are single-crystalline and SAD patterns in different zone axis enabled the determination of the growth direction of each material. Based on these results, it was possible to propose that SnO nanobelts grow following a self-catalytic vapor-liquid-solid (VLS) mechanism and Sn3O4 and SnO2 nanobelts grow by a vapor-solid (VS) process. The growth mechanism of discs are not completely known, but is related to the solidification of SnO vapor. While SnO2 is a well-known material in literature, there are only few works on SnO and Sn3O4 structures. Based on it we also studied the sensor response of materials to reducing and oxidizing gases and it was possible to correlate the sensor mechanism of materials with the exposed planes of structures.


We would like to thank the funding agencies FAPESP and CNPq for supporting this work.

Fig. 1: SEM image of SnO nanobelts. In the inset there is a TEM image of a single belt.

Fig. 2: SEM image of SnO microdiscs. In the inset there a SEM image of a single disc.

Fig. 3: a) SEM general view of Sn3O4 nanobelts. b-d) Details of Sn3O4 nanobelts.

Fig. 4: a) SEM general view of SnO2 nanobelts. b-d) Details of SnO2 nanobelts.

Type of presentation: Poster

MS-1-P-3419 Characteristaion of local surface plasmon resonance modes of metallic nanostructures with cathodoluminescence and electron energy loss spectroscopy in the (S)TEM

Stowe D. J.1, Wilkinson N.1
1Gatan Inc, R&D Headquarters, Pleasanton, CA, USA
dstowe@gatan.com

There is significant interest in understanding and controlling surface plasmon resonance modes of metallic nanostructures for use in chemical sensing and nano-optics applications. The optical effects of local surface plasmon resonance (LSPR) modes are only observable at a length scale below the diffraction limit of far field optical experiments meaning that alternative techniques are needed to study individual nanostructures. A focussed electron beam of a (scanning) transmission electron microscope (S)TEM can be used to generate local surface plasmon resonance (LSPR) modes; cathodoluminescence (CL) and electron energy loss spectroscopy techniques can be used to observe the generation and propagation of LSPRs with high spatial (<1nm) and spectral resolution (<2meV) and allow direct correlation with particle morphology. Here we report the use of CL and EELS in the characterisation of gold and silver nanoparticles using the Gatan Vulcan CL detector and Gatan GIF Quantum EEL spectrometer.
Figure 1 shows luminescence patterns of gold prisms acquired by (S)TEM-CL. Gold prisms of different size reveal very different LSPR patterns; a gold prism with sides of 250nm shows strong emission at 1.80eV (660nm) at the three corners (Figure 3c) whereas an 800nm particle shows luminescence dominated by a strong central LSPR node (1.53eV, 803nm) with other lower intensity LSPR modes along the edges (1.43 and 1.70eV, 870 and 705nm).
Light emitted from the specimen in the forward and backwards directions (in the direction of, and in the opposite direction to the fast electron) was measured (simultaneously) enabling differences in the intensity and spectral characteristics to be determined. A simple model system of a 120nm diameter sliver nanosphere was investigated. Preferential emission of the quadrapole LSPR mode was observed in the forward scattered direction whereas the dipole LSPR mode was preferentially emitted in the backward scattered direction, in agreement with Mie theory.


The authors would like to thank Dr M Bosman and Dr Ziyou Li for the provision of specimens and Dr G. Fern and Dr. D McComb for access to microscope facilities.

Fig. 1: Figure 1. (a) and (b) bright field STEM image of a two gold prisms (250 and 800nm sides respectively); (c) and (d) panchromatic STEM-CL images reveal ‘bright’ local surface plasmon resonance modes (e) and (f) CL spectra acquired from selected positions. 

Type of presentation: Poster

MS-1-P-3425 ELECTRON MICROSCOPY STUDIES OF 1D TiO2 NANOSTRUCTURES WITH LOW IMPROVED PHOTOCATALYTIC ACTIVITY

Acosta D.1, Cabrera J.2, Rodríguez J.2, Candal R.3, Alarcón H.2, López A.2, Arenas J.1
1Instituto de Física, UNAM, México D.F., 2Universidad Nacional de Ingeniería, Lima, Perú, 3Universidad de Buenos Aires, Argentina
jarenas@fisica.unam.mx

Nanowire/nanorods TiO2 nanostructures (NS) of approximately 8 nm in diameter and around 1000 nm long were synthesized by alkaline hydrothermal treatment of sol-gel made TiO2 or P-25 TiO2. Anatase like 1D TiO2 NS were obtained in both cases. The 1D NS made using seeds from Sol Gel TiO2 nanopowders turn on rod-like NS and presents lower surface area than the NS made from commercial TiO2 P-25 (97 y 279 m2/g, respectively).In both cases, the 1D NS showed lower photocatalytic activity than P25 nanoparticles. However, the rod-like NS obtained from TiO2 Sol Gel seeds displayed slightly higher efficiency than the original seeds. Despite the higher surface area shown by the NS, the photocatalytic efficiency did not improve with respect to their precursor seeds. This phenomenon can be associated with the lower crystallinity of 1D TiO2 in both materials.
From SEM and TEM micrographs of sol-gel TiO2 synthesized nanoparticles, large and compact aggregates (Figure 1a) with radius around 7 nm are detected. After 18 h of exposing the TiO2 sol-gel made seeds, to alkaline hydrothermal treatment, the NPs, turned to tube-like NS (Figure 1b) with inner and external diameter in average of approximately 5.6 and 8 nm respectively. After the hydrothermal treatment the samples were acid treated to exchange Na+ by H+; it was observed the tubular structure was conserved in spite of the acid treatment. Figure 1c displays SEM and TEM micrographs of TiO2 NS obtained by alkaline hydrothermal treatment of P-25 seeds during 18 or 24 h, followed by acid exchange and annealing at 400ºC. The 1D structures are preserved. As in the previous case, tubes were obtained by both hydrothermal treatment time (18 and 24 h). But in this case the tubular structure remains after annealing process. The images clearly shows that as consequence of the annealing the tube like NS turned into short rod-like particles. These results suggest that during the annealing the structure of the tubes collapsed, cutting the tubes in smaller pieces but preserving in part their original morphology.
The growth mechanism of 1D TiO2 NS synthesized by alkaline hydrothermal method from TiO2 nanoparticles is still under discussion. It has been suggested that it take place by the rolling of hydrogen titanate laminar structures during the ion exchange step (Kasuga, 1998).Other authors, suggest the 1D structure formation during the treatment of TiO2 in NaOH aqueous solution (Zhao, 2010). Our results seem to be in agreement with the last authors because the tubes were well formed, as it is observed in Figure 1b, before the application of acid treatment .


The financial support of CONCYTEC and Alianza del Pacifico program to J.Cabrera is here recognized. Also we thanks to Roberto Hernandez for technical help and the financial support of DGAPA-UNAM IN 105541 project to Dr. D. Acosta laboratory work.

Fig. 1: Figure 1. a) Electron micrographs of TiO2 1-D NS obtained from sol-gel and hydrothermally treated for 18 h and after acid treatment. b) 1-D NS from sol-gel nanoparticles, hydrothermally treated and after annealing at 400°C. c) TiO2 NS by alkaline hydrothermal treatment of P-25 seeds during 18 hrs, after acid treatment and annealed at 400°C.

Type of presentation: Poster

MS-1-P-3432 Lutetium Oxide Nanoparticles: Synthesis by a green method.

Schabes -Retchkiman P. S.1
1Instituto de Fisica, UNAM, Mexico City, Mexico
pabloschabes@yahoo.com.mx

The nanostructured systems present promising properties , optical, magnetic and catalyic properties. [1]. In particular a search for new designs at the nanometric scale has increased during the later years and its uses diversified.the interest in metallic oxides is due to applications in catalysis optical and electronic devices  [2-5]. Lutetium oxide has uses in laser crystals optical devices and ceramics, because of the opossibility of obtaining with high purity it is considered as target for Xray emission because of its high density.

In general this compound is used as a dopant in the formation of granate for laser crystals. Since the oxides . are electrically non-conductive but by developing certain structures the electrical conductivity can be achieved and used in the cathode of Fuell-Cells and Oxygen generation systems. Being an anhydrous compound of basic character, it allows redox reactions, which helps in wáter electrolysis. Furthermore, since it is an insoluble compound, and very stable it can be used in the fabrication of cathodes of light ceramic structure in fuel-cells saving on weight. For the use of this oxide three important parameters are looked after: size, structure and elementary composition.

Lutetium Oxide nanoparticles were obtained in a high yield by means of biosynthesis with alfalfa. The morphology and crystal structure were characterized by high resolution transmission electron microscopy, EELS and Z-contrast microscopy. It is shown that the synthesis method employed produces small Lutetium Oxide nanoparticles, (2-5 nm), Lu2O3, as characterized from the optical diffractograms of individual particles.

The synthesis of  Lu2O3 nanoparticles was done by the bioreduction method similar to the one proposed by G. Canizal [6], for gold nanoparticles . The synthesis was designed in principle to obtain nanoparticles of metallic Lu, but due to the characteristics of Lu the metallic oxide was obtained. Figure 1 shows size distribution function for Lutetium NP. Figure 2 shows pH7 HAADF of Lutetium NPs, Figure 3 presents HRTEM of Lu oxide particles, showing some have coalesced for pH7. Conclusion; The bioreduction method has been ueful to produce Lu2O3 NP of quantum size character.
1. H. S. Nalwa, Handbook of Nanostrucutred Materials and Nanotechnology, Academic Press, San Diego, CA (2000).
2. M. L. Wu, D. H. Chen and T. C. Huang, J. Colloid Interf. Sci. 243, 102 (2001).
3. B. Veisz, L. Toth, D. Teschner, Z. Paal, N. Gyorffy, U. Wild and R. Schlogl, J. Mol. Catal. A-Chem. 238, 56 (2005).
4. M. O. Nutt, J. B. Hughes and M. S. Wong, Environ. Sci. Technol. 39, 1346 (2005).
5. R. Narayanan and M. A. El-Sayed, J. Phys. Chem. B 109, 12663 (2005).
6. G. Canizal, J. A. Ascencio, J. Gardea-Torresday, M. Jose´-Yacaman, J. Nanopart. Res. (2001) 475.


The authors thank Ms Indira Blanco, Mr. L. Rendón Vázquez and the LACMIF-UNAM personnel for technical help. This research has been partially supported by DGAPA-UNAM project # . IN120006

Fig. 1: Size distribution function for Lu Oxide NPs

Fig. 2: HAADF showing Lutetium Nanoparticles distinct contrast.

Fig. 3: Lutetium oxide nanoparticles, typical shapes and structures

Type of presentation: Poster

MS-1-P-3446 Structure and Morphology Study of Pure and Mixed ZnO and ZnO2 Nanoparticles

Paraguay-Delgado F.2, Roman E.1, Gómez M. M.1, Solís J. L.1, Antunez-Flores W.2
1Universidad Nacional de Ingeniería, Faculty of Science, Av. Túpac Amaru 210, Lima 25, Perú, 2Centro de Investigación en Materiales Avanzados S. C., CIMAV, Miguel de Cervantes 120, CP 31109 Chihuahua, Chih. México.
francisco.paraguay@cimav.edu.mx

Synthesis of ZnO2 nanoparticles was performed via a sol–gel technique assisted with UV irradiation. One gram of zinc acetate dehydrate, Zn(CH3COO)2.2H2O, was dissolved under vigorous stirring in a mixture of 50 ml distilled water and 5 ml of 30% H2O2. The resulting solution was then irradiated with a 300W Ultra-Vitalux lamp (Osram), positioned 10 cm above the solution, for 30 min at ambient temperature. This procedure resulted in the formation of a white zinc peroxide colloidal suspension. The ZnO2 nanoparticles were precipitated by centrifugation. The precipitate was then washed using distilled water until a pH of 8 was reached. Finally the resultant white solid was dried at 80 °C for 12 h, similar to follow in the reference [1]. The resultant powder was annealed between 100 and 220°C for 1 h in an oven with air atmosphere. The morphology, structure and domain size of the nanoparticles were determined by X-ray diffraction, and scanning transmission electron microscopy. By X-ray diffraction, all patterns can be indexed to the zinc peroxide phase for samples prepared up to 120°C. For a sample prepared at 160°C we had a mixture of ZnO2 and ZnO, while for particles treated at 220°C all the material was pure ZnO.
Micrographs shows STEM images for zinc oxide and zinc peroxide nanoparticles. Fig 1 shows rounded ZnO particles, with an average grain size of 18±5 nm. The inset displays that the ZnO d-space was 2.8 Å. Fig 2 shows an image of ZnO and ZnO2 mixture, in the inset figure can be appreciated rund conglomerated particles. There are two types of particles, the bigger ones belong to ZnO and the smaller ones belong to ZnO2. The information of the atomic columns acquired by HAADF detector indicated that ZnO d-spaces were between 2.8 Å and 2.6 Å. This parameter must be connected to synthesis conditions of the material. In any case the average diameter size was 145 ± 55 nm.
Figures 3 and 4 belong to images of pure ZnO2 particles acquired by HAADF and BF detectors respectively. At low magnification can be observed spherical shapes with broad size dispersion between 40 and 287 nm. The average diameter was 130±64 nm. At higher magnification these conglomerates displays small grains (≈ 5 nm). Figure 1d confirmed that each small grain had d-space values which belong to ZnO2.
Using electron microscope techniques we have studied in detail the morphology and the structure of ZnO nanoparticles, ZnO2 nanoparticles and a mixture of both. The ZnO2 nanoparticle are of great interest, because they had interesting microbiological characteristics [2].

References
[1] R Colonia, J L Solís and M Gómez, Adv. Mat. Sci.: Nanotechnol 5 (2014) 015008 (4pp).
[2] R Colonia, V. Martinez, J. L. Solís and M. M.Gómez, Rev. Soc. Quim. Peru 79(2)2013 126 (10pp)


Thanks to P. Pisa and E. Torres for their technical help, at nanotech Cimav.

Fig. 1: STEM image by HAADF and BF images for ZnO particles, inset image show d space of ZnO.

Fig. 2: ZnO-ZnO2 particles, d-sapace belongs to ZnO.

Fig. 3: ZnO2 particles conglomerate in a spherical shape, inset can be notice the nanoparticles.

Fig. 4: BF image showing d-spaces of ZnO2.

Type of presentation: Poster

MS-1-P-3458 Blue luminescence related to stacking faults in ZnO:Ag

Tsiaoussis I.1, Potin V.2, Bourgeois S.2, Khranovskyy V.3, Eriksson M.3, Yakimova R.3
1Department of Physics, Solid State Physics Section, Aristotle University Thessaloniki,, 2I.C.B. - Laboratoire Interdisciplinaire Carnot de Bourgogne, UMR 6303 CNRS - Université de Bourgogne 21078 DIJON CEDEX, France , 3Department of Physics, Chemistry and Biology (IFM), Linköping University, 58183, Linköping, Sweden
tsiaous@auth.gr

Ag was used as catalyst for growth of vertical ZnO nanorods. Moreover, Ag also acts as an amphoteric dopant, existing both on substitutional Zn sites and in the interstitial sites, as substitutional acts as acceptor. Therefore, the effects of Ag on the optical properties of ZnO is of importance. We have grown the ZnO microrods by MOCVD with Ag catalyst on Si (100) substrates at Ts = 773 K. The elongated quasi single crystal structures were observed to have corrugated side facets.

The presence of SFs affects the luminescence properties of ZnO by creating additional peak at 3.321 eV (at 4 K). This peak co-exists with the common donor bound excitonic (D0X) and free excitonic recombination’s peaks as well as their respective phonon replicas (Fig.1). Furthermore, the SFs related peak is stable at least up to 350 K, providing splitting of the near band edge emission of ZnO into two peaks: FX emission appears at 375 nm and this of SFs at 386 nm (Fig.2). Therefore, visualizing of the two emissions was performed by CL mapping: the two types of emission are spatially resolved. These radiative recombination processes are under investigation by time-resolved PL. It is proposed, that at high Ag concentrations, the SFs formation is favored.

High concentrations of basal SFs were found to be responsible for the surface corrugation. A TEM study of the cross-section of ZnO/Ag/Si has revealed additional unusual contrast in bright field mode. The featured lines were located parallel to the substrate plane, e.g. perpendicular to the c-axis of the NR. Presuming that the reason for this is the extended defects, we have studied individual NRs by HRTEM. A number of basal plane [0001] stacking faults were observed, penetrating the NRs perpendicular to its c-axis (Fig. 3). BSFs were found to be quasiperiodically inserted every 5 - 10 nm along the NRs. It has to be mentioned that SFs are observed in both types of NRs. We attribute the appearance of BSFs as due to the Ag dopants.

Theoretical study in wurtzite GaN by Schmidt et al. [T. Schmidt et al. Phys. Rev. B 65 033205 (2002)] has shown that it is energetically favourable for Mg atom to reside at some distance from the fault plane. We believe, that in our case Ag atoms behave similarly, and their availability in the proximity of SFs transform them to a radiative recombination centre, existing up to room temperature.


We acknowledge the Linköping Linnaeus Initiative for Novel Functional Materials (LiLi-NFM) for supporting this work. Dr. Ioannis Tsiaoussis would like to thank Dr. Valerie Potin for enabling the TEM experiments at the University of Burgundy.

Fig. 1: LT PL spectrum (4 – 100 K) of ZnO:Ag microrods (the SEM image is as inset)

Fig. 2: The micro-PL spectra taken along the microrod: the transition of the PL intensity from UV to Blue emission is shown. The light emissions are visualized by CL mapping (inset).

Fig. 3: High concentration of basal SFs were found to be responsible for the surface corrugation as it is seen in the HRTEM image.

Type of presentation: Poster

MS-1-P-3482 Study of contact behaviour of TiO2 nanoparticle agglomerates by means of AFM and in-situ TEM

Salameh S.1, Mädler L.1, Seo J. W.2
1Foundation Institute of Materials Science (IWT), Department of Production Engineering, University of Bremen, Bremen, Germany, 2Department of Materials Engineering, KU Leuven, Leuven, Belgium
maria.seo@mtm.kuleuven.be

Adhesion forces between individual nanoparticles play an important role in many different processes such as fluidization, agglomeration and coating. Currently, adhesion between particles is interpreted in terms of continuum models that are able to take into account the effects of capillary forces, surface roughness and electrostatics. This approach is generally suitable for particle sizes in the micrometer range. However, for smaller particles with characteristic sizes in the range of 10 nm, more subtle effects beyond continuum theories can influence and even dominate the adhesion behavior. We have studied adhesion forces and contact behaviour of TiO2 nanoparticles with a diameter in the range of about 10 nm. These nanoparticles were produced in a flame spray reactor using the liquid precursor consisting of 0.5 molar Ti(IV) isopropoxide in xylene. Inside a TEM, we studied stretching and de-agglomeration behaviour of TiO2 nanoparticle agglomerates using an AFM/TEM holder. These in-situ observations were correlated with the force measurements obtained from AFM force spectroscopy. To be precise, the AFM data were based on the statistical analysis of the force peaks measured in repeated approaching/retracting loops of an AFM cantilever into a film of nanoparticle agglomerates. The in-situ TEM data revealed sliding and rolling events first leading to local rearrangements in the film structure when subjected to tensile load, prior to its final rupture caused by the reversible detaching of individual nanoparticles. The associated contact force of about 2.5 nN is in quantitative agreement with the results of Molecular Dynamics simulations of the particle-particle detachment [1]. Our results indicate that the contact forces are dominated by the structure of water layers adsorbed on the particles’ surfaces at ambient conditions. This leads to non-monotonous force-displacement curves that can be explained only in part by classic capillary effects, and highlight the importance of considering explicitly the molecular nature of the adsorbates [1].

We also studied the size dependent contact behavior of nanoparticle agglomerates by using four different size-fractionated agglomerates, with median values in the range of 78 to 161 nm. Force-distance curves of AFM as well as in-situ TEM observations show that the length of the chains and the amount of rearrangements depend on the agglomerate sizes. Larger agglomerates require more work for their aggregate rearrangement before the final breakage is induced [2].

[1] S. Salameh , J. Schneider, Jens Laube, A. Alessandrini, P. Facci, J. W. Seo, L. Colombi Ciacchi and L. Mädler, Langmuir 28 (2012), p. 11457. Langmuir 28 (2012), p. 11457.

[2] Salameh, R. Scholz, J.W. Seo and L. Mädler, Powder Technology 256 (2014) p. 345.


We acknowledge the Flemish Hercules Stichting (HER08/25), KU Leuven STRT1/08/025 and DFG for funding this project SPP 1486 under grants MA 3333/3 and CO 1043/3.

Type of presentation: Poster

MS-1-P-3503 Imaging and microanalysis of plasmonic Ga nanoparticles

Suvorova A. A.1, Losurdo M.2, Brown A. S.3, Rubanov S.4, Bruno G.2
1Centre for Microscopy, Characterisation and Analysis, The University of Western Australia, Perth, Australia, 2Institute of Inorganic Methodologies and Plasmas at National Council of Research (CNR), Bari, Italy, 3Electrical and Computer Engineering Department, Duke University, Durham, North Carolina 27705, United States, 4Electron Microscope Unit, Bio21 Institute, University of Melbourne, Melbourne, Victoria 3010, Australia
alexandra.suvorova@uwa.edu.au

Plasmonic nanoparticles (NPs) are of considerable interest due to plasmon tunability and potential applications as biosensors, photonic, optoelectronic and photovoltaic devices. These applications rely on the fabrication of metallic NPs on technologically important substrates and on the possibility to control the surface plasmon resonance (SPR) properties. The ability to create tunable (from the UV to the visible) plasmonic nanosystems using Ga NPs is differentiating gallium from the commonly used noble (gold and silver) metals. In our previous work, we have demonstrated the efficiency of Ga NP-based platforms in localized surface plasmon resonances (LSPR) tunable over the UV to the near IR spectral range1-3.

Here we describe a range of imaging and microanalysis electron microscopy techniques that are highly suitable for the study of the Ga-based plasmonic nanosystems. Ga nanoparticles were deposited onto sapphire, silicon, glass and graphene substrates in a Veeco GEN II molecular beam epitaxial system under ultrahigh vacuum conditions at room temperature with a constant Ga flux. TEM cross-sectional samples were prepared by the FEI Nova dual beam focused ion beam (FIB) system. A range of microanalytical electron microscopy techniques can be used to characterise the NPs at the nanoscale level. Scanning electron microscopy (SEM) imaging has been applied to study morphology and growth dependent modifications of the Ga NPs. Transmission electron microscopy (TEM) and associated analytical tools have been used to determine the structural and compositional properties of the nanostructures at a subnanometer scale. High resolution imaging revealed crystalline core/ amorphous shell structure for Ga NPs grown on sapphire and amorphous  Ga structure for Ga NPs grown on other substrates. Energy-filtered imaging  showed compositional uniformity of the Ga core and the presence of oxide layer on the NPs surface. Low-loss imaging confirms the presence of Ga, with particle contrast being maximised close to the Ga plasmon energy (13.8eV). The Ga plasmon signal is significantly higher for the crystalline core of the Ga NPs. In summary, a range of electron microscopy techniques can be used to identify and characterise Ga NPs at the nanoscale level. Such information is important for understanding structural and optical properties of Ga-based nanosystems.

1Wu, P. C.; Kim, T. H.; Brown, A. S.; Losurdo, M.; Bruno, G.; Everitt, H. O. Appl.Phys. Lett. 2007, 90, 103119.

2Yi, C.; Kim, T. H.; Jiao, W.; Yang, Y.; Lazarides, A.; Hingerl, K.;Bruno, G.; Brown, A. S.; Losurdo, M. N. Small 2012, 8, 2721–2730.

3M Losurdo, C Yi, A Suvorova, S Rubanov, T-Ho Kim, M M. Giangregorio, W Jiao, I Bergmair, G Bruno, A. S. Brown ACS Nano 2014 in press


Fig. 1: SEM image of Ga NPs grown on sapphire.

Fig. 2: TEM image showing the Ga NPs in cross-section.

Fig. 3: EFTEM imaging: zero-loss image of  Ga NPs.

Fig. 4: Low loss imaging of Ga NPs showing core/shell structure

Type of presentation: Poster

MS-1-P-3507 STEM Electron Diffraction and High Resolution Used in the Determination of Thiolated Gold Clusters

Ponce A.1, Bahena Uribe D.1, Tlahuice Flores A.1, Santiago U.1, Whetten R.1, Jose Yacaman M.1
1University of Texas at San Antonio
arturo.ponce@utsa.edu

Gold nanoparticles protected by thiolate ligands have attracted extensive research activity because of their enhanced optical, electrochemical, and other application-related properties. The desired physicochemical properties are strongly dependent upon size, composition, and structure. It is therefore critical for improved materials design to correlate the cluster structure and bonding with the properties sought for applications. However, the determination of atomic structures of nanomaterials is a challenging task even for composition elucidated gold−thiolate nanoclusters where gold cores span various symmetries. Determination of the total structure of molecular nanocrystals is an outstanding experimental challenge that has been met, in only a few cases, by single-crystal X-ray diffraction. Described here is an alternative approach that is of most general applicability and does not require the fabrication of a single crystal. The method is based on rapid, time-resolved nanobeam electron diffraction (NBD) combined with high-angle annular dark field scanning/transmission electron microscopy (HAADF-STEM) images in a probe corrected STEM microscope, operated at reduced voltages. In the current presentation, we will show the new structures of Au130 and Au144 thiolated clusters explored in a combined experiment-theory approach [1-2]. A full map in reciprocal space has been simulated and compared with the experimental patterns using STEM diffraction as well as atomically resolved images obtained through aberration-corrected STEM-HAADF images. The nanobeam diffraction (NBD) through the STEM imaging mode is controlled by the condenser lens system. The combination of probe-corrected STEM imaging and quasi-parallel beam diffraction (D-STEM) is obtained by positioning the beam in the STEM image at a single nanoparticle using the Digiscan control. The scan is stopped and positioned arbitrarily at a xy position on the screen. Subsequently, the electron diffraction pattern is recorded using a digital charge couple device (CCD) camera. D-STEM mode works in the diffraction plane, the overlapping of the convergent disks is optimized by a compensation of the last condenser lens (C3) and the use of the adaptor lens (ADL) at the hexapole coils of the CEOS corrector.

[1] A. Tlahuice-Flores, U. Santiago, D. Bahena, E. Vinogradova, C.V. Conroy, T. Ahuja, S.B.H. Bach, A. Ponce, G. Wang, M. Jose-Yacaman and R.L. Whetten, J. Phys. Chem. A, 2013, 117 (40), pp 10470–10476
[2] D. Bahena, N. U. Santiago, A. Tlahuice, A. Ponce, S.B.H. Bach, B. Yoon, R.L. Whetten, U. Landman and M. Jose-Yacaman, J. Phys. Chem. Lett., 2013, 4 (6) , 975-981.


This project was supported by grants from the National Center for Research Resources (5 G12RR013646-12) and the National Institute on Minority Health and Health Disparities (G12MD007591) from the National Institutes of Health. The authors would also like to acknowledge the NSF PREM # DMR 0934218.

Type of presentation: Poster

MS-1-P-3512 Electron beam induced surface modification of semiconductor nanowires in a chlorine environment - A new route to electrical tailoring of nanodevices

Wanzenboeck H. D.1, Mika J.1, Shawrav M. M.1, Gavagnin M.1, Ismail B.1, Zeiner C.1, Lugstein A.1, Stoeger-Pollach M.1, Bertagnolli E.1
1Vienna University of Technology, A-1040 Vienna, Austria
heinz.wanzenboeck@tuwien.ac.at

Scanning electron microscopy is not only a high-resolution imaging technique for nanocharacterisation of materials, but the focused beam of electrons can also be used for inducing chemical reactions in the nanometer-regime. For focused electron beam induced processing (FEBIP) precursor gas is introduced into the vacuum chamber and the electrons interact with precursors adsorbed on the sample surface. It has already been demonstrated, that metalorganic precursors lead to deposition of materials including noble metals such as Pt or Au as well as magnetic metals such as Fe or Co.

We have recently introduced a controlled etching process that is sustained by the irradiating electron beam. With the semiconductors Si and Ge we have not observed spontaneous etching, while material could be etched in the areas exposed to the electron beam. A clean vacuum chamber is a prerequisite for this process and has been achieved with an in-situ ozone cleaning procedure of the chamber.

In this work we report on the controlled etching of Si-nanowires and of Ge-nanowires. Nanowires themselves are smart nanomaterials with very promising characteristics and may be used for nanoelectronic devices, innovative sensor concepts and for photovoltaics applications. Focused electron beam induced etching (FEBIE) offers a further alternative to modify these nanomaterials in-situ in a SEM. With dynamic experiments in the SEM we have investigated the chemical reactions on the nanoscale. The material modification with regard to, its composition and its electrical has been investigated.

The custom-designed tailoring of electrical properties of nanowires is essential for the development of new devices. FEBIE is a versatile approach for trimming of Si-nanowires as the low-energy electrons inflict no significant crystallographic damage and cause no contamination of the silicon nanowire. This in-situ preparation allows to keep specimens as close as possible to their native state.

With chlorine as etch gas even without geometrical thinning of nanowires the short-term irradiation was observed to result in a change of electrical properties towards a diode-like characteristics. The effect of electron exposure under the presence of molecular chlorine was investigated. Additional to structural studies also an electrical characterisation of contacted Si-nanowires and a TEM nanostructure analysis was performed.

FEBIE has been established as a novel approach that allows for tailoring of material properties by controlled in-situ modification of nano-scaled materials. Potential future applications of FEBIE to design and to develop of new nanomaterials for sensor applications and for photonics applications will be discussed.


We acknowledge financial support by the Austrian Science Fund (FWF) under project P24093. TEM analyses were carried out at the University Service Centre for Transmission Electron Microscopy, Vienna University of Technology.

Fig. 1:  Scanning Electron Microscope LEO 1530 VP with the custom-built gas injection system for chlorine

Fig. 2: Schematic illustration of nanowire modification by focused electron beam induced etching.

Fig. 3: Left: Setup for the electrical measurement of the semiconductor nanowire. The top image shows a SEM image in top view Right: Electrical behavior of the Si-nanowire. The I-V-curve shows the electrical properties before and after FEBIE modification in a chlorine environment.

Type of presentation: Poster

MS-1-P-5699 Microscopic and spectroscopic techniques as useful tools in the construction of nanobiosensors based on carbon nanotubes

Guadarrama L.1, Chanona J.1, Hernandez H.1, Manzo A.2, Martínez A.3, Calderon G.1, Suarez E.1
1Escuela Nacional de Ciencias Biológicas, Instituto Politécnico Naciona, 2Escuela Superior de Ingeniería Química e Industrias Extractivas, Instituto Politécnico Nacional, 3Centro de Investigación en Computación, Instituto Politécnico Nacional
jorge_chanona@hotmail.com

A biosensor is a device, which converts a biological response between a target analyte and a bioreceptor into an electrical signal. The bioreceptor can be a microorganism, organelle, cell, enzyme, antibody, nucleic acid etc. All these kind of sensors can exploit the advantages from high surface-to-volume-ratio property of nanomaterials. Carbon nanotubes (CNT) present outstanding electrical and chemical properties and could interact with organic and inorganic compounds therefore can be functionalized with supramolecular complex. The morphology and quality of CNT can be determined by the use of transmission (TEM) and scanning (SEM) electron microscopy. Raman and X-ray photoelectronic (XPS) spectroscopy allow identify the type of CNT and verify the chemicals modifications produced on CNT during the functionalization process. These techniques provide useful information about the biosensor construction process like homogeneity of the CNT network or wide of the layer. This project presents the microscopic and spectroscopic characterization of multi-walled CNT simultaneously purified and functionalized through an acid treatment with HNO3-H2SO4 with the objective of make them more reactive trough the formation of acid carboxylic groups on the CNT and then use it as support for amyloglucosidase (AMG) as a probe molecule and check if the enzyme is still active. All CNT were analyse using SEM, TEM, Raman and XPS. Raman spectra allow observe how the acid treatment removes impurities of the CNT (Fig 1). Chemical functionalization with carboxylic groups is evidenced by XPS spectra showing a peak at 288.5 eV characteristic of carboxylic groups and a shoulder at 287 eV possibly associated with peptide bonds (Fig 2). By the TEM micrographs (Fig 3) it is possible to observe the enzyme onto CNT contrasted with uranyl acetate (1%). The results of enzymatic assay prove that the AMG preserve 50 % of its activity compared against native enzyme. The double-layer capacitance, obtained from the current versus potential characteristics at different scan rates (mV/s), was also obtained for the CNT in each step of functionalization and finally the CNT/AMG system was tested in optimal conditions for the AMG catalyze the substrate. The different responses in the electrochemical capacitance and the results of enzymatic kinetics provide evidence of an adequate functionalization of CNT for their use as electrochemical biosensor. Microscopic and spectroscopic techniques prove that are not only useful but necessary tools for biosensors construction.


CONACyT and PIFI-IPN for the scholarship. Financial through the projects 20131864, 20130333, 20140387 and Red-SIP-NTC-FET at IPN-Mexico and from CONACyT 161793, 133102. The TEM images were obtained with the project supported by a grant from the National Institute on Minority Health and Health Disparities (G12MD007591) from the National Institutes of Health.

Fig. 1: Raman spectra for A) raw carbon CNT and B) purified CNT

Fig. 2: XPS spectra for raw CNT, purified CNT and system CNT-AMG. Inset: zoom in the range 286-291 eV.

Fig. 3: TEM image of a CNT covered with AMG and contrasted with uranyl acetate at 1%

Type of presentation: Poster

MS-1-P-5750 A Comprehensive study of the nucleation and growth mechanism of Au anisotropic nanoparticles: From Seeds to Bipyramids

Ihiawakrim D.1, Ersen O.1, Hirlimann C.1, Florea I.2, Treguer-Delapierre M.3, Majimel J.3, Spuch-Calvar M.3
1IPCMS, CNRS-Université de Strasbourg, Strasbourg, France, 2LPICM, CNRS-Ecole Polytechnique, Palaiseau, France, 3ICMCB, CNRS-Université de Bordeaux, Pessac, France
dris.ihiawakrim@ipcms.unistra.fr

In the framework of the development of nano-sized materials with new optical properties induced by the size effect or by a specific morphology, the study of noble metal nanoparticles takes nowadays a prominent position, due in particular to their plasmonic properties. Attention has been paid to the quest for a synthesis method able to provide Au nanostructures with precise shape and crystallographic orientation. Thus, it was demonstrated that by using a seed mediated technique [1] and tuning the Au seeds concentration that play the role of nucleation centers, one can synthesize Au bipyramids (BPs) with various aspect ratios inducing various symmetries in the basal plane. From a fundamental point of view, it is expected that changing the morphology of these NPs may induce modifications of their optical properties. To synthetize Au BPs with controlled aspect ratios a good understanding of the nucleation and growth mechanisms is needed. The goal of this work is to perform a comprehensive analysis based on an approach combining modern TEM techniques: conventional TEM imaging mode, HR imaging using both TEM and STEM HAADF modes [2] and STEM-HAADF electron tomography. This type of analysis provides reliable information regarding the morphology and the crystallographic structure of both Au seeds and Au BP and thus allows elaborating reliable hypothesis on the nucleation and growth processes of these anisotropic NPs. We present here a detailed study performed on Au NPs presenting aspect ratios of 2, 3, and 5. An icosahedral (penta-twinned decahedron) shape of Au seeds NPs presenting a 4 nm size was firstly evidenced (Fig. 1). In a second step, a detailed HRTEM analysis on Au BP allowed us to directly observe the highly stepped nature of the BP surface constituted by {151} lateral facets (Fig. 2). Finally, the analysis of the reconstructed volumes obtained by electron tomography showed that the bipyramidal morphology is preserve for all the studied nanoparticles (Fig. 3a). However, some differences between them can be observed regarding the shape of their tips, the symmetry of the equatorial plane and the characteristics of the steps present on the surface (Fig. 3b). Particularly, the larger the BP is, the sharper the tips and higher the surface steps are. In addition, the analysis of the transversal sections for each volume showed that the symmetry of the equatorial plane changes from a hexagonal one for high aspect ratios to a pentagonal one for low ratios.

References:

[1] Liu, M.;, Guyot-Sionnest, P.J., J. Phys. Chem. B, 2005, p 22192.

[2] Burgin, J. et al. Nanoscale 2012 p.1299.


The authors gratefully acknowledge funding from the ANR under Grant number ANR-BLANSIMI10-LS-100617-15-01.

Fig. 1: (a) HR-STEM HAADF image of a 4 nm Au seed NP; (b) Projection at 0° extracted from the tilt series used to reconstruct the volume of an area containing several Au seeds NP;(c) XY slice through the reconstructed sub-volume of an individual Au seed NP evidencing its icosahedral morphology .

Fig. 2: (a), (b) High Resolution TEM micrographs of an Au bipyramid showing the crystallographic nature of the stepped lateral facets; (c) Schematical representation illustrating the oriented assembling of Au seeds icosahedrons.

Fig. 3: (a) 3D Models of Au bipyramids with three various aspect ratios obtained by electron tomography; (b) Illustration of the presence of steps on the external surface of the BP.

Type of presentation: Poster

MS-1-P-5771 Application of EFTEM and EELS for investigations of electronic and structural properties of nanostructures

Sobczak K.1, Borysiuk J.1, Li T.1, Dabrowski J.1, Dluzewski P.1
11Institute of Physics PAS al. Lotników 32/46 ,PL 02668 Warsaw
ksobczak@ifpan.edu.pl

Transmission electron microscope offer wide range of measurement possibilities like high efficient electron energy loss spectroscopy (EELS) and energy filtered transmission electron microscopy (EFTEM)[1]. That allows, among others, the measurement of surface plasmons resonance (SPR) and in the case of structures with the size below tens nanometers – Localized Surface Plasmons Resonance (LSPR)[2].
The phenomena of electron energy loss can be exploited in both scanning and imaging working mode of TEM. In the scanning mode a spectrum is recorded at a given beam position and therefore the spatial resolution is determined by beam size and specimen thickness, which indicate a volume from which the spectrum is collected. The energy resolution depends on monochromaticity of the incident electron beam and the quality of a spectrometer. Nowadays  the energy resolution of monochromator is 0,2 eV at a 300 keV and about 0,15 eV at a 80 keV. STEM mode connection with such a good energy resolution allows the measurement of the plasmon resonance eve n in nanostructures with dimensions less than 5 nm.
EFTEM technique is to form an image with electrons within a certain kinetic energy range. The EFTEM standard procedure for elemental mapping is based on recording three images: two pre-edge images with electron energy loss window before an absorption edge and one post-edge image with energy window after the absorption edge. The element mapping is obtained by the post-edge image after removing a background extrapolated from two pre-edge images. The intensity of the resulting image is proportional to the concentration of the element for which the absorption edge was used. Fig. 1 and 2 shows the application of the EFTEM elements mapping procedure in the case of AlN/GaN heterostructure.
EFTEM is very useful method to mapping of the distribution of elements in the investigating sample. Fig.3 and 4. present maps obtained by EFTEM method for Ag nanoparticles. An interesting application of EFTEM is mapping of electronic properties with a use of plasmon absorption [3]. It is well known that plasmon excitation energy depends on the mobility and density of charge carriers. The energy of plasmon’s peak is correlated with local structure and therefore gives information about relation between electronic properties and structural defects.


[1] F.J. Garcia de Abajo. Optical excitations in electron microscopy Rev. Mod. Phys. 2010;82:209-275
[2] JA. Scholl, AL. Koh, JA. Dionne. Quantum plasmon resonances of  individual metallic nanoparticles Nature 2012;483:421-428
[3] J. Nelayah, M. Kociak, O. Stephan, FJG. de Abajo, M. Tence, L. Henrard, D. Taverna, I. Pastoriza-Santos, LM. Liz-Marzan, Ch. Colliex. Mapping surface plasmons on a single metallic nanoparticle. Nature Physics 2007;3:348-353


The project was financed by The National Science Centre Nr DEC-2012/05/N/ST3/03163, DEC-2011/03/B/ST5/02698 and No. POIG.02.01-00-14-032/08.

Fig. 1: a) HRTEM image of GaN dopped with Al

Fig. 2: b) Mapping of the Al by EFTEM method.

Fig. 3:  a) HRTEM image of silver nanoparticles

Fig. 4: b) Mapping of N line of the Ag using The EFTEM method.

Type of presentation: Poster

MS-1-P-5778 The use of Microscopy in the development of certified reference materials for nanotechnology

Gerganova T.1, Roebben G.1, Kestens V.1, Kerckhove G.1, Vissers R.2, Boydens E.2, Held A.1, Emons H.1
1European Commission, Joint Research Centre, Institute for Reference Materials and Measurements (IRMM), Retieseweg 111, 2440 Geel, Belgium, 2Umicore Group R&D, Analytical Laboratory, Watertorenstraat, B- 2250 Olen, Belgium
Tsvetelina-Ivanova.Gerganova@ec.europa.eu

The use of Microscopy in the development of certified reference materials for nanotechnology

Authors: Tsvetelina Gerganova1, Gert Roebben1, Vikram Kestens1, Giovani Kerckhove1, Rita Vissers2, Eddy Boydens2, Andrea Held1, Hendrik Emons1

1) European Commission, Joint Research Centre, Institute for Reference Materials and Measurements (IRMM), Retieseweg 111, 2440 Geel, Belgium

2) Umicore Group R&D, Analytical Laboratory, Watertorenstraat, B- 2250 Olen, Belgium                       

Abstract: Nanotechnology is an enabling technology, which has the potential to greatly improve many areas of human life through newly developed nanomaterial-based products. The safe use of such products requires an understanding of the possible release of nanoparticles from such products, and an assessment of the potential hazards resulting from the exposure to these particles. This understanding and assessment must be underpinned by accurate measurement data of both the chemical and physical properties of nanoparticles.

Amongst others, size and shape have been identified as important properties of nanoparticles. In this context microscopy is a powerful technique to investigate and to quantitatively assess both metrics. The accuracy of the results obtained with microscopy techniques can only be guaranteed if the measurements are performed in a metrologically sound manner. Certified reference materials (CRMs) are indispensable tools for method performance verification and validation, as they enable a quantitative assessment of the method trueness and the uncertainty of measurement results. This work discusses the value and limitations of different microscopy techniques in the development of CRMs for use in the size and shape measurements of nanoparticles. 


Type of presentation: Poster

MS-1-P-5796 3D structural & chemical analysis of ternary PtRh/SnO2 catalysts for Ethanol Oxidation in Direct Ethanol Fuel Cells

Parlinska M.1,2, Li M.3, Kowal A.4,5,6,7
1Facility for Electron Microscopy & Sample Preparation, Univ. of Rzeszow, Poland, 2Int. Centre of Electron Microscopy for Materials Science, Dept. of Physical & Powder Metallurgy, Faculty of Metal Engineering & Industrial Computer Science, AGH University of Science and Technology, Krakow, Poland , 3Dept of Chemistry, Brookhaven National Laboratory, Upton, New York 11973, United States, 4Dept of Mat. & Natural Sci., Univ. of Rzeszow, Poland, 5Center for Synthesis and Char. of Nanomaterials, Univ. of East Sarajevo, Bosnia & Herzegovina, 6Central Lab. of Batteries & Cells, Forteczna 12, Poznan, Poland , 7Elcatak, ul. Pod Sikornikiem 8, Krakow, Poland
nckowal@cyf-kr.edu.pl

One possible solution for new, efficient, and environmentally-friendly technologiy for transforming chemical energy into electricity are fuel cells [1]. Ethanol seems to be an ideal fuel, as it is a non-toxic liquid and can be produced cheaply and efficiently from grasses. The best performance in ethanol oxidation reaction, was obtained for PtRh/SnO2 nanoparticles designed and synthesized by the Adzic group [2-5]. Additives such as F or Sb can be added into SnO2 in order to enhance its catalytic activity [6]. Structural and chemical investigations of the nanocatalysts with different Pt:Rh:Sn ratio were performed by TEM Osiris FEI operating at 200 kV and equipped with Super-EDX. Fig. 1(a) shows a STEM HAADF image of the PtRh/SnO2 catalyst with a Pt:Rh:Sn ratio = 1:1/3:1, which is relatively homogeneously distributed on the Vulcan carbon substrate. Unfortunately, a direct distinction between the tin oxide and PtRh particles is not possible, Fig. 1(b). Individual particles with a size from 2-10 nm are observed. Fig.2 shows the HAADF image of the PtRh/SnO2 particles with quantified EDX maps of Pt, Rh and Sn. Individual Pt particles of 2-5nm in the map are distinguishable, what is not the case for the SnO2 particles. The Rh signal is rather weak and is located in the same areas as Pt. The SnO2 particle size starts from 4-5 nm for individual particles. In the sum EDX map of Pt, Rh and Sn, it can be clearly seen, that the SnO2 particles are not completely coated by PtRh particles, therefore areas with pure tin are visible (blue in Fig. 3). The Pt:Rh ratio determined by EDX in this sample is 3:1. The right image in Fig. 3 shows the HRTEM image of the PtRh/SnO2 particles, which are visible as dark dots on the amorphous, circular Vulcan carbon substrate.

[1] V.S. Bagotsky, Fuel Cells: Problems and Solutions 2nd Ed.; John Wiley & Sons: Hoboken, New Jersey, 2012 pp. 3-5
[2] A. Kowal, M. Li, M. Shao, K. Sasaki, M.B. Vukmirovic, J. Zhang, N. S. Marinkovic, P. Liu, A. Frenkel, R. R. Adzic, Nature Materials, 2009, 9, 325-330.
[3] A. Kowal, S.Lj. Gojković, K.-S. Lee, P. Olszewski, Y.-E. Sung, Electrochem. Comm. 2009, 11, 724-727.
[4] R. Adzic, A. Kowal, (Brookhaven National Laboratory), Patent Application Publication, Pub. No. US2009/0068505 A1 (Mar. 12, 2009).
[5] M. Li, A. Kowal, K. Sasaki, N. Marinkovic, D. Su, E. Korach, P. Liu, R. Adzic, Electrochima Acta, 2010, 55, 4331-4338.
[6] M. Parlinska-Wojtan, R. Sowa, M. Pokora, A. Martyła, K. S. Lee and A. Kowal, Surf. & Interfaces Anal. Published on-line 15 Feb. 2014, DOI: 10.1002/sia.5384


Fig. 1: HAADF STEM images of the PtRh/SnO2 catalyst deposited on the Vulcan carbon substrate: (a) overview image showing a uniform distribution of the catalyst nanoparticles; (b) magnified view – the PtRh and tin oxide particles are not directly distinguishable in the HAADF detector.

Fig. 2: (left to right) STEM HAADF image of the PtRh/SnO2 particles and the corresponding quantified EDX maps of Pt, Rh and Sn.

Fig. 3: left image: EDX map of the catalyst showing the distribution of Pt, Rh and Sn in the sample; right image: HRTEM image of the PtRh/SnO2 (dark dots) particles on amorphous carbon.

Type of presentation: Poster

MS-1-P-5798 Ozone decomposition over supported Ni/Pd catalysts synthesized by extractive-pyrolytic method

Batakliev T.1, Georgiev V.1, Serga V.2, Anachkov M.1, Rakovsky S.1
1Institute of Catalysis, Bulgarian Academy of Sciences, 1113 Sofia, Bulgaria, 2Riga Technical University, Institute of Inorganic Chemistry, 2169 Riga, Latvia
todor@ic.bas.bg

The full text of the abstract is not available. Please contact the presenting author.


Type of presentation: Poster

MS-1-P-5805 Silver coated perlite catalyst for ground-level ozone degradation

Georgiev V.1, Blaskov V.2, Stambolova I.2, Batakliev T.1, Shipochka M.2, Vassilev S.2, Anachkov M.1, Eliyas A.1, Rakovsky S.1
1Institute of Catalysis, Sofia, Bulgaria 1, 2Institute of General and Inorganic Chemistry, Sofia, Bulgaria
vlado@ic.bas.bg

The full text of the abstract is not available. Please contact the presenting author.


Type of presentation: Poster

MS-1-P-5831 Benefit of HRTEM to explain the catalytic behavior of gold catalysts on Co- and Fe-doped ceria supports

Petrova P.1, Zanella R.2, Ivanov I.1, Pantaleo G.3, Venezia A M.3, Ilieva L.1
1Institute of Catalysis, Bulgarian Academy of Sciences, Sofia, Bulgaria, 2Centro de Ciencias Aplicadas y Desarrollo Tecnológico Universidad Nacional Autónoma de México, 3Istituto per lo Studio di Materiali Nanostrutturati, CNR, I- 90146 Palermo, Italy
petia@ic.bas.bg

The dispersion of gold is one of the key factors for the high catalytic performance of gold based catalysts. Nanosized gold catalysts supported on Co- and Fe-modified ceria were studied in important for environmental point of view processes: i) complete benzene oxidation (CBO) over gold (3wt%) catalysts on Co-doped ceria supports (5, 10 and 15 wt% Co3O4) prepared by mechanochemical mixing (MM); ii) CO-free hydrogen production for fuel cells application via WGS and PROX over gold (3wt%) catalysts on Fe-doped ceria supports (5, 10 and 20 wt% Fe2O3) prepared by mechanochemical mixing (MM) or impregnation (IM). The catalysts were characterized by different methods (XRD, XPS, TPR). Gold dispersion was evaluated by means of high resolution transmission electron microscopy (HRTEM) and high angle annular dark field (HAADF) measurements. The study is focused on the relationship between gold dispersion and catalytic activity, depending on the method of supports preparation and the dopant amount.


Very high catalytic activity in CBO was observed over gold catalyst on MM prepared ceria with 10 wt% Co3O4. It was significantly higher compared to 5 wt% or 15 wt% dopant. The HRTEM/HAADF results revealed that the doping with 10 wt.% Co3O4 was favorable for the highest gold dispersion: the highest part of very small particles (0.5 nm) and the lowest amount of bigger particles (3.5 nm and above) can be seen for Au10CoCeMM catalyst (fig. 1). It correlates with the highest reducibility and the highest oxidation activity in CBO.

The MM or IM preparation methods of Fe-doped ceria supports of gold catalysts affected in different way the catalytic behavior in the WGS and PROX reactions. Gold catalysts on IM supports exhibited WGS activity lower than that of gold/ceria. Significantly better WGS performance was demonstrated using MM. The observed differences were explained mainly by the differences in gold dispersion determined by the preparation method (considering the features of the multicomponent supports as well): HRTEM/HAADF results in accordance with XRD and XPS data showed higher gold dispersion in the case of supports prepared by MM. The gold dispersion is not such a key factor for PROX because the variations in the CO conversion as well as the selectivity dependence on the preparation method were not very substantial.


Important role for PROX at realistic conditions with CO2 and water in the gas feed played the formed using IM method nanosized hematite particles. This Fe-phase was evidenced by detailed analysis of HRTEM images in agreement with the Mössbauer results at LNT. Covering the ceria grains it leads to lower surface basicity and by this way they could improve the resistance toward CO2 deactivation.


This work was supported by ESF (Grant BG051PO001-3.3.06-0050).

Fig. 1: Catalytic activity, HRTEM/HAADF images and size distribution histograms of gold paricles for the Au5CoCeMM, Au10CoCeMM and Au15CoCeMM catalysts

Type of presentation: Poster

MS-1-P-5833 Combining electrons and ions: a (S)TEM-FIB microscopy application to unveil the final surface structure of a washcoated monolithic catalyst

Hernández-Garrido J. C.1, Gaona D.1, Gómez D. M.1, Vidal H.1, Gatica J. M.1, Sanz O.2, Rebled J. M.3, Peiró F.4, Calvino J. J.1
1Universidad de Cádiz, Cádiz, Spain, 2Universidad del País Vasco, San Sebastián, Spain, 3Institut de Ciència de Materials - CSIC, Barcelona, Spain, 4Universitat de Barcelona, Barcelona, Spain
jcarlos.hernandez@uca.es

The use of structured materials consisting of catalyst-coated honeycomb-type monoliths offers some advantages in comparison with the powder catalysts: e.g. higher exposed surface area and/or improvement of the active site-reactant contact. The dip-coating is the most extended procedure to load the powdered catalyst onto the monolith. Although the structural and chemical characterization at the nanometer scale of powder catalysts by (scanning) transmission electron microscopy, (S)TEM, techniques is in most cases well established, the characterization at such scale of catalytic devices, as it is the case of coated monoliths, poses currently a few challenges, some of them related to basic strategies to obtain representative information from the observations and others related to the sample preparation steps. In any case the approach is mandatory to determine the influence of the preparation procedure used to load the monolith on the final structure of both the active catalyst powder and of the coating itself.

There are several techniques to characterize devices at the micrometer scale but they present some disadvantages related to get a sample without mechanical damages. More recently, Focused ion beam (FIB) offers a solution of this particular (S)TEM characterization challenge, because it allows extracting precisely positioned, nanometer-sized sections of this type of devices, suitable for high resolution studies by different (S)TEM techniques. In this contribution we report how the combination of scanning electron microcopy (SEM)+(S)TEM + FIB studies yields valuable information from specific areas of a Co3O4/La-modified-CeO2 catalyst deposited by means of washcoat techniques on honeycomb cordierite (400 cells/in2) and Fe-C-based monoliths, designed for catalytic combustion of volatile organic compounds.

X-EDS maps in STEM-mode provided detailed information about the spatial distribution of these components. Remarkable, a Co3O4 layer was detected the external surface of the washcoat surrounding an inner core made up of an ensemble of the nanosized CeO2 support crystallites, suggesting that during the washcoating process the two components of the initial active catalysts, a ceria-supported cobalt oxide, segregate in space giving rise to a stratified structure in the coating layer of the catalytic device, far too different from the initial powder in which the two components were intimately mixed. These findings were confirmed in both cases, ceramic or metallic monolithic-supports. We remark the potential of the combined FIB-STEM characterization to give detailed important information about the surface of the monolith, which it is not possible to obtain by macroscopic characterization techniques and even by Electron Microscopy techniques.


Fig. 1: Catalyst-coated ceramic (left) and metallic (right) honeycomb-type monoliths.

Fig. 2: SEM images recorded at the different steps of the preparation of the FIB sample of the monolith washcoated. Adapted from Hernandez-Garrido et al. J Phys Chem C, 117(25):13028. Adapted with permission.

Fig. 3: X-EDS element (Co, Ce and Al) distribution maps recorded on the catalysts powder (a), the waschcoating suspension (b) and the final FIB sample from the ceraminc monolith. Adapted from Hernandez-Garrido et al. J Phys Chem C, 117(25):13028. Adapted with permission.

Type of presentation: Poster

MS-1-P-5885 TiO2 nanoparticles obtained by laser ablation in water: influence of pulse energy and duration on the crystalline phase

Canton P.1, Giorgetti E.2, Marsili P.2, Muniz-Miranda M.3, Vergari C.4, Giammanco F.4
1Department of Molecular Sciences and Nanosystems, University of Venezia, Mestre, 30170, Italy cantonpa@unive.it, 2Istituto dei Sistemi Complessi - CNR, Via Madonna del Piano 10, Sesto Fiorentino (Firenze), Italy, 3Department of Chemistry “U. Schiff”, University of Firenze, Firenze, Italy , 4Department of Physics "E. Fermi", University of Pisa, Pisa, Italy.
cantonpa@unive.it

Pulsed laser ablation (PLAL) can be used for production of stable and unprotected TiO2 nanoparticles (NPs) in pure solvents [1]. In general, rutile or anatase phase are required, depending on the applications, the first one being more attractive in the development of pigments and the second one more appropriate in photocatalysis. However, in spite of the large amount of work described in the literature, the control of the crystalline phase of the obtained samples is still a challenging tasks.
For this purpose, we performed a thorough characterization of the ablation of a Ti target in deionized water, by using the 1064 nm fundamental wavelength of a ns or a ps Nd:YAG laser and by tuning the energy per pulse and the fluence on target. We analyzed the colloids by UV-vis and Raman spectroscopy and by SAED, NBD, HRTEM and found some experimental rules which allow the control of the different phases of the oxide. According to Raman tests, we obtained the characteristic bands at 440 and 605 cm-1 of rutile NPs with ns pulses and prevalently rutile or anatase NPs with ps pulses, being anatase (395, 506 and 625 cm-1) more abundant when ps ablation is carried out with high energy pulses (see Fig. 1, 2 showing BF and NBD of three nanoparticles evidencing crystalline and amorphous structures).
The previous experimental results were compared with a theoretical model, which gives a detailed description of the ablation process during the laser pulse and the subsequent time-space evolution of key parameters of yields, i.e. pressure and temperature, in the surrounding solvent. By using simplified rate equations and phase diagrams of Ti oxides, the model not only allows to explain the observed energy and pulse-width dependence of TiO2 crystalline phase, but also provides a guide to choose the experimental parameters required to isolate the different crystalline species .

References.
[1] J. S. Golightly and A. W. Castleman, Jr.; J. Phys. Chem. B 110 (2006) pp. 19979-19984


Fig. 1: BF of 8mJ ps ablation specimen

Fig. 2: NBD of the three particles shown in figure 1

Type of presentation: Poster

MS-1-P-5946 Self-assembly of nanoparticles into 3D supraparticles

Kister T.1, Kraus T.1
1INM – Leibniz Institute for New Materials, Saarbruecken, Germany
thomas.kister@inm-gmbh.de

In the last decade, nanoparticles (NPs) have been applied in optics, magnetics and electronics. For example, fluorescent semiconductor NPs and metal NPs are used as markers for cells. Uniform NPs have been shown to self-assemble into regular structure like 2D or 3D superlattices. Depending on conditions such as solvent and temperature, the resulting superlattices can have different crystalline structures [1]. Supraparticles (SP) are self-assembled 3D clusters of NPs which are stably dispersed in a solvent. They can be produced using oil-in-water emulsions. Nanoparticles are confined inside the dispersed nonpolar phase. Upon evaporation of the oil phase, NPs arrange into SPs [2]. Their exact arrangement depends strongly on the surfactant that stabilizes the emulsion [3]. The SP retain most properties of the NPs; coupling between the closely packed NPs leads to plasmon shifts and energy transfer. SPs can also be formed from dispersions containing different NPs. The resulting SPs combine the properties of its constituents. For example, binary SP from gold and cadmium selenide (CdSe) NPs exhibit optical plasmon absorption due to the gold NPs and fluorescence due to the CdSe quantum dots.
Electron microscopy is the only available technique that provides sufficient resolution to study shape and structure of the SPs. It can resolve the particles’ arrangement inside the SP. SPs containing two types of NPs may be binary crystals, random mixtures of the constituent or Janus-like structures. Electron microscopy resolves such differences. Figure 1 shows a TEM picture of a SP produced from gold NPs. The structure of the SPs is similar to that of minimum energy particle arrangements known as Lennard–Jones clusters. We found that SPs from monodisperse gold or silver NP tend to exhibit such cluster-like structures. CdSe NPs and mixtures of CdSe and gold NPs exclusively assembled into SPs with glassy structures (Fig. 2).
So far, the structure of SPs was estimated from their 2D projections obtained from TEM. Currently, electron tomography measurements are performed to reconstruct precise 3D NP arrangements. Projections of a single SP are recorded at different angles in the range of -45° to 45°. We employ the “BART” and the “SIRT” to reconstruct a 3D model in order to quantitatively characterize the inner part of SPs. The main technical difficulty is posed by the dense gold NP cores. Under certain angles, the low-density spacing between the cores is sufficient for electrons to penetrate the entire SP. We will exploit this property to improve reconstruction.

1. E. Shevchenko et al. J. Am. Chem. Soc., 2006, 128, 3620−3637
2. J. Lacava et al. Nano Letters, 2012. 12(6), 3279-3282
3. J. Lacava et al. Soft Matter, 2014, 10, 1696-1704


Fig. 1: Ordered supraparticles from gold nanoparticles

Fig. 2: Binary supraparticles from gold and cadmium selenide

Type of presentation: Poster

MS-1-P-5995 SEM, HRTEM and HAADF analysis of nickel-doped ceria nanorods obtained by hydrothermal method

Romero-Núñez A.1, Hernández-Cristobal O.1, Díaz G.1
1Instituto de Física, Universidad Nacional Autónoma de México, México
araromero@fisica.unam.mx

A series of NixCe(1-x)O2 nanorods with different nickel contents were synthesized via a simple hydrothermal method. The aim of this work is to simultaneously control the composition and morphology of cubic ceria (CeO2) structure. These concepts are important defining catalytic properties of CeO2 or any other material [1]. Even though it is already known that composition and morphology are critical to improve catalytic properties a simultaneous control over such factors has been barely approached. Tuning the morphology into one-dimensional shapes leads a preferential exposure of reactive facets which improve catalytic performance [2]. SEM and TEM images of a NixCe(1-x)O2 nanorods sample are shown in figure 1a and 1b respectively. A one dimensional rod-like morphology, which is expected to exhibit {110} and {110} reactive ceria planes, is directly observed. In figure 1c HRTEM analysis of the same sample shows a [110] direction growth that corroborates the preferential exposure of {110} and {110} surface planes. Figure 2 show the STEM analysis of NixCe(1-x)O2 nanorods. HAADF image, figure 2a, also confirms the rod-like morphology. Cerium and nickel EDS elemental mapping, figure 2a and 2b, show a dispersed and homogeneous distribution of Ni species in the ceria host structure. Ni species distribution is a critical factor for catalytic properties as selectivity, activity and stability [3]. Catalytic performance for CO oxidization is superior in the doped NixCe(1-x)O2 nanorods samples than in undoped ceria nanorods. This is in agreement with the extrinsic formation of defects, which is inherent of the formation of the solid solution, and with the high dispersion of Ni that was corroborated by EDS mapping analysis. Nickel-doping and one-dimensional morphology are tuned together for the first time on ceria structure showing good catalytic properties. A full understanding of catalytic performance could only achieved with the careful structure analysis provided by microscopic techniques. EELS analysis is currently in progress and will be also presented to complement the present work.

[1] Li T, Xiang G, Zhuang J, Wang X. Chem Commun 2011;47:6060–2.
[2] Nolan M, Parker SC, Watson GW. Surf Sci 2005;595:223–32.
[3] Barrio L, Kubacka A, Zhou G, et al. J Phys Chem C 2010;114:12689–97.


Authors want to thank financial support of CONACyT-176509, PAPIIT-IN107512, PAEP-PCeIM and CONACyT-BN-332648, and technical support of Antonio Gómez-Cortés, Roberto Hernández, Manuel Aguilar and Antonio Morales.

Fig. 1: Electron microscopy characterization of NixCe(1-x)O2-NR. 1D rod-like morphology of samples is confirmed by (a) SEM and (b) TEM images. (c) HRTEM image view along [110] in which interplanar distances and angles corroborate the [110] growth direction.

Fig. 2: STEM analysis of NixCe(1-x)O2 nanorods (a) HAADF image, EDS spectroscopic mapping of (b) Cerium and (c) Nickel.

Type of presentation: Poster

MS-1-P-6007 Microstructure of Pt nanoparticles deposited on stainless steel in simulated boiling water reactor environment

Veleva L. V.1, Grundler P. V.1, Ramar A.2, Ritter S.1
1Paul Scherrer Institut, Nuclear Energy and Safety Research Department, 5232 Villigen PSI, Switzerland, 2GE-Global Research, EPIP zone, Bangalore-560037, INDIA
lyubomira.veleva@psi.ch

In boiling water reactors (BWR) radiolysis products of water (O2/H2O2) generate a highly oxidising environment which may result in an increased susceptibility to stress corrosion cracking (SCC) of reactor components. The technology of online noble metal chemical addition (OLNC) was developed by General Electric, where noble metal compounds are injected into reactor feed water of a BWR to mitigate SCC on internals and recirculation systems. Upon injection into the hot water (220-288°C), Na2Pt (OH)6 decomposes to form Pt nanoparticles. These nanoparticles deposit on all water wetted reactor components, and in presence of H2, catalyse the reduction of O2/H2O2, thus decreasing the electrochemical corrosion potential.

In order to assess this SCC mitigation technique, the Pt particle distribution and deposition behaviour on stainless steel coupon specimens, exposed to simulated BWR water conditions, using a sophisticated high-temperature water loop, has been investigated in detail at PSI. The Pt treated stainless steel coupons were first studied by field emission gun SEM. To better understand the catalytic behaviour and the bonding of the Pt particles to the oxide film, the microstructure and morphology of single Pt nanoparticles was studied by STEM and TEM techniques, including EDS. TEM specimens were prepared by using a replica technique, allowing the removal of Pt particles together with some of oxide crystals from the outer part of the oxide layer.

Preliminary results have shown that the Pt particles are homogeneously distributed on the surface of the oxide layer, with sizes in the nanometric range (Figure 1). It has been observed that the Pt particles precipitate in various locations on the oxide surface, including different crystals, edges, and facets (Figure 2). STEM and TEM high resolution observations confirmed that the Pt particles (Figures 3 and 4), have a crystalline structure and different shapes: round shape without or with facets (Figure 3), or rhomboidal shape (Figure 4). Edges, steps, corners and twin boundaries are often sites of high catalytic activity; therefore particles rich in such features are likely to be more efficient catalysts.

The difference in the observed Pt particles including shape, size and orientation to the oxide matrix could indicate various mechanisms involved in the nucleation of the Pt precipitates, as well as their binding to the oxide surface.


The financial support by ENSI, the contributions of the nuclear power plants KKL, KKM, and the microscopy centre of ETHZ to this work, are gratefully acknowledged.

Fig. 1: SEM BSE image of Pt nanoparticles precipitates on the oxide surface

Fig. 2: STEM HAADF Z-contrast image of Pt nanoparticles on an oxide crystal, occupying different facets, edges and corners.

Fig. 3: STEM bright field image of two Pt nanoparticles, one with facets oriented in [111] direction and the second one with a round shape and no visible facets.

Fig. 4: TEM bright filed image of a Pt nanoparticle with a rhomboidal shape, attached to the oxide matrix and oriented in [111] direction.

Type of presentation: Poster

MS-1-P-6011 Electrochemically deposited CeO2-NiO powders for solid oxide fuel cells: an in situ transmission electron microscopy study

Catalano M.1, Taurino A.1, Zhu J.2, Crozier P. A.3, Mele C.4, Bozzini B.4
1IMM-CNR, Via Monteroni, 73100 Lecce, Italy, 2Center for Solid State Science, ASU, Tempe, AZ 85287, USA, 3School for Engineering of Matter, Transport and Energy, ASU, Tempe, AZ 85287-6106, 4Dip. Ing. Innovazione, Università Salento, via Monteroni, 73100 Lecce, Italy
massimo.catalano@le.imm.cnr.it

Chemical and structural modifications of electrochemically deposited cermet precursor of Ni/NiO/CeO2, for the realization of solid oxides fuel cells (SOFC), were studied by conventional and in situ transmission electron microscopy (TEM). These cells are devices for effective conversion of chemical energy into electricity and heat, with low environmental impact. Cermets offer high ionic and electronic conductivity and high reforming and electrocatalytic activity. There is considerable interest in lowering the operating temperature of such devices, and doped cerias represent one possible materials choice. Doping ceria with oxides of lanthanides (Dy and Tb) improves the ionic and electronic conductivity and also increases the electrocatalytic activity of cermets [1-2]. In this work, the following CeO2-NiO samples were studied: (i) pure, (ii) Dy-doped, (iii) Tb-doped and (iv) co-doped with Dy and Tb. High spatial resolution in situ TEM techniques were used to monitor the changes in chemical/physical properties of these materials at the nanoscale [3], during thermal treatments employed to fabricate active cermets. High spatial resolution TEM observations were performed in a JEOL 2010F, an aberration corrected ARM200F and an environmental FEI Tecnai F20 transmission electron microscopes, with the combined use of TEM analytical techniques. The materials were initially analyzed in their as deposited form, then after ex situ heat treatments at 600°C for 1 hour in a furnace. After assessing the general modifications of the materials, they were subjected to in situ cycles of aging (temperature up to 700°C for maximum 300 minutes in an oxygen atmosphere) and the changes of their structural properties were monitored. Fig. 1(a) shows a typical high resolution image obtained from the pure and untreated sample, consisting of CeO2 nanograins (size < 5nm) apparently laying over a polycrystalline NiO layer, with much larger grain size. The insets show the Fast Fourier Transform (FFT) from the areas marked in the figure. CeO2 in the cubic phase was detected; evidence for the rhombohedral NiO phase was found. Upon ex situ heat treatment, coarsening of the grains occurs, the CeO2 ones reaching sizes up to 10-20 nm. Large NiO crystals can be observed, as shown in figure 1(b). Doping results in an amorphization of the samples. Fig. 1 shows the images from the co-doped sample, before (c) and after (d) ex situ annealing. The images of the same sample after the in situ treatment are shown ((e) and (f)). The details of the ex situ and in situ treatments will be compared and discussed.

[1] B. Bozzini et al., Electrochem. Commun. 24 (2012) 104

[2] C. Mele and B. Bozzini, Energies 5 (2012) 5363

[3] R. Wang , P. A. Crozier and R. Sharma, J. Phys. Chem. C 113 (2009) 5700


We acknowledge support from US NSF DMR-1308085 and the use of the electron microscope at the LeRoy Eyring Center for Solid State Science at ASU.

Fig. 1: High resolution images of: pure untreated a) and treated b) samples and FFTs from marked areas, with zone axis identification and attribution to the CeO2 (1, 2….labels) and NiO (1’, 2’….labels); co-doped untreated c) and treated d) samples and relevant diffraction patterns; co-doped sample untreated e) and in situ treated d) at 700°C for 170 min

Type of presentation: Poster

MS-1-P-6019 The influence of oxygen on the metal seeded growth of SnO2 nanowires

Krekeler T.1, Mader W.2
1Technische Universität Hamburg-Harburg, Hamburg, Germany, 2Rheinsche Friedrich-Wilhelms-Universität Bonn, Bonn, Germany
krekeler@uni-bonn.de

Tin dioxide (SnO2) is the most common sensor material for detection of reducing gases like CO, H2, CH4 etc. [1]. Many contributions about gas sensors based on SnO2 nanowires have been published over the last decade, showing the great potential of nanostructured sensor materials [2]. One versatile method for the synthesis of nanowires is the well investigated vapor-liquid-solid (VLS) mechanism with gold particles as a catalyst [3]. The exact growth mechanism of metal oxide nanowires is however still a matter of discussion because of the insolubility of oxygen in gold.

In this work we describe the synthesis of SnO2 nanowires by MOCVD technique and determine the influence of oxygen on the nanowire growth by methods of electron microscopy. The nanowires are grown via a reaction of TMT (Sn(CH3)4) with oxygen (O2) on fused silica substrates with gold particles as a seed. Synthesis using optimized reaction parameters (t = 10 min, T = 800 °C, p = 1 Pa) and a TMT:O2 molar ratio of 1:35 yields SnO2 nanowires of 3 µm length and 30 nm width. The nanowires are terminated by a facetted particle with corresponding width (Fig. 1a). The growth direction was determined to be <101> of the cassiterite modification of SnO2 from HRTEM (Fig. 1b-d). EDS measurements show that the particles are tin-free gold particles (Fig. 2).
If the TMT:O2 ratio is raised to 1:1.3 by keeping the TMT flow rate constant and reducing the oxygen flow rate, the growth speed strongly increases by a factor of 15. The terminating particles are now of a round shape and about twice the diameter of the nanowires (Fig. 3). EDS measurements show significant amounts of tin in the gold particle, indicating the formation of a Sn-Au alloy.

The relationship between growth speed and TMT:O2 ratio is supposed to be tied to the state of aggregation of the catalytic particle. A high TMT:O2 ratio allows for accumulation of tin in the particle, leading to a liquefied particle and therefore higher surface diffusion rates to the particle-nanowire interface resulting in faster growth. A low TMT:O2 ratio averts the accumulation of tin in the catalytic particle and liquefaction of the particle. This leads to lower diffusion rates of the reactants and thus slower growth of the nanowire.

These results clearly demonstrate that, in contrast to the classic VLS-mechanism, growth of SnO2 nanowires can occur without liquefaction of the catalyst particle. Therefore the growth mechanism of metal oxide nanowires can be better described as a surface mediated "metal seeded growth" [4].

1. J. Watson, Sens. Actuators (1984), 5, 29-42.
2. B. Wang, J. Phys. Chem. C (2008), 12, 6643-6647.
3. R.S. Wagner, W.C. Ellis, Appl. Phys. Lett. (1964), 4, 89.
4. W. Mader, Cryst. Growth Des. (2013), 13, 572−580.


All research leading to this contribution has been done at the Rheinische Friedrich-Wilhelms-Universität Bonn.

Fig. 1: (a) TEM BF-image of SnO2 nanowire, (b) closeup of the interface between catalyst particle and nanowire, (c) closeup of the tetragonal packed tin cations in SnO2, (d) FT of c. Zone axis determined to [111], growth axis to <101>.

Fig. 2: TEM BF-image and EDS-Spectrum of facetted Au particle on SnO2 nanowire. Red asterisk marks origin of EDS-Spectrum.

Fig. 3: TEM BF-image and EDS-Spectrum of round shaped Sn-Au particle on SnO2 nanowire. Red asterisk marks origin of EDS-Spectrum.

Type of presentation: Poster

MS-1-P-6029 Synthesis and Heterostructures of Metal Dichalcogenides Monolayer

Lee Y.1
1National Tsing Hua University
yhlee.mse@mx.nthu.edu.tw

Recently, monolayers of layered transition metal dichalcogenides (TMDc), such as MX2 (M = Mo, W and X = S, Se, Te), have been reported to exhibit excellent optoelectronic performances and diverse interesting properties. Monolayers in this class of materials offered a burgeoning field in fundamental physics, energy harvesting, electronics and optoelectronics. However, growth mechanisms and transfer of CVD-TMD monolayers remain challenge issues.[1~3] Hence, a feasible synthetic process and transfer techniques to overcome the challenges are essential. Here, we demonstrate the growth of high-quality TMD monolayers using chemical vapor deposition (CVD) with seeding promoter of aromatic molecules. The growth of monolayer TMD single crystals is achieved on various surfaces and a possible growth mechanism of the seed-activated growth would be presented.

We would like to demonstrate some techniques in transferring the TMD monolayers to diverse surfaces, Some characterization techniques and applications of vdw heterostructures were presented, which may which may stimulate the progress on diverse hybrid structures with TMDc monolayers.

Reference

[1] Yi-Hsien Lee, et al., Adv. Mater., 24 (17), p.2320-2325 (2012)

[2] Yi-Hsien Lee, et al. Nano Lett., 13 (4), 1852–1857 (2013)

[3] Xi-Ling, Yi-Hsien Lee*, et al., Nano Lett., 14 (2), p.464–472 (2014)

[4] Lili Yu, Yi-Hsien Lee, X. Ling, E. Santos, Y.C. Shin, Y. Lin, M. Dubey, E. Kaxiras, J. Kong, H. Wang, T. Palacios, Nano Lett, 14 (6), p.3055-3063 (2014)

[5] Xin-Quan Zhang et al, (in preparation)


We thank the Ministry of Science and Technology of the Republic of China (103-2112-M-007-001-MY3) for partial support of this research.

MS-2. Carbon-based nanomaterials, nanotubes, fullerenes, graphenes

Type of presentation: Invited

MS-2-IN-2482 Irradiation-induced Modifications and Beam-driven Dynamics in Low-dimensional Materials

Meyer J. C.1, Eder F. R.1, Mangler C.1, Kaiser U.2, Kotakoski J.1
1University of Vienna, Physics department, Vienna, Austria, 2University of Ulm, Central Facility for Electron Microscopy, Ulm, Germany
jannik.meyer@univie.ac.at

Irradiation-induced phenomena open a plethora of pathways for material modifications far beyond the thermal equilibrium and which are beyond the reach of direct synthesis. By using electron beams, such modifications can be induced and simultaneously observed at the atomic level. Moreover, in the analysis of 2-D materials, the position of every atom (rather than the atomic column) can be discerned in a high-resolution image.

For example, the introduction of multi-vacancy defects in graphene can be readily observed under 100kV aberration-corrected HRTEM [1], and the combination of vacancy creation and bond rotations can be used to convert graphene into two-dimensional amorphous carbon [1,2]. Atom loss can be directly counted in the images as a function of dose and in this way we have measured the knock-on sputtering cross sections for carbon atoms in graphene [3]. Recently, we have carried out a statistical analysis of the 2-D amorphous carbon (or carbon glass) structures with variable degree of disorder [4], enabled by an automated image analysis. For the first time, this provides atomic configurations for a continuous transition from a crystalline to an amorphous state, which were used for a statistical analysis based on experimentally obtained atomic coordinates.

At lower energies, the formation of defects in the pristine lattice is less likely, but existing defects convert from one configuration to another and migrate under the beam. Fig. 1 shows results from a 60kV STEM experiment under ultra-high vacuum conditions where double vacancies in graphene are extraordinarily stable; meaning that they neither convert into higher-numbered vacancies nor trap carbon and convert back to a pristine lattice, for long sequences of images. Nevertheless, the defects rapidly move via beam-driven bond rotations [5].

Although beam-driven dynamics are useful to modify materials under direct observation, these effects are also a major obstacle for the analysis of the pristine state of the sample. Using simulated data, we have shown a new approach to extract information from very low dose exposures [6], which will be discussed in the second part of the presentation. If this can be achieved also with experimental data, it may provide a novel route to circumvent the limitations of radiation damage in the analysis of materials.

[1] J. Kotakoski et al., Phys. Rev. Lett. 106 (2011), p. 105505. [2] J. Kotakoski et al., Phys. Rev. B. 83 (2011), p. 245420. [3] J. C. Meyer et al., Phys. Rev. Lett. 108 (2012), p. 196102. [4] F. Eder et al., Scientific Reports, in press (2014).  [5] J. Kotakoski et al., submitted (2014). [6] J. C. Meyer et al., Ultramicroscopy in press (DOI: 10.1016/j.ultramic.2013.11.010)


Austrian Science Fund (FWF: P25721-N20, M1481-N20 and I1283-N20), European Research Council (ERC) project PICOMAT, German Ministry of Science (DFG), Ministry of Research and the Arts (MWK) of the State of Baden-Wuertternberg within the SALVE project, Computational time from the Vienna Scientific Cluster.

Fig. 1: (a-d) Four subsequent frames of a di-vacancy defect. The transition from (a) to (b) requires at least four bond rotations (indicated by arrows) while only one bond rotation is sufficient from (c) to (d). (e-h) partial STEM images of different di-vacancy configurations, indicating that a transformation has occurred during the scan.

Fig. 2: (a) Simulated STEM data for infinite dose, and (b) processed for 500 electrons per square Angstrom. (c) Maximum-likelihood reconstruction from a larger area of low-dose data [6].

Type of presentation: Invited

MS-2-IN-2939 Defect structure and dynamics in two dimensional crystals under in situ heating and biasing conditions

Alem N.1
1Materials Science and Engineering Department, Center for Two Dimensional and Layered Materials, Penn State University, University Park, PA, USA
nua10@psu.edu

The past decade has seen incredible progress in the ability to isolate and manipulate two dimensional (2D) crystals. Such crystals are made of a network of atoms with strong bonds in the crystal plane, and much weaker out-of-plane van der Waals bonds. Due to this unique structure and dimensionality, charge carriers can be confined in two dimensions resulting in peculiar physical, chemical, and electronic properties. Such unexplored properties can be controlled and tuned through defects, step edges, interfaces, and grain boundaries. In this study, ultra-high resolution aberration-corrected electron microscopy is used to investigate the chemical and atomic structure of the edges, defects and grain boundaries in atomically thin two dimensional crystals, i.e. graphene, hexagonal boron nitride (h-BN) and tungsten disulfide (WS2). In addition, we use ultra-high resolution TEM to probe the structure of defects, edges, grain boundaries and their stability and dynamics under in situ thermal and electrical conditions [1-2]. Using high resolution electron microscopy imaging coupled with exit wave reconstruction technique, we have observed reconstruction of the edges in graphene into an armchair structure and formation and growth of 5-5-8 line defects originating from the holes in a monolayer graphene under joule heating conditions [1]. Fig. 1 shows an example of a hole in graphene under such conditions and formation of a line defect originating from the hole. In contrast to graphene, we observe holes with a zigzag structure in a monolayer of hexagonal boron nitride (Fig. 2) under in situ heating conditions. Unusual defect structures, such as 5-7 defects, are formed in h-BN and along the grain boundaries at high temperatures, although their formation is not expected in hexagonal boron nitride [2]. This talk will also address the stability and migration dynamics of grain boundaries in a monolayer WS2 as opposed to defect dynamics in graphene [3].

References:

[1] J. H. Chen, G. Autès, N. Alem, F. Gargiulo, A. Gautam, M. Linck, C. Kisielowski, O. V. Yazyev, S. G. Louie and A. Zettl, Controlled Growth of a Line Defect in Graphene and Implications for Gate-Tunable Valley Filtering, PRB, In press.

[2] A. L. Gibb, N. Alem, J.H. Chen, J. K. Erickson, J. Ciston, A. Gautam, M. Linck, and A. Zettl, Atomic Resolution Transmission Electron Microscopy of Grain Boundaries in Chemical Vapor Deposition Hexagonal Boron Nitride, JACS, 135 (18), (2013) 6758–6761.

[3] A. Azizi, X. Zou, P. Ercius, Z. Zhang, A. L. Elías, N. Perea-López, M. Terrones, B. I. Yakobson, and N. Alem, Dislocation and Grain Boundary Migration in 2D Transition Metal Dichalcogenides, submitted.


This work was performed at the physics department at University of California Berkeley, as well as the Materials Science and Engineering Department and the Center for Two Dimensional and Layered Materials at Penn State University. TEAM microscope at the National Center for Electron Microscopy at Lawrence Berkeley National Lab was used for this study.

Fig. 1: HREM image of a hole in graphene film reconstructed to maintain armchair edges under joule heating conditions. A 5-5-8 line defect (shown with arrow) originating from the hole is formed and extended into the defect free region of the film. The atoms are white. Scale bar is 1 nm.

Fig. 2: HREM image of a monolayer of hexagonal boron nitride under in situ heating at 450 °C. Holes are mostly observed with zigzag structure, while mono-vacancies are considered the main defects at this temperature. The atoms are white. The scale bar is 1 nm.

Type of presentation: Invited

MS-2-IN-3128 TEM and cathodoluminescence with nanometric spatial resolution on BN nanostructures

Pierret A.1, 2, Schué L.1, 3, Fossard F.1, Moldovan S.4, Ersen O.4, Ducastelle F.1, Barjon J.3, Loiseau A.1
1LEM, ONERA-CNRS, Châtillon, France, 2CEA-CNRS-UJF group , 3GEMaC, UVSQ-CNRS, Versailles, France, 4IPCMS, Univ. Strasbourg-CNRS, Strasbourg, France
annick.loiseau@onera.fr

Hexagonal boron nitride (h-BN) is a wide band gap semiconductor (6.4eV), which can be synthesized, as its carbon analog graphite, as bulk crystallites, nanotubes and nanosheets. Investigation of their optoelectronic properties is made difficult because of the paucity of high quality samples and suitable investigation tools. These structures meet nevertheless a growing interest for deep UV LED and graphene engineering. A deeper understanding of the interplay between the structural and luminescence properties of different BN structures and how these properties can be further exploited for their characterization are therefore highly needed.

Such studies are now possible thanks to the recent development of dedicated photoluminescence (PL) and cathodoluminescence (CL) experiments running at 4K and adapted to the detection in the far UV range (up to 6eV) [1]. We can also combine various TEM techniques and CL experiments in a FEG-SEM with a spatial resolution of 3nm on the same nano-object. With these tools, we investigated the structure and luminescence of various structures, from high quality crystals [2], exfoliated nanosheets to multi-wall nanotubes [3].

As a result, BN materials present original optical properties, governed by excitonic effects in the 5.5–6eV energy range. Two kinds of excitonic luminescence have been identified and are called S and D lines [4]. As revealed from CL-TEM analyses, D lines are issued from defective areas (Fig 1), so that D/S ratio can be used as a qualification parameter of the defect densitiy [5]. This procedure has been applied to understand the first luminescence studies of few layers individual BN flakes [5].

Concerning nanotubes, CL images reveal that the luminescence in the 5.5–6eV energy range is strongly inhomogenous and oscillating. Thanks to a deep investigation combining different TEM techniques, we have shown that the tubes display a complex twisted faceted structure and that the twist period is correlated with the luminescence oscillations (Fig 2). Furthermore, we could show that excitons, responsible for the spectacular localization of the luminescence, are trapped to specific defects, twisted along with the faceting structure.

Finally, low-loss EELS providing an alternative approach to the nature of electronic excitations [6], we will show how it is an efficient tool to investigate the local structure and optical properties with an energy resolution below 100meV of different BN layers and nanotubes.

[1] P. Jaffrennou et al., PRB 77 (2008) 235422

[2] T. Taniguchi et al., J. Cryst. Growth 303 (2007) 525

[3] C. Tang et al., Chem. Commun. 12 (2002) 1290

[4] K. Watanabe et al., PRB 79 (2009) 193104

[5] A. Pierret et al., PRB 89 (2014) 035414

[6] R. Arenal et al., PRL 95 (2005) 127601


The research leading to these results has received funding from the European Union Seventh Framework Programme under grant agreement n°604391 Graphene Flagship. D. Golberg, T. Taniguchi and K. Watanabe from NIMS Japan are warmly acknowledged for providing samples.

Fig. 1: (a) SEM image of a h-BN crystallite; (b), (c) Corresponding CL images recorded (b) on the main S line (S3-S4), and (c) on the main D line (D4). (d) Map of the D/S ratio. (e) CL spectra recorded in the areas #1 (grain boundary) and #2 (middle of the grain), indicated in (a).

Fig. 2: Images of a BN nanotube: (a) 3nm spatially resolved CL image recorded at 5.49 eV (226 nm); (b), (c) Corresponding TEM images in (b) bright-field mode, and (c) dark-field mode on the (100) reflection. (d) Heptagonal tube cross-section obtained by tomography experiment. (e) Structure of the tube as deduced from (b-d) images.

Type of presentation: Oral

MS-2-O-1592 Illuminating the electronic properties of individual carbon nanotubes with photons and electrons

Rossouw D.1,3, Bugnet M.1, Najafi E.2, Lee V.1, Hitchcock A. P.1, Botton G. A.1
1McMaster University, Hamilton, Canada, 2California Institute of Technology, California, USA, 3Present address: University of Cambridge, Cambridge, United Kingdom
dr418@cam.ac.uk

Carbon nanotubes (CNTs) exhibit unique physicochemical properties that have led to their use in a variety of novel materials science applications. Despite rapid progress in the theoretical and experimental investigation of CNTs, techniques capable of studying the structural and electronic properties of individual tubes are limited. Here, the spectral signature of carbon is used to identify the electronic character of individual single-walled CNTs. In addition, a newly built laser-TEM system is used to study light-induced structural and electronic distortions in individual CNTs.
Using high-resolution EELS, we differentiate metallic and semiconducting SWCNTs based on the fine structure of the recorded carbon K edge [1]. While the overall features in the C-K edge are similar for metallic and semiconducting tubes, differences are observed in the fine structure of the π* peak between 284 and 286 eV (Fig. 1); semiconducting nanotubes have a shoulder to the left of the π* peak, metallic to the right. Results from scanning transmission X-ray microscopy performed on the same electronically pure SWCNTs are in good agreement with EELS and are of comparable spectral resolution. The quality of the EEL spectra of individual SWCNTs opens up the possibility to probe the electronic state of single-SWCNT devices.
The study of light driven electronic and structural changes in matter is fundamental to understanding materials properties and performance. While ultra-fast and time-resolved experiments provide unique information based on measurements from very short-time intervals, not much is known on the steady-state response of nanomaterials to an intense continuous beam of light [2]. To address this, a unique system has been built to deliver a focused and continuous laser spot coincident with the electron beam inside a TEM. The laser-TEM system allows the study of structural and electronic modifications in nanomaterials under intense light irradiation. Structural and electronic distortions in individual CNTs have been studied in-situ [3]. When illuminated, a multi-walled CNT expands radially with coupled changes in its σ* conduction band (Fig. 2), as well as in its π* plasmon spectral band. Such observations may aid our understanding of the unique photoconductivity and luminescence properties of CNTs.

[1] D. Rossouw, G.A Botton, E. Najafi, V. Lee, A.P. Hitchcock. ACS Nano, 6, 10965 (2012).
[2] A. Howie. The European Physical Journal - Applied Physics, 54, 33502 2011.
[3] D. Rossouw, M. Bugnet, and G.A. Botton. Physical Review B, 87, 125403 2013.


D.R. acknowledges support from the University of Cambridge and the Royal Society in the form of a Newton International Fellowship.

Fig. 1: Differentiating between individual metallic (a) and semiconducting (b) SWCNTs by their EELS carbon K edge (c).

Fig. 2: (a) The boxed region of a multi-walled CNT extending over a hole in the support grid selected for analysis (b) contains 12 tubules and is free of any obvious defects. (c) Changes in the C-K edge of the tube during laser illumination are strongest in the vicinity of the σ* peak and are fully reversible.

Type of presentation: Oral

MS-2-O-1827 Nanometric resolved cathodoluminescence on few layers h-BN flakes

Bourrellier R.1, Amato M.2, Meuret S.1, Tizei L.1, Giorgetti C.2, Gloter A.1, Heggie M.3, March K.1, Stephan O.1, Reining L.2, Kociak M.1, Zobelli A.1
1Laboratoire de Physique des Solides, Univ. Paris-Sud, CNRS UMR 8502, F-91405, Orsay, France, 2Laboratoire des Solides Irradies, Ecole Polytechnique, Route de Saclay, F-91128 Palaiseau and European Theoretical Spectroscopy Facility (ETSF), France, 3Department of Chemistry, University of Surrey, Guildford GU2 7XH, United Kingdom
romain.bourrellier@u-psud.fr

Within the latest years number of layered materials at reduced dimensions have demonstrated remarkable optical properties. However most studies focused on perfect system and the role of defects as optical active centers remain unexplored. Hexagonal boron nitride(h-BN) is one of the most promising candidates for light emitting devices in the far UV, presenting a single strong excitonic emission at 5.8 eV. However, a single line appears only in pure monocrystals, obtained through complex process[1]. Common h-BN samples present more complex emission spectra that have been attributed to the presence of structural defects. Despite a large number of experimental studies up to now it was not possible to attribute specific emission features to well identify defective structures.
Here we address this fundamental questions by adopting a theoretical and experimental approach combining few nanometer resolved Cathodoluminescence (CL) techniques with high resolution TEM images and state of the art quantum mechanical simulations.
Recently, the Orsay team has developed a CL detection system integrated within a STEM[2]. This unique experimental set up is now able to provide full emission spectra with a resolution as low as few tens of meV associated with an electron probe size of 1nm. A CL spectrum-image can thus be recorded in parallel with an HAADF image.Nanometric resolved CL on few-layer chemically exfoliated h-BN crystals have shown that emission spectra are inhomogeneous within individual flakes. Emission peaks close to the free exciton appear in extended regions. Complementary investigations through high resolution TEM allow to associate these emission lines with extended crystal deformation such as stacking faults and folds of the planes[3].
By means of ab-initio calculations in the framework of Many Body Perturbation Theory (GW+BSE) we provide an in-depth description of the electronic structure and spectroscopic response of bulk hexagonal boron nitride in the presence of extended morphological modifications. In particular we show that, in a good agreement with the experimental results, additional excitons are associated to local symmetry changes occurring at crystal stacking faults.
Additional features appearing within the band gap present a high spatial localization, typically less than 100 nm, and thus they can be related to individual point defects. When addressed individually through a highly focused electron probe they might have a single photon emitter quantum character. This hypothesis has been recently confirmed by experiments combining our CL system with an Hanbury Brown and Twiss interferometer.

[1] K. Watanabe et al, Nat. Mater. 3 (2004) p. 404
[2] L. Zagonel et al Nano Lett. 11 (2011) p. 56
[3] R Bourrellier et al, arXiv:cond-mat/1401.1948 (2014)


Fig. 1: a Bright field and b dark field images of an individual BN flake. c Overall emission spectrum of the flake andindividual spectra taken at specific probe positions indicated in panel b. d-h Emission maps for individual emission peak.Intensity is normalized independently within each individual map

Type of presentation: Oral

MS-2-O-1850 Atomic scale investigations of electron beam sensitive molecules embedded in carbon nanotubes

Biskupek J.1, Chamberlain T. W.2, Skowron S. T.2, Bayliss P. A.2, Bichoutskaja E.2, Khlobystov A. N.2, Kaiser U.2
11. University of Ulm, Central Facility of Electron Microscopy, Electron Microscopy Group of Materials Science, Albert Einstein Allee 11, D-89069 Ulm, Germany, 2School of Chemistry, University of Nottingham, University Park, Nottingham NG7 2RD, United Kingdom
johannes.biskupek@uni-ulm.de

Aberration corrected high resolution transmission electron microscopy (AC-HRTEM) at conventional accelerating voltages of 200 or 300 kV allows atomic structural investigations with sub-Ångstrøm point resolution. Specimens made of light atoms such as carbon nanostructures or Li-based materials are easily subjected to knock-on damage at 200 keV and higher energies. The nowadays state-of-the-art is to operate the CS-corrected microscopes at 80 keV energies to lower the damage of carbon nanostructures as demonstrated for graphene and carbon nanotubes below the knock-on threshold. Molecular organic structures including fullerenes (e.g. C60), tetrathiafulvalene (TTF, sulphur rich molecule) or coronene can be embedded into carbon nanotubes [1, 2] or sandwiched between graphene sheets [2] not to only act as the host “sample holder” but also to minimize charging effects. However, the typical energies of 80 keV used for AC-HRTEM investigations are eventually still too high and may damage and modify the delicate embedded molecules, especially if they contain hydrogen atoms in their structure. Therefore, the investigation of molecular organic structures in their pristine states is still a challenge.

We explore the possibilities of reduction of electron energies to 40 keV or even 20 keV to enhance the stability of molecules. The low electron energies require the correction of 5th order geometric aberrations and the correction of chromatic aberrations to compensate for the higher elctron wavelength.

We also study the advantages of dedicated low-dose techniques to minimize beam damage. Also modifications of the nanostructures prior to the actual imaging process are applied to enhance the stability against the electron irradiation.

Fig.1 shows an example HRTEM images of C60 molecules at different stages of electron irradiation. At 80 kV already a relatively small accumulated dose of 5×107 e-/nm2 is sufficient to form first dimers of C60 . However, no visible changes of the structure of C60 molecules are visible at 40 kV after irradiation with the same accumulated electron dose. A coalescence of the C60 molecules is clearly visible at a electron dose of 2×108 e-/nm2 and 80 kV irradiation. It requires almost two orders of magnitude higher dose (40 times) to initiate observable coalescence of the C60 molecules at 40 kV

[1] A. Chuvilin et al., Nature Materials 10 (2011) 687

[2] T. W. Chamberlain et al., ACS Nano 6, (2012) 3943

[3] G. Algara-Siller et al., APL 103 (2013) 203107


We acknowledge financial support from the German Research Foundation (DFG) and the Ministry of Science, Research and Arts (MWK) of the state Baden-Wurttemberg within the Sub-Ångstrøm Low-Voltage Electron Microscopy project (SALVE). We are grateful to P. Hartel and M. Linck (CEOS company) for assisting CC/CS corrected HRTEM experiments.

Fig. 1: HRTEM images of C60 molecules in single-walled carbon nanotubes imaged at 80 kV (left, CS-corrected) and 40 kV (right, CC/CS-corrected). At 40 kV it requires an almost two orders of magnitude increase in dose (40 times) to get first indications of coalescence and rupturing of the C60 molecules.

Type of presentation: Oral

MS-2-O-1977 In-situ electron mciroscopy of carbon atom chains

Banhart F.1, La Torre A.1, Cretu O.2
1Institut de Physique et Chimie des Matériaux, University of Strasbourg, 23 rue du Loess, 67034 Strasbourg, France, 2National Institute of Advanced Industrial Science and Technology, Central 5, 1-1-1 Higashi, Tsukuba, Ibaraki 305-8565, Japan
florian.banhart@ipcms.unistra.fr

Carbon chains are sp-hybridized strings of carbon atoms; they may may be considered as the elements of a one-dimensional phase of carbon. Atomic carbon chains have been proposed since a long time until they were observed by electron microscopy. According to theory, the chains may be bonded by either alternating single/triple carbon-carbon bonds (polyyne) or by double bonds throughout the chain (cumulene). Their electrical and mechanical properties and their stability have been subject of many theoretical studies; however, no experimental information has been available. Now, by using an STM stage (Nanofactory) in a TEM in an in-situ study, carbon atom chains have not only been made but also characterized electrically [1]. The chains were obtained by establishing a contact between a metallic tip and graphene ribbons. Retracting the tip while an electrical current flowed through the contact led to the unraveling of carbon atoms from the graphene ribbons. Figure 1 shows the simplified principle of the experiment and figure 2 the development of a typical carbon chain, spanning here between two graphene filaments. The electrical conductivity of the chains could be measured in such a way and was found to be much lower than predicted for ideal chains. Figure 3 shows the measured current-voltage characteristics of a chain. Theory predicts that strain in the chains determines their conductivity in a decisive way. Indeed, carbon chains are always under varying non-zero strain that transforms their atomic structure from cumulene to polyyne, thus inducing a tunable band gap. The modified electronic structure and the characteristics of the contact to the graphitic periphery explain the conductivity of the locally constrained carbon chains. New experiments show the local chemistry and the bonding at contacts between metals and carbon chains as well as characteristic current-voltage curves, depending on the type of contact. Dedicated experiments show qualitatively that the chains have an outstanding mechanical strength, in accordance with theory. The results show a perspective toward the synthesis of carbon chains and their application as the smallest possible interconnects or even as one-dimensional semiconducting devices.

[1]  O. Cretu, A. R. Botello-Mendez, I. Janowska, C. Pham-Huu, J.-C. Charlier and F. Banhart, Nano Lett. 13, 3487 (2013)


Funding by the French Agence Nationale de Recherche (projects NANOCONTACTS, NT09 507527 and NANOCELLS, ANR12 BS1000401) is gratefully acknowledged.

Fig. 1: Unraveling an atomic carbon chain from a graphene ribbon by passing a current through the junction and retracting the STM tip.

Fig. 2: Formation of a carbon chain between two graphene ribbons (FLG). The time scale as well as the length of the chain are indicated. In (f) the chain is broken.

Fig. 3: Carbon chain (arrowed) and its current-voltage characteristics.

Type of presentation: Oral

MS-2-O-2002 Measuring the temperature dependence of the Debye–Waller Factor in Graphene using electron diffraction techniques.

Allen C. S.1, Fan Y.1, Liberti E.1, Kim J.1, Warner J. H.1, Kirkland A. I.1
1Department of Materials, University of Oxford, OX1 3PH. UK.
christopher.allen@materials.ox.ac.uk

Within the thermodynamic limit lattice vibrations in a two dimensional (2D) crystal should destroy any long range order. As such prior to the isolation of monolayer graphene in 2004, 2D crystals were thought impossible to realise.[1] In order to explain the surprising stability of 2D crystals it is necessary to establish a detailed understanding of their lattice vibrations and how they might be affected by the properties of the crystal, for example domain size or defect density.
To date there have been many theoretical predictions of the phonon band structure of graphene with supporting experimental evidence from Raman spectroscopy measurements.[2] Recently the mean-square displacement (or Debye-Waller factor) of graphene atoms from their equilibrium lattice position has been measured from electron diffraction patterns.[3]
Using in-situ heating and cooling TEM holders we have recorded diffraction patterns from single crystals of mono-layer graphene at varying tilt angles and temperatures ranging from 100K – 1500K (figure 1). Careful analysis of these diffraction patterns has allowed us to extract values for the in-plane mean square displacement (Debye-Waller factor) of atoms over the whole temperature range. By studying the tilt dependence of the diffraction spot intensity we have also measured the out of plane atomic displacements relating to the flexural phonon modes of the graphene lattice.
We compare our results to theoretical predictions for the Debye-Waller factor based on calculations of the phonon-dispersion relation of graphene and comment on the validity of these models over the temperature range investigated.

[1] L.D. Landau et al. Statistical Physics . Part I (Butterworth-Heinemann, Amsterdam, 2003)
[2] E. Pop et al. MRS Bull. 37, 12, 1273-1281 (2012)
[3] B. Shevitski et al. Phys. Rev. B. 87,045417, (2013)


The research leading to these results has received funding from the European Union Seventh Framework Programme under Grant Agreement 312483 - ESTEEM2 (Integrated Infrastructure Initiative¬I3).
References

Fig. 1: Figure 1. a. Selected area diffraction pattern of mono-layer graphene taken at room temperature and zero tilt. b. Real space TEM image of the graphene sample imaged through the selected area aperture used for diffraction.

Type of presentation: Oral

MS-2-O-2010 XEDS-STEM tomography of multilayer nanotubes for 3D chemical characterization

Kurttepeli M.1, Bladt E.1, Deng S.2, Cott D. J.3, Detavernier C.2, Bals S.1, Van Tendeloo G.1
1EMAT, University of Antwerp, Groenenborgerlaan 171, B-2020 Antwerp, Belgium, 2Department of Solid State Science, Ghent University, Krijgslaan 281/S1, B-9000 Ghent, Belgium, 3Imec, 75, Kapeldreef, B-3001 Leuven, Belgium
mert.kurttepeli@uantwerpen.be

During the last decade, there has been an increasing demand on the 3D characterization of materials, which led to the development of different electron tomography techniques. BFTEM and HAADF-STEM based electron tomography are among those that are commonly performed in materials science. However, these techniques are mainly used to obtain the 3D morphologies of the nanostructures rather than the chemical information. Through the 3D composition mapping using STEM coupled with the X-ray energy dispersive spectrometry (XEDS) via symmetrically arranged XEDS detector design, it is now possible to resolve the 3D distribution of elements in nanoscale materials and to elucidate the 3D chemical information in a large field of view of the TEM sample [1].
We present the application of the XEDS-STEM tomography technique for 3D chemical imaging of nanoscale materials. We performed this technique to investigate the 3D chemical distribution of titanium dioxide (TiO2) and vanadium oxide (VOx) coated carbon nanotubes (CNT). Figure 1 and 2 show the results of 3D tomography applied to CNT-TiO2-VOx-TiO2 using both HAADF STEM and XEDS-STEM techniques. The comparison of simultaneously acquired HAADF-STEM and XEDS-STEM tomography results shows that XEDS-STEM tomography succeeds to provide 3D chemical information of the material in addition to the 3D morphology, in spite of the low, neighboring atomic numbers of Ti, V and C. As presented in Fig 2, XEDS-STEM tomography resolves the individual Ti, V and C containing layers and reveals that the coating of CNT by TiO2-VOx-TiO2 was uniform and conformal. One important advantage of XEDS-STEM tomography is the decreased electron beam induced damage in the TEM samples. With the improved XEDS detectors which give higher collection efficiencies, the specimen damage is minimized significantly [1]. This enabled the precise 3D nanoscale chemical characterization of fine structures as in our case without the shape and size changes of the object during acquisition. Through XEDS-STEM tomography technique it is therefore possible to resolve 3D compositional variations at nanoscale with high accuracy.

[1] A. Genc, L. Kovarik, M. Gu, H. Cheng, P. Plachinda, L. Pullan, B. Freitag and C. Wang, Ultramicroscopy 131, 24 - 32 (2013).


The authors acknowledge financial support from European Research Council and Sim-Flanders.

Fig. 1: Volume rendered 3D visualizations of HAADF-STEM reconstruction, revealing only the 3D morphology.

Fig. 2: Volume rendered composite 3D visualizations of XEDS reconstructions, revealing the 3D elemental distribution of Ti (assigned green), V (assigned red) and C (assigned blue).

Type of presentation: Oral

MS-2-O-2037 Local boron environment in B-doped diamond films studied by advanced TEM and spatially resolved EELS

Turner S.1, Lu Y.1, Idrissi H.1, Janssens S. D.2,3, Haenen K.2,3, Sartori A. F.4, Schreck M.4, Verbeeck J.1, Van Tendeloo G.1
1EMAT, University of Antwerp, Groenenborgerlaan 171, 2020 Antwerp, Belgium, 2Institute for Materials Research (IMO), Hasselt University, Wetenschapspark 1, B-3590 Diepenbeek, Belgium , 3IMOMEC, IMEC vzw, Wetenschapspark 1, B-3590 Diepenbeek, Belgium, 4Universität Augsburg, Institut für Physik, D-86135 Augsburg, Germany
stuart.turner@uantwerpen.be

Diamond is an attractive material for many technological applications, because of its extreme hardness, chemically inert surfaces, high Young’s modulus and large band gap of 5.5 eV. One of the most commonplace synthesis methods for nanocrystalline diamond (NCD) and epitaxial, single crystal thin films is microwave plasma assisted chemical vapour deposition (MPCVD). Many of the technological applications of diamond require specific semiconducting properties of the material and therefore doping is necessary. The most effective doping with p-type character is obtained by inserting boron in-situ during the growth process. Boron doping of diamond allows from mild p-type character for low [B] to a metallic regime and also superconducting properties at liquid helium temperatures for very high [B].1

However, much debate surrounds the question of the position and coordination of the B dopants in the diamond, especially in defective regions of the material. As B doping leads to an increase in defects in diamond grains and films upon growth, it is plausible that the boron dopants are preferentially embedded in defective regions.

In this work, conducting films of B-doped nanocrystalline diamond and single crystal diamond grown by MPCVD have been investigated in both plan-view and cross-section orientation by a combination of aberration-corrected (scanning) transmission electron microscopy (HR-ADF-STEM) and spatially resolved electron energy-loss spectroscopy (STEM-EELS) performed on a state-of-the-art aberration corrected instrument. Using these tools, the B concentration, distribution and the local B environment in this type of thin nanocrystalline diamond films have been determined.

Concentrations of ~1 at.% of boron are found to be embedded within the pristine diamond lattice. Boron distribution maps however clearly reveal a preferential enrichment of boron at defective areas like twin boundaries, incoherent defects and even dislocations in diamond thin films. Inspection of the EELS fine structure reveals a distinct difference in coordination of the B dopants in “pristine” diamond areas and in defective regions, identified through comparison of the experimental EELS fine structure to density functional theory (DFT) calculated fine structure signatures.2,3,4

1) Ekimov E.A. et al. (2004) Nature, 428, 542-545

2) Turner S. et al. (2012) Nanoscale, 4, 5960-5964

3) Lu Y.-G. et al. (2012) Applied Physics Letters, 101, 041907

4) Lu Y.-G. et al. (2013) Applied Physics Letters, 103, 032105


S.T. gratefully acknowledges financial support from the Fund for Scientific Research Flanders (FWO).

Fig. 1: B:NCD film. (a)&(b) Overview ADF-STEM images. (c) Image of a single defected diamond grain. (d)&(e) Survey image and quantitative B distribution map. B is clearly enriched at defects. (f) B-K edge fine structure from a diamond (black) and defect region (red). (g) C-K edge fine structure from a diamond (black) and defect region (red).

Fig. 2: Epitaxial B:diamond thin film. (a) ADF-STEM image of the thin film on a diamond substrate. A high density of dislocations is present in the film. (b)&(c) ADF image of a single dislocation and B map. B is enriched at the dislocations. (d) B-K edge from a pristine film region (black), a dislocation-rich region (red) and an etch pit (green).

Type of presentation: Oral

MS-2-O-2104 The making and electrical biasing of graphene nanoribbon-based devices inside the TEM

Rodríguez-Manzo J. A.1, Puster M.1, Qi Z. J.1, Balan A.1, Charlie Johnson A. T.1, Drndić M.1
1Department of Physics and Astronomy, University of Pennsylvania, Philadelphia, USA
rjulio@sas.upenn.edu

The band gap of graphene nanoribbons (GNR) makes them suitable candidates for sensor devices. Their dimensions (one-atom thickness and widths of tens of nanometers or less) make it difficult to experimentally correlate their electrical properties with changes in their width, edge structure or defect concentration. In this context, we discuss two examples in which GNR-based devices were characterized and also modified within a TEM with the electron beam while an electrical bias was applied.

In the first example, we describe how to fabricate GNR-nanopore devices, which are promising candidates for next-generation DNA sequencing, with the converged electron beam of a TEM [1]. Such devices normally comprise a 2-10 nm diameter pore formed with the beam at the edge or in the center of a 100 nm-wide GNR on a 50 nm-thick silicon nitride membrane. We discuss the changes on GNR conductance when such devices are irradiated with a 200 keV beam and the differences between irradiating with a homogenous (TEM mode) versus a scanned condensed beam (STEM mode). By minimizing the electron dose at 200 kV in STEM mode we were able to prevent electron beam-induced damage and make nanopores in highly conducting GNR. The resulting devices, with unchanged resistances after nanopore formation, can sustain micro ampere currents at low voltages (∼ 50 mV) in buffered electrolyte solution and exhibit high sensitivity, with a large relative change of resistance upon changes of gate voltage, similar to pristine GNR without nanopores (see Figure 1).

It is a truism that before characterizing GNR one must fabricate them, and this is a challenge by itself. In the second example, we describe how to use the condensed beam of a TEM with corrected spherical-aberrations to sputter carbon atoms from predefined areas in electrically-connected free-standing graphene sheets to obtain GNR with sub-10 nm widths [2]. This approach allows us to correlate the lattice and edge structure of sub-10 nm wide GNR with their electrical properties (see Figure 2).

These two examples illustrate the advantages of combining standard TEM observation of GNR-based devices with their electrical biasing as well as the challenges involved in this type of in situ TEM experiment, where chips fabricated with standard lithographic techniques are coupled to TEM sample holders through electrical contacts.

[1] Towards sensitive graphene nanoribbon-nanopore devices by preventing electron beam induced damage. M. Puster, J. A. Rodríguez-Manzo, A. Balan, M. Drndić, ACS Nano 7, 11283 (2013). [2] Correlating atomic structure and transport in suspended graphene nanoribbons. Z. J. Qi, J. A. Rodríguez-Manzo, Andrés R. Botello-Méndez, et al. Submitted for publication (2014).


This work was supported by NIH Grant R21HG004767 and NBIC through NSF NSEC DMR08-32802. We acknowledge use of TEM facilities at Pennsylvania and Rutgers Universities. JZQ and CJ acknowledge SRC contract #2011-IN-2229, NSF AIR Program ENG-1312202. Part of this work was done at the CFN in BNL, supported by the U.S. DOE, Contract No. DE-AC02-98CH10886 (FEI-Titan ACTEM through proposal 31972).

Fig. 1: (a) Chip-carrier and (b) detail of 200 um-wide SiNx window containing 4 GNR. (c) GNR-nanopore device. (d) HAADF STEM image of a nanopore next to a GNR. (e) HAADF STEM image of a 100 nm-wide GNR. Inset: positioning of the beam at the edge of the GNR with a precision of ~ 4 nm. (f) Resistance of a GNR during nanopore formation.

Fig. 2: (a) Sample holder with mounted chip and (b) detail of 500 nm-wide free-standing GNR. (c) Reduction of GNR width by carbon sputtering with the condensed electron beam. (d) HRTEM image of biased sub-10 nm GNR. (e) Resistance changes as a function of GNR width reduction.

Type of presentation: Oral

MS-2-O-2415 Evidence of re-crystallization in graphene probed by atomic scale imaging

Okuno H.1, Tyurnina A.2, Pochet P.1, Dijon J.2
1CEA-Grenoble, INAC/SP2M, Greoble, France, 2CEA-Greboble, LITEN/DTNM, Grenoble, France
hanako.okuno@cea.fr

Graphene shows great potential for future nanoelectronics due to its extraordinary electronic properties and structure-engineerable nature. Recently CVD based graphene production technology has given insights to the possibility of large scale application. Since the CVD grown films are typically polycrystalline with numerous grain boundaries (GBs), it is important to characterize and control grain size and GBs, generally believed to limit the transport properties of graphene film. In this work, we report on a peculiar grain evolution occurring during the growth process. This evolution is revealed by means of Raman spectroscopy that gives access to a statistically characterization of the graphene structure (grain size, presence of defects etc.). Besides, we need some other complementary techniques to fully understand the detailed atomic structures. This is done by means of atomic-resolution TEM. Indeed, following the invention of aberration corrector (AC), HRTEM direct imaging has become possible on one atom thick layer of carbon. This technique allows us to analyse the detailed atomic structures of and around the GBs together with fundamental information of each grain. In this work, graphene continuous films synthesized on platinum substrate with a specific configuration of CVD set-up are atomically characterized using AC-HRTEM imaging. Fig. 1 shows Raman spectra with corresponding HR-TEM images of our samples at different stages of the growth process. The evolution of both orientation and size of the grains is observed during the process. Nanometer size grains already connected with various orientations (see FFT on the inset) at stage I are further re-oriented and merged together along some pre-dominant directions in stage II and finally form large single crystal domains in the last stage. In stage II, two neighboring grains are typically aligned zig-zag to armchair at the boundary (Fig. 2a), which might be a low-energy crystallographic configuration. Slightly misoriented neighboring domains are connected with the presence of some dislocations (Fig. 2b). A lot of small domains are observed enclosed within larger ones (Fig. 2c) with the same zig-zag to armchair misorientation. Generally in CVD processes, the orientation and achievable size of grains are determined at the early nucleation stage. In our case, the small grains already form a continuous film at the early stage and further transform to low defective large crystals. We infer that such large scale re-crystallization of graphene is enhanced thanks to the platinum substrate. The latter hypothesis is supported density functional theory (DFT) calculation.
[1] O. Lehtinen et al., Nature Communications, DOI : 10.1038 /ncomms3098 (2013)


Fig. 1: Raman spectra of graphene continuous films at different stages of the synthesis process. HR-TEM images (a-c) correspond to typical atomic structures of stage I, II and III, respectively. Fourier transform shown on insets is taken from 50x50nm2 area of each sample. Scale bar is 2 nm.

Fig. 2: The upper panels show low-pass Fourier filtered images and the lower panels with maximum filtering [1]. Typical structures observed at stage II of (a) boundary between grains aligned zigzag to armchair, (b) dislocation between two grains with slight angle mismatch and (c) small domains inside another large grain. Scale bar is 1 nm.

Type of presentation: Oral

MS-2-O-2445 Atomic Resolution Structure Study of Fluorinated Graphene by Phase Restoration of Focal Series of Images

J Kashtiban R.1, Dyson M. A.1, Raveendran-Nair R.2, Zan R.4, Bangert U.3, Sloan J.1, Geim A. K.2
1University of Warwick, Coventry, UK, 2University of Manchester, Manchester, Uk, 3University of Limerick, 4Niğde University, Niğde, Turkey
r.jalilikashtiban@warwick.ac.uk

One approach towards band gap engineering of the graphene(Gr) which is scalable and cost effective is the chemical modification route to produce 2D derivatives of wonder material of suitable quality for monolayer devices although few such phases exist. GO and (CH)n are disordered or unstable while stoichiometric fluorographene(FG) exhibits significant corrugation. Fluorination of a Gr sheet results in FG, a Gr derivative with each fluorine atom connected to one carbon atom at the basal plane of Gr by a stronger SP3 bonding. Chair-C2F is a highly ordered material that demonstrates selective alternating fluorination and presents with an undistorted 2D morphology in contrast with stoichiometric but corrugated CF with the consequence that the former is a potentially much more tractable material for 2D device fabrication [1].A structural and morphological study of this material by means of atomic resolution TEM has yet not been reported. Exit wave restoration(EWR) by means of focal series of HRTEM images [2] can be performed to recover phase at the exit plane of the sample. Here we reveal, by atomic resolution EWR for the first time , that chair-C2F is a stable Gr derivative and demonstrates long-range order limited only by the size of a functionalized domain. Monolayer C2F was produced by partially fluorinating a suspended CVD grown Gr sample using direct fluorination method [3]. Focal series of images of Gr and chair-C2F were obtained at 80 kV in an aberration-corrected TEM instrument. EWR images reveal that imaged single carbon atoms and carbon-fluorine pairs in chair-C2F alternate strictly over domain sizes of at least 150 nm2 with electron diffraction indicating ordered domains up to ~0.16 μm2. Our results[4] also indicate that, within an ordered domain, functionalization occurs on one side only as theory predicts[1].

Figure 1 EWR and SIM EWRs of pristine Gr and monolayer chair C2F and figure 2 is a demonstration of long-range order within a 64 nm2 domain of C2F.

Refrences:

[1] Şahin, H, et al., Phys. Rev. B, 83 (2011)
[2] Coene, W.M.J., et al., Ultramicroscopy 64(1996) 109
[3] Nair, R.R., et al., Small, 6(2010) 2877
[4] Kashtiban et al., submitted.


We thank the EPSRC for funding through a studentship for M. A. D. and for a P. D. R. A. Fellowship for R. J. K. and additional support provided by the Warwick Centre for Analytical Science (EP/F034210/1).

Fig. 1: Figure 1 a-b, Graphene(Gr) and C2F models.c-d, Equivalent domains of EXP. and SIM. phase for Gr and chair-C2F respectively. e,f line profiles for the EXP and SIM phase for Gr and chair-C2F.g overlaid full plots of the EXP and SIM phase contrast for Gr, respectively.h overlaid full plots of the EXP and SIM phase contrast for chair-C2F, respectively.

Fig. 2: EWR image of a 250 nm2 sheet of highly ordered sheet C2F with an unrippled 64 nm2 domain highlighted. b,  64 nm2 domain in a exhibiting a high degree of order (scale bar = 2 nm). c, Surface plot from b in which the orange-yellow apexes correspond to ordered –CF< units. d, Line profiles (I-III) obtained through either >C–CF< or >C–CF< dumbbells.

Type of presentation: Oral

MS-2-O-2630 Unstacked double-layer templated graphene for high-rate lithium-sulfur batteries

Zhao M. Q.1, Zhang Q.1, Huang J. Q.1, Tian G. L.1, Nie J. Q.1, Peng H. J.1, Wei F.1
1Beijing Key Laboratory of Green Chemical Reaction Engineering and Technology, Department of Chemical Engineering, Tsinghua University, Beijing 100084, China
zhang-qiang@mails.tsinghua.edu.cn

Preventing the stacking of graphene is essential to exploiting its full potential in energy-storage applications. The introduction of spacers into graphene layers always results in a change in the intrinsic properties of graphene and/or induces complexity at the interfaces. Here, we show the synthesis of an intrinsically unstacked double-layer templated graphene via template-directed chemical vapor deposition. The as-obtained graphene is composed of two unstacked graphene layers separated by a large amount of mesosized protuberances and can be used for high-power lithium-sulfur batteries with excellent high-rate performance. Even after 1000 cycles, high reversible capacities of ca. 530 and 380 mA h g-1 are retained at 5 and 10 C, respectively. The preparation of the unstacked DTG is scalable and serves as a general strategy for the fabrication of a broad class of electrode materials for supercapacitors and lithium-ion, lithium-sulfur, and lithium-air batteries. We expect these DTG materials to have potential applications in the areas of environmental protection, nanocomposites, electronic devices, and healthcare because of their intrinsic large surface area, extraordinary thermal and electric conductivity, robust 3D scaffold, tunable surface chemistry, and biocompatible interface. Because unstacked layered nanostructures are not limited to graphene, we foresee a new branch of chemistry evolving in the stabilization of nanostructures through 3D topological porous systems.

References: Zhao MQ, Zhang Q, Huang JQ, Tian GL, Nie JQ, Peng HJ, Wei F. Nature Communications 2014, 5, 3410


We acknowledge the FEI company and Carl Zeiss Microscopy for their technical assistance. This work was supported by the National Basic Research Program of China (973 Program, 2011CB932602) and the Natural Scientific Foundation of China (No. 21306102).

Type of presentation: Oral

MS-2-O-2634 Oxidation resistance of reactive atoms in graphene

Chisholm M. F.1, Duscher G.2, Windl W.3
1Oak Ridge National Laboratory, Oak Ridge, TN, USA, 2University of Tennessee, Knoxville, TN, USA, 3The Ohio State University, Columbus, OH, USA
chisholmmf@ornl.gov

Carbon layers down to a thickness of a single layer have been known to form on metal surfaces for more than 50 years. Graphene on metal surfaces is also known to lead to catalytic deactivation.[1,2] However, it has not yet been recognized that graphene can also offer a level of oxidation protection to individual atoms or small clusters of atoms in/on the graphene layer. Here we will show using annular dark field imaging in a scanning transmission electron microscope and electron energy-loss spectroscopy (EELS) that Si and Fe atoms incorporated in a graphene layer are not oxidized. We further combine the microscopy data with first-principles density-functional calculations of the reaction of the impurity with graphene and the impurity with oxygen. Interestingly, our density-functional theory calculations explain these observations are due to preferential bonding of O to non-incorporated atoms and H-passivation effects.

Figure 1 shows annular dark-field images of a single Si atom segregated to a carbon vacancy and a single Fe atom segregated to a carbon divacancy. Si has also been found in graphene divacancies, but Fe has not been imaged in graphene single vacancy sites. The observed intensities are consistent with there being just a single impurity atom incorporated in the graphene defects, i.e. no additional oxygen. EELS confirms the identification of these impurity atoms as Si and Fe. The spectrum from a Si atom (Fig. 2c) shows an edge onset of about 102eV which is less than the oxide value and close to the edge onset for Si in of SiC. No oxidation features are apparent in the Si-L2,3 edge and no O-K edge was detectable in the spectrum. The core loss spectrum from a Fe atom (Fig. 2f) not only confirms the identification it also indicates that the Fe atom is not oxidized. The ratio of the Fe L3/L2 edges is lower than any seen for the various forms of Fe oxide. There have been no previously published experimental or theoretical studies on the oxidation resistance of iron or silicon incorporated in graphene to our knowledge. This is a potentially important discovery. Improved resistance to oxidation has important consequences for some catalytic reactions and small devices based on single atoms or small clusters of non-noble metals. Graphene as a substrate appears to protect single atoms and small clusters of atoms from oxidation without completely isolating them.[3]

References

1. S. Hagstrom, H.B. Lyon, G.A. Somorjai, Phys. Rev. Lett. 15 491 (1965).

2. R. Schlogl in Handbook of Heterogeneous Catalysis, vol. 1 G. Ertl, H. Knozinger, F. Schuth, J. Weitkamp (eds.) Wiley-VCH Weinheim 2008 p. 357.

3. M.F. Chisholm, G. Duscher, W. Windl, Nano Lett. 12 4651 (2012).


This research was supported by the Materials Sciences and Engineering Division of the Office of Basic Energy Sciences, U.S. Dept. of Energy and by NSF Award Number DMR-0925529 and the Center for Emergent Materials at The Ohio State University, a NSF MRSEC (Grant DMR-0820414). W.W. acknowledges support from the Ohio Supercomputer Center under project PAS0072.

Fig. 1: ADF images of graphene. The as recorded data (a,b) were corrected to remove noise and probe tail effects (c,d). Images a and c show a Si atom in a C vacancy site. Images b and d show an Fe atom in a C divacancy site. The scale mark on each image corresponds to 0.2 nm. Taken from Ref. 3.

Fig. 2: EEL spectrum image data from single atoms incorporated in graphene. ADF image of a Si atom in graphene (a), Si composition map (b) and spectrum obtained over the Si atom (c). ADF image of an Fe atom (d), Fe composition map (e) and the spectrum obtained over the Fe atom (f). The scale marks correspond to 0.2 nm. Taken from Ref. 3.

Type of presentation: Oral

MS-2-O-2651 Dislocations in bilayer graphene — Materials science meets physics

Butz B.1, Dolle C.1, Niekiel F.1, Weber K.2, Waldmann D.3, Weber H. B.3, Meyer B.2, Spiecker E.1
1Center for Nanoanalysis and Electron Microscopy (CENEM), Friedrich-Alexander University Erlangen-Nürnberg (FAU), Erlangen, Germany, 2Interdisciplinary Center for Molecular Materials (ICMM) and Computer Chemistry Center (CCC), Friedrich-Alexander University Erlangen-Nürnberg (FAU), Erlangen, Germany, 3Chair for Applied Physics, Friedrich-Alexander University Erlangen-Nürnberg (FAU), Erlangen, Germany
benjamin.butz@ww.uni-erlangen.de

Dislocations represent one of the most fascinating and fundamental concepts in materials science. First and foremost, they are the main carriers of plastic deformation in crystalline materials. Furthermore, they can strongly affect the local electronic and optical properties of semiconductors and ionic crystals. In materials with small dimensions they experience extensive image forces, which attract them to the surface in order to release strain energy. However, in layered crystals like graphite dislocation movement is mainly restricted to the basal plane. Thus the dislocations cannot escape enabling their confinement in crystals as thin as only two monolayers. To explore the nature of dislocations under such extreme boundary conditions, the material of choice is bilayer graphene, the thinnest imaginable quasi-2D crystal, in which such linear defects can be confined. Homogeneous and robust graphene membranes (Figure 1a, b) derived from high-quality epitaxial graphene on SiC [1] provide an ideal platform for their investigation.

Here we report on the direct observation of basal-plane partial dislocations (Burgers vector 1/3<1-100>) in freestanding bilayer graphene (Fig. 1a, b) by transmission electron microscopy and their detailed investigation by diffraction contrast analysis (Figure 1c, Burgers vector analysis 2c) and atomistic simulations (Figure 2a, b, and e) [2]. Our investigation reveals striking size effects. First, the absence of stacking fault energy, a unique property of bilayer graphene, leads to a characteristic dislocation pattern (Figure 1c, center), which corresponds to an alternating AB ↔ BA change of the stacking order (Figure 1c, right). Most importantly, our experiments in combination with atomistic simulations reveal a pronounced buckling of the bilayer graphene membrane (Figure 2a-d), which directly results from accommodation of strain (Figure 2e). In fact, the buckling completely changes the strain state of the bilayer graphene and is of key importance for the electronic properties. Our findings will significantly contribute to the future understanding of the structural, mechanical and electronic properties of bilayer and few-layer graphene.

[1] D. Waldmann, B. Butz, S. Bauer, J.M. Englert, J. Jobst, K. Ullmann, F. Fromm, M. Ammon, M. Enzelberger, A. Hirsch, S. Maier, P. Schmuki, T. Seyller, E. Spiecker, H.B. Weber, ACS Nano 7 (2013) 4441-4448

[2] B. Butz, C. Dolle, F. Niekiel, K. Weber, D. Waldmann, H.B. Weber, B. Meyer, E. Spiecker, Nature 505 (2014) 533-537


We thank J. Müller, E. Bitzek and A. Kohlmeyer for discussion and P. Schmuki for use of equipment. The research was supported by the DFG (SFB953, Cluster of Excellence EXC 315).

Fig. 1: a) Graphene membranes on SiC. b) One membrane at higher magnification: number of layers indicated. c) Series of bright-field (BF) and dark-field (DF) TEM images of same area. {11-20} images show pronounced contrast due to the presence of partial dislocations (dark lines), while {2-200} images depict respective changes of stacking sequence AB ↔ BA.

Fig. 2: a) Membrane topography and b) side-/top-view of pair of partial dislocations (change of stacking sequence enlarged shown). c) Burgers vector analysis using {11-20} DF images. d) Validation of Burgers vector analysis and atomistic model by DF-image simulation. e) Atomistic-strain distributions and derived disregistry/Burgers vector distributions.

Type of presentation: Oral

MS-2-O-2924 Thermal contact resistance between multiwalled carbon nanotubes and supporting substrates measured using electron thermal microscopy

Nilsson H.1, Voskanian N.1, Cumings J.1
1Department of Material Science and Engineering University of Maryland, College Park
hnilsson@umd.edu

Multiwalled carbon nanotubes (MWCNTs) have a high intrinsic thermal conductivity which makes them a promising material for heat management in nanoscale electronic devices. However, experimental results have shown that the total conductance is strongly limited by microscopic thermal resistances. These include the contact resistance between the nanotube and the substrate it rests upon, as well as the contact resistance with metal contacts. Due to characterization difficulties, exact values of the contact resistances have not been determined. In fact, literature estimates vary greatly. Some of these characterization difficulties include spatial resolution and the inability to separate resistance values within the system. An in-situ TEM technique developed by our group, called Electron Thermal Microscopy (EThM) allows us to obtain thermal maps with spatial resolution on the order of 10s of nanometers. The technique uses a specialized holder to locally heat an individual nanotube either directly by biasing or passively by a connected palladium heater wire. Indium nanoislands deposited on the backside of the sample membrane act as local temperature probes; their phase transition from solid to liquid at 156 degrees Celsius can easily be seen with dark field TEM imaging. Previous results using this technique have determined that the contact resistance with the silicon nitride is at least 250 Km/W (K. H. Baloch, N. Voskanian, M. Bronsgeest, and J. Cumings, "Remote Joule heating by a carbon nanotube," Nature Nanotechnology, 2012.). Here we present new device geometries featuring slits in the membranes to control the spread of heat through the membrane. This allows us to separate the heat transferred to the substrate via the MWCNT and via metal contacts. With more precise determination of the contact resistances it will be possible to get a better understanding of the thermal transport physics of MWCNTs.


Fig. 1: Schematic showing sample device geometry. A nanotube is heated passively via a palladium heater wire. Both sides of the tube are anchored with palladium pads. The slit through the membrane underneath one side of the nanotube is used to control heat transport across the membrane so that the contact resistance can be determined.

Fig. 2: Dark field TEM image demonstrating the melting of indium nanoislands when current is passed through a simple palladium heater wire.

Fig. 3: TEM image showing a palladium heater wire patterned on top of a nanotube.

Type of presentation: Oral

MS-2-O-3121 Imaging functional groups in graphene oxide at atomic resolution

Moreno M. S.1, Boothroyd C. B.2, Duchamp M.2, Kóvacs A.2, Monge N.3, Morales G. M.3, Barbero C. A.3, Dunin-Borkowski R. E.2
1Centro Atómico Bariloche, 8400 - S.C. de Bariloche, Argentina, 2ER-C and PGI-5, Forschungszentrum Jülich, D-52425 Jülich, Germany, 3Universidad Nacional de Río Cuarto, X5804BYA Río Cuarto, Argentina
meketo@gmail.com

Graphene oxide is a form of graphene with its surface modified by the addition of functional groups such as carboxylic groups, ketones and hydroxyl. The structure and distribution of the functional groups depends on the synthesis method used and they affect its chemical, electrical and mechanical properties. Of particular interest is the chemical composition and spatial distribution of these functional groups. Direct identification of the groups by high-resolution imaging is not possible at present, but they can be made visible by bonding heavy atoms, such as Ba, to selected groups and imaging the distribution of the heavy atoms.
High-resolution images of Ba doped graphene oxide taken at 80kV in a Cs and Cc corrected transmission electron microscope (TEM) show the structure of the graphene oxide with minimal electron beam damage but are not able to identify the Ba atoms. We attempted to use energy-filtered electron microscopy to image the Ba M5 and M4 edges at 781eV but radiation damage over the long exposures required prevented location of the Ba atoms to better than a few nm. Compositions measured using energy loss spectroscopy (EELS) revealed a progressive loss of O with increasing temperature and that 1% O is retained at 800°C.
High-resolution scanning transmission electron microscopy (STEM) at room temperature proved to be impossible due to contamination from migration of the functional groups on the graphene surface. Contamination was overcome by imaging with the graphene oxide above 400°C in a stable heating holder.
Simultaneous STEM energy-loss spectrum images (SI) and high-angle annular dark-field (HAADF) and bright-field (BF) images acquired at a specimen temperature of 800°C allowed us to correlate the location of the Ba atoms with features in the high-angle dark-field images at atomic resolution.
Simultaneously acquired HAADF and BF-STEM images (Fig. 1) show bright and dark dots at the same positions. These bright dots are identified as Ba atoms by EELS as shown in Fig. 2a. The corresponding elemental maps extracted from the SI are shown in Fig. 2b-d. The maps also show that Ca and O occur together and that Ba is not associated with a significant concentration of O. The positions of Ba atoms attached to functional groups on graphene oxide can therefore be mapped with atomic spatial resolution by using a combination of STEM and TEM techniques. The fact that Ca was observed to correlate strongly with O suggests that it could be used as a marker for the positions of the O-containing groups.


We acknowledge support from the European Union under the FP7 and a contract for an Integrated Infrastructure Initiative (ESTEEM2) and CONICET (Argentina) with a CONICET-DFG grant.

Fig. 1: Simultaneously acquired (a) HAADF and (b) bright-field STEM images at a specimen temperature of 800°C after recording the spectrum image (Fig.2), whose approximate area is marked by the box in (a). The bright dots in the HAADF image (a) correspond to the dark dots in the bright-field STEM image (b).

Fig. 2: (a) EEL spectrum extracted from a 3x3 pixel (0.25x0.25 nm) part of the spectrum image centred on a single Ba atom marked by the arrow in (b). (b-d) Elemental maps of the Ba M4,5, O K and Ca L2,3 edges obtained from the SI (200x200 spectra, dwell time 0.02s per spectrum, after VCA processing).

Type of presentation: Oral

MS-2-O-3205 Graphene Nanoribbons with atomically well-defined edges through Scanning Transmission Electron Microscopy

Vicarelli L.1, Xu Q.1, Zandbergen H. W.1
1Kavli Institute of Nanoscience, Delft University of Technology, Delft, The Netherlands
l.vicarelli@tudelft.nl

We recently demonstrated a controllable and reproducible method to obtain suspended monolayer graphene nanoribbons with atomically defined edge shape [1]. Our method exploits the electron-beam of a Scanning Transmission Electron Microscope (accelerated at 300 kV) to create vacancies in the lattice by knock-on damage and pattern graphene in any designed shape.

The small beam spot size (0.1 nm) enables close-to-atomic cutting precision, while heating graphene at 600o C during the patterning process avoids formation of beam-induced Carbon deposition and allows self-repair of the graphene lattice. Self-repair mechanism is essential to obtain well-defined (zig-zag or armchair) edge shape and, if the electron beam dose is lowered, to perform non-destructive imaging of the graphene nanoribbons.

Drawing the electron-beam path with a software script, we were able to obtain reproducible graphene nanoribbons with a minimum width of 2 nm and defined edges (see Fig. 1 and 2). In order to unravel some of the predicted properties of these graphene nanoribbons, we are currently exploring their transport properties through in-situ electrical measurement inside a Transmission Electron Microscope.

Early results show that large-area suspended graphene is stable over gaps of ~1 μm size. Both CVD grown and exfoliated graphene have been used. Performing 2 wire measurements, we saw that contact resistance between graphene and gold contact pads has a non-neglectable influence on the measurements, although it can be greatly reduced with in-situ thermal annealing above 300°C.

References:
[1] Q.Xu, M. Wu, G. F. Schneider, L. Houben, S.K. Malladi, C. Dekker, E. Yucelen, R.E. Dunin-Borkowski, and H.W. Zandbergen, ACS Nano 2013 7 (2), 1566-1572


The research leading to these results has received funding from the European Research Council, ERC Project n. 267922

We thank Kavli NanoLab Delft for the support provided in the fabrication of our NEMS devices. 

Fig. 1: Annular dark-field STEM image of a nanoribbon array, illustrating the reproducibility of the patterning. These four patterns were created using a script-controlled electron beam.

Fig. 2: HRTEM image of a nanoribbon in monolayer graphene sculpted at 300 kV and 600o C and imaged at 80 kV and 600o C. The yellow line indicates a zig-zag edge. An atom structure model for zig-zag edge is given as inset in the figure.

Type of presentation: Oral

MS-2-O-3438 3D insight on the catalytic nanostructuration of few-layer graphene

Melinte G.1, 2, Janowska I.2, Baaziz W.2, Florea I.1, Moldovan S.1, Arenal R.3, 4, Wisnet A.5, Scheu C.5, Begin-Colin S.1, Begin D.2, Pham-Huu C.2, Ersen O.1
1Institut de Physique et Chimie des Matériaux de Strasbourg, CNRS-Université de Strasbourg, 23, rue du Loess, 67037 Strasbourg, France , 2Institut de Chimie et Procédés pour l’Energie, l’Environnement et la Santé, CNRS, ECPM, Université de Strasbourg, 25, rue Becquerel, 67087 Strasbourg, France, 3Laboratorio de Microscopias Avanzadas, Instituto de Nanociencia de Aragon, Universidad de Zaragoza, 50018 Zaragoza, Spain, 4ARAID Fundation, Calle Mariano de Luna, 50018 Zaragoza, Spain, 5Department of Chemistry and Center for NanoScience, Ludwig Maximilians University, Butenandtstr. 11, 81377 Munich, Germany
georgian.melinte@ipcms.unistra.fr

The catalytic cutting of few-layer graphene (FLG) is attracting nowadays an increase attention due to its potential applications in both the field of catalysis and graphene nanoribbons (GNRs) fabrication.[1,2] The nanopatterning of FLG sheets with open and subsurface channels develops during a chemical reaction between the metal nanoparticles (NPs) and the carbon substrate under a hot gaseous atmosphere (H2, O2).[3,4] The use of the FLG powder in the field of catalysis is limited by the restacking process which significantly decreases its surface accessibility. By channeling the FLG sheets the restacking effect is significantly reduced. This is due to the creation of a porous network which will not only increase the surface accessibility but also will create a network of defects that will further serve as anchorage sites for the surface decoration. Moreover, the nanopatterning of FLG has the potential of creating nanosheets with well-defined shapes and edge configurations which can be transformed in single-layer GNRs by simple techniques as for instance the chemical exfoliation.
To characterize the channeling process and the obtained nanostructures we used an initial system consisting in Fe3O4 NPs dispersed on FLG sheets. The FLG/Fe3O4 composite has undergone a heating treatment in a H2 atmosphere. HR-TEM shows that the well-defined channels are not randomly oriented but follow specific crystallographic directions i.e. <11-20> and <10-10> (Figure 1), leading to the formation of two types of edge morphologies, zigzag and armchair, respectively. The electron tomography analyses reveal interesting features on both the nanopattering mechanisms and properties of the nanostructured FLG sheets. Figure 2 indicates the effect of a topographical step-up of the FLG sheet . Accordingly, one observes that after interaction the cutting direction remains unchanged but the depth of the open-surface channel is changing proportional with the height of the step-up. Figure 3 displays the impact of a step-up event with the height larger than the size of the NP. As previously, the cutting direction remains unmodified and the result is the formation of subsurface channel. When topographic step-down events are encountered by the active NPs, the particle either stops or changes the direction.
[1] Mei-xian Wang et al., J. Phys. Chem. Lett., 4, 1484−1488 (2013).
[2] Ci Lijie et al., Nano Res. 1, 116-122 (2008).
[3] Tim J. Booth et.al, Nano Lett. 11, 2689–2692 (2011).
[4] Datta S. et. al Nano Lett. 8(7), 1912-1915 (2008).


Fig. 1: HR-TEM images illustrating the preferential crystallographic orientation of the tranches with a zigzag (a) and armchair (b) edge morphology (scale bars 5 nm).

Fig. 2: a) TEM projection of a selected channel (scale bar 50 nm). b) XY slice through the reconstructed volume. c) XZ slices through the selected channel at the positions numbered in Figure 1a. d) YZ slice taken through a region indicated in Figure 1a with a yellow dotted line. Scale bars 20 nm.

Fig. 3: a) TEM projection of two subsurface channels (scale bar 50 nm). b) XY slice through the reconstructed volume (scale bar 20 nm). c) and d) YZ slices taken at the position indicated in Figure 1a with the yellow line (scale bar 20 nm).

Type of presentation: Oral

MS-2-O-3480 Plasmon Tailoring in Graphene through Lattice Impurities and Ad-Atoms

Pan C. T.1, Pierce W.2, Boothroyd C.3, Ramasse Q.4, Kepaptsoglou D.4, Bangert U.5
1School of Physics and Astronomy, The University of Manchester, Manchester M13 9PL, UK, 2School of Materials, The University of Manchester, Manchester M13 9PL, UK, 3Ernst Ruska-Centre for Microscopy and Spectroscopy with Electrons and Peter Gruenberg Institute, Juelich Research Centre, D-52425 Juelich, Germany, 4SuperSTEM Laboratory, STFC Daresbury Campus, Daresbury WA4 4AD, UK, 5Department of Physics and Energy, University of Limerick, Limerick, Ireland
Ursel.Bangert@ul.ie

Although there is considerable documentation on efforts to tailor and employ plasmons to merge photonics and electronics, and use surface plasmons for subwavelength optics [1,2] and enhancement of the photovoltaic conversion efficiency, especially by making use of the surface plasmons at metal nano-clusters [3], little has been reported on single atom plasmon effects [4]. Graphene’s potential for terahertz nano-scale plasmonic devices, has so far only been realised via gating and patterning [5,6]. However, defects in the graphitic plane, including vacancies and dopant atoms, can intrinsically alter the electronic structure and hence lead to effects such as plasmon enhancement and change of the plasmon energy in the uv/vis region, a phenomenon that can be exploited for coupling with light.

We present observations, using high resolution (S)TEM in combination with electron energy loss spectroscopy and energy filtered imaging, of the effects of single or few-atom impurities on plasmons in the uv/pi-plasmon energy regime in graphene. We accompany the experiments by WIEN calculations, which reveal new transitions in graphene for various metal ad-atoms species (Ti, Pd) and also for Si (fig. 1) and substitutional dopants such as B and N: a peak at around 1-2 eV is introduced which is not present in energy loss spectra of pristine graphene. Both, position and intensity of this peak change according to doping/dosing levels. The increase of the latter shifts this peak towards the uv regime (3eV). These transitions are mostly ascribed to single particle (SPE) and intraband excitations or to SPE-π plasmon coupling and not to the creation of new plasmon peaks in the graphene-dopant system. The same applies to defect and edge-states. Our experimental observations are in general support of the above predicted additional absorption features in the uv. More so, we observe intensity enhancement around metal atoms (e.g., Pd) at graphene edges (fig. 2), which we also find, although to a much lesser extent, at pristine graphene edges. This intensity increase does, however, not arise from new spectral features and is ascribed to the enhancement of intrinsic low loss features of graphene, where metal atoms/ defects act as atomic antennae, due to donation of d-electrons, in the case of transition metals. The efficiency of this process appears to vary with the transition metal, and seems to be high for, e.g., Pd.

[1] W L Barnes et al, Nature , 424 (2003) 824-830

[2] S A. Maier, H A Atwater, J. Appl. Phys 98 (2005) 011101

[3] J Nelayah, et al Nature Physics, 3 (2007) 348-352

[4] W Zhou et al, Nature Nanotech, 7 (2012), 161-165

[5] T J Echtermeyer et al, Nature Comms (2011), DOI: 10.1038/ncomms1464

[6] L Ju et al Nature Nanotech, DOI: 10.1038/NNANO (2011) 146


Fig. 1: Simulated in-plane (left) and out-of-plane (right) EEL spectra of a single (a) Au, (b) Pd, (c) Cr, (d) Ti and (e) Si adatom on (solid curves) and spectra after the carbon atoms of the graphene are removed (red dashed curves). Spectra are shifted along the Y axis and are all on the same scale. The two bottom spectra are of pristine graphene.

Fig. 2: Images from an EFTEM image series obtained in a monochromated triple -corrected Titan-PICO: a) enhancement at 3.5-4 eV and b) depletion at 5-5.5 eV of the loss intensity at a hole in graphene with Pd deposit, c) HREM image prior to the EFTEM series with magnified boxed area, showing Pd atoms (arrowed) decorating the edge of the hole.

Type of presentation: Poster

MS-2-P-1483 Structural Analysis of Electron-Beam-Irradiated C60 Single Crystal Films Using High-Resolution Transmission Electron Microscopy and Electron Diffraction

Masuda H.1, Yasuda H.2, Onoe J.1
1Tokyo Institute of Technology, 2Osaka University
hidmasuda@nr.titech.ac.jp

  We have found that one-dimensional (1D) uneven peanut-shaped C60 polymer is formed by electron-beam (EB) irradiation of a pristine C60 film [1], and exhibits novel physical properties arising from 1D metal, such as the geometric curvature effects on the Tomonaga-Luttinger liquid states [2]. For the polymer structure, in situ infrared (IR) spectra and density-functional calculations suggested the 1D polymer has a cross-linked structure (Fig. 1(c)) roughly close to that of the P08 C120 isomer (Fig. 1(b)) obtained from the general Stone-Wales transformation [3]. Although the previous results indicated the polymer to have 1D peanut-shaped structure, we have examined the structure of the 1D polymer formed from a C60 single crystal (SC) film more precisely, using HRTEM and electron diffraction (ED).
  The 1D C60 polymer film was formed on a mica substrate by EB irradiation of a pristine C60 SC film in an UHV chamber. After confirmed that all C60 molecules were polymerized to form the 1D polymer using in situ IR spectroscopy, we ripped the film off the mica and mounted it on a Cu mesh in air, and observed the film by TEM.
  Figure 2 shows HRTEM images and ED patterns of the pristine and the EB-irradiated C60 films. The C60 film is a FCC SC with [111] orientation, which contains twins as stacking faults on (111), and shows weak spots E1 and E2 as 1/3 and 2/3 of 422 series, respectively. The EB-irradiated C60 SC film shows three new features. Firstly, E1 and E2 become intense, indicating symmetry reduction and FCC transferred to HCP. Secondary, spots of 220 series become doublet. Since the corresponding distance of these spots is 5.0 Å and 4.6 Å, respectively, the intermolecular distance (di) between adjacent C60 molecules is estimated to be 10.0 Å and 9.3 Å for each. Finally, each spot becomes an arc-like stretched line of ca. 9.2°. This arises from a slight loss of the orientation. These results show the asymmetric shrinkage of crystal structure along one given direction.

  C60 FCC structure changes to 1D polymer BCO (Fig. 1(d)). Furthermore, judging from the intense E1 spot, BCO changes to HCP-m (Fig. 1(e)). Figs. 2(e, f) show the simulated ED patterns of BCO and HCP-m based 1D C60 polymer model (the di of 9.3 Å) with 3-fold symmetry derived from three possible polymerization directions on (111) of FCC C60 SC film, using QSTEM code [4]. Since each pattern well reproduces the experimental ED pattern, a mixed stacking model of BCO and HCP-m 1D C60 polymer structures is suitable for the EB-irradiated C60 SC film [5].
[1] J. Onoe et al., Appl. Phys. Lett. 82, 595 (2003).
[2] J. Onoe et al., Europhys. Lett. 98, 27001 (2012).
[3] A. Takashima et al., J. Phys. D: Appl. Phys. 45, 485302 (2012).
[4] C. Koch, Ph.D. thesis (2002).
[5] H. Masuda et al., to be submitted.


This work was supported by “Advanced Characterization Nanotechnology Platform (MEXT)” at Osaka University and by the collaborative research fund of J-Power.

Fig. 1: Molecular models of C60 (a), P08 C120 isomer (b), 1D C60 polymer (c). The unit cell of 1D C60 polymer with body-centered orthorhombic structure (BCO) (d), and hexagonal closed-packed-based monoclinic structure (HCP-m).

Fig. 2: Experimental HRTEM images (insets: FFT patterns) and ED patterns of pristine C60 SC (a, c), EB-irradiated C60 SC (b, d), simulated ED patterns using 1D C60 polymer model with intermolecular distances of 9.3 Å in polymer direction (e, f).

Type of presentation: Poster

MS-2-P-1536 Mapping electronic states in Graphene

Löffler S.1,2, Pardini L.3, Hambach R.4, Kaiser U.4, Schattschneider P.1,2, Draxl C.3
1Institute of Solid State Physics, Vienna University of Technology, Austria, 2University Service Centre for Transmission Electron Microscopy, Vienna University of Technology, Austria, 3Department of Physics, Humboldt University Berlin, Germany, 4Electron Microscopy Group of Materials Science, Ulm University, Germany
stefan.loeffler@tuwien.ac.at

Graphene and similar carbon-based materials are currently the focus of much research. In particular, their peculiar electronic properties are arousing a lot of interest. At the same time, the question of the influence of defects – such as vacancies or dopant atoms – is of particular practical importance. Recently, it has been reported that different dopant atom configurations change the charge distribution [1] and, therefore, give rise to different electron energy-loss spectrometry (EELS) signals [2]. Likewise, introducing vacancies changes the local crystal structure [3] and, hence, also the local charge distribution and EELS signal.

This gives rise to the hope to map the electronic states using energy-filtered transmission electron microscopy (EFTEM) [4]. In this work, we present predictions regarding the possibility of direct mapping of electronic states in both ideal, pristine Graphene and Graphene with defects using EFTEM. To that end, calculations of the electronic states of Graphene, with and without defects, were carried out using the full-potential all-electron density functional theory (DFT) package "exciting" [5]. Subsequently, its output was used to calculate the inelastic scattering kernels which, combined with elastic scattering calculations, ultimately result in EFTEM images.

The EFTEM images were calculated for a variety of acceleration voltages and lens aberration functions to simulate realistic conditions, as well as investigate the optimal experimental conditions. Fig. 1 shows EFTEM simulations for Graphene with and without core-hole effects included in the DFT calculations. It is clearly visible that this greatly alters the expected EFTEM images. Fig. 2 goes a step further and shows the images to be expected when mapping a vacancy. This demonstrates that EFTEM at high spatial resolution could become an invaluable tool for the study of electronic states in carbon-based materials.

[1] Meyer et al., Nat Mater 10 (2011) 209
[2] Zhou et al., PRL 109 (2012) 206803
[3] Meyer et al., Nano Lett 8 (2008) 3582
[4] Löffler et al., Ultramicroscopy 131 (2013) 39
[5] http://exciting-code.org/


The authors acknowledge financial support by the FWF (I543-N20), the DPG and the MWK Baden-Württemberg.

Fig. 1: Simulated EFTEM images for 40 keV electrons for pristine multi-layer Graphene without (left) and with (middle) core-hole effects under ideal imaging conditions. In both cases, an energy loss of 6.5 eV above the C K-edge onset was used. Additionally, the partial density of states is shown (right).

Fig. 2: Simulated EFTEM images for Graphene with a vacancy. The images show π* (left) and σ* (right) states. The circle marks the vacancy. Both images were simulated for a beam energy of 80 keV and ideal imaging conditions.

Type of presentation: Poster

MS-2-P-1730 In-Situ Environmental TEM Observation of the Formation of Defects in Growing Carbon Nanotubes

Yoshida H.1, Takeda S.1
1The Institute of Scientific and Industrial Research, Osaka University
h-yoshida@sanken.osaka-u.ac.jp

As-grown carbon nanotubes (CNTs) generally have various grown-in defects, such as vacancies, pentagon-heptagon pairs, bending, and irregular interlayer spacing. It is well known that the electronic and mechanical properties of CNTs are affected by these grown-in defects. An understanding of the formation mechanism of the CNT grown-in defects is required for the growth of defect-free CNTs exhibiting ideal properties and CNTs with intentionally induced defects exhibiting modified properties. Recent environmental transmission electron microscope (ETEM) [1] observations of chemical vapor deposition (CVD) growth of CNTs have provided us with knowledge of the growth mechanism. We have clarified that CNTs grow from structurally fluctuating iron carbide Fe3C and iron molybdenum carbide (Fe,Mo)23C6 nanoparticles [2-4]. However, in situ studies on the formation of defects in growing CNTs are limited. In this study, we have elucidated the origin of grown-in defects in CNTs, such as bending, irregular interlayer spacing, change in the diameter, and change in the number of graphitic layers, by in situ atomic-scale ETEM observations of the CVD growth of CNTs [5].

Figure 1 shows the growth of a CNT with a drastic disorder of the interlayer spacing. We also observe large changes in the CNT diameter during growth as shown in Fig. 2. Our ETEM observations clearly demonstrate that deformation of the nanoparticle catalysts during CNT growth triggers the formation of these grown-in defects [5]. The small deformation of nanoparticle catalysts at the interface with CNTs gives rise to the formation of bends and disorder of the interlayer spacing (Fig. 1) in CNTs. The changes in the diameter (Fig. 2) and number of graphitic layers in CNTs are caused by the large protrusion on and shrink deformations of nanoparticle catalysts. Based on the ETEM observations, we will propose the formation mechanism of grown-in defects in CNTs.

References

[1] S. Takeda and H. Yoshida, Microscopy, 62 (2013) 193.

[2] H. Yoshida, S. Takeda, T. Uchiyama, H. Kohno, and Y. Homma, Nano Lett., 8 (2008) 2082.

[3] H. Yoshida T. Shimizu, T. Uchiyama, H. Kohno, Y. Homma, and S. Takeda, Nano Lett., 9 (2009) 3810.

[4] H. Yoshida, H. Kohno, and S. Takeda, Micron, 43 (2012) 1176.

[5] H. Yoshida and S. Takeda, Carbon, 70 (2014) 266.


This work was supported by JSPS KAKENHI Grant Number 24710117.

Fig. 1: Drastic disorder in the interlayer spacing of graphitic layers in a growing CNT from a nanoparticle catalyst [5]. The observation time is shown in the images.

Fig. 2: Large Change in the diameter of a growing CNT from a nanoparticle catalyst [5]. The lattice images in the nanoparticle catalyst corresponds to a (Fe,Mo)23C6-type structure [3,4]. The observation time is shown in the images.

Type of presentation: Poster

MS-2-P-1745 Electron microscopic evidence for a tribologically induced phase transformation as the origin of wear in diamond

Zhang X.1, Schneider R.1, Müller E.1, Mee M.2, Meier S.2, Gumbsch P.1, 2, Gerthsen D.1
1Karlsruhe Institute of Technology (KIT), Karlsruhe, Germany, 2Fraunhofer Institute for Mechanics of Materials IWM, Freiburg, Germany
xinyi.zhang2@kit.edu

The origin of wear and the low friction coefficient of diamond is still an intensely debated problem in tribology. Here we study coarse-grained diamond films, deposited by plasma-enhanced chemical vapor deposition, which were tribologically loaded on a ring-on-ring tribometer against a similar diamond counterpart. The microstructure of worn and unworn regions of the diamond film was studied by transmission and scanning electron microscopy. Amorphous carbon (a-C) layers are observed on both as-deposited and on tribologically tested diamond, but differ significantly as far as thickness and morphology are concerned. The a-C layer with a thickness of up to several 100 nm on as-deposited diamond is attributed to the plasma deposition process. For the tribologically tested region of the film, the TEM images (Fig. 1) demonstrate that the µm-sized grains at the rough original diamond surface are almost completely flattened indicating that a significant amount of material must have been removed including the residual a-C layer from the deposition process. In contrast to the as-deposited a-C residue, the tribo-induced a-C layer is comparably uniform with a thickness below 100 nm. The TEM sample from the wear track prepared by conventional techniques (Fig. 2) confirms the findings of the FIB-prepared sample. A few of the TEM samples containing a tribo-induced a-C layer contain grain boundaries of the underlying polycrystalline diamond in the electron transparent region. It is found that the thickness of the a-C layer changes quite abruptly on grains with different crystallographic orientations (white arrow in Fig. 2). Fig. 3 clearly shows that the interface between the crystalline diamond and the tribo-induced amorphous a-C layer is not crystallographically flat but displays a nm-scale roughness. The anisotropic phase transformation and the small roughness of the interface are regarded as evidence for an atom-by-atom wear process. Quantitative electron energy loss spectroscopy of the C-K ionization edge, performed in a transmission electron microscope, reveals the transition from sp3-hybridized C-atoms in diamond to a high fraction (65 %) of sp2-hybridized C-atoms in the tribo-induced a-C layer within a region of less than 5 nm thickness.


XZ acknowledges funding from China Scholarship Council (CSC) (No. 2010606030). PG acknowledges support from Deutsche Forschungsgemeinschaft DFG (project grant Gu 367/30).

Fig. 1: Overview TEM image of a cross-section FIB-lamella taken from the wear-track region.

Fig. 2: Overview TEM image of a conventional cross-section TEM sample prepared from the wear-track region.

Fig. 3: HRTEM image of the interface region between crystalline diamond and the tribo-induced amorphous carbon layer with the diamond oriented along the [111] zone axis. The approximate position of the interface is marked by a dashed line.

Type of presentation: Poster

MS-2-P-1752 Atom-by-atom STEM EELS investigation of n- and p- doped graphene.

Kepaptsoglou D. M.1, Seabourne C. R.2, Hardcastle T.2, Nicholls R.3, Pierce W.4, Zan R.4, Bangert U.4, Scott A. J.2, Ramasse Q. M.1
1SuperSTEM Laboratory, STFC Daresbury Campus, United Kingdom, 2Institute for Materials Research, SPEME, University of Leeds, Leeds, United Kingdom, 3Department of Materials, University of Oxford, Oxford, United Kingdom, 4School of Materials, University of Manchester, Manchester, M13 9PL, United Kingdom
dmkepap@superstem.org

Graphene, or the miracle material as it has become known, has promised to revolutionize the world of electronics by replacing Si-based technology [1,2]. Graphene however is an excellent conductor – often described as a ‘zero bandgap’ semiconductor, an attribute which so far limits its widespread application in devices. Among various solutions to tailor its properties for practical implementation, the introduction of dopants in the graphene lattice is predicted to have a drastic effect on graphene's band structure [3], such as the opening of an optical bandgap or an increase in charge carrier density resulting in n- or p-type doping, with carrier concentrations allowing practical transistor applications. The introduction of dopants such as N in graphene is most commonly achieved during the chemical growth process, with varying levels of success regarding the purity of the samples, which often contain contaminants, defects and secondary impurities. We have recently demonstrated an alternative, cleaner method by successfully doping freestanding single layer graphene with N and B through low energy ion implantation [4], achieving retention levels of the order of ~1%.

In this work we use STEM-based spectroscopy [5,6], to study the impact of single N or B dopant atoms on the electronic structure of the graphene membrane. Z-contrast imaging and atomically resolved electron energy loss spectroscopy were performed in a Nion UltraSTEM100 dedicated STEM instrument and were used to unambiguously identify single dopant atoms (fig. 1) and to determine the doping levels as a function of ion implantation energy and flux. Furthermore, the electronic structure modifications due to the presence of these dopant B or N atoms are strikingly demonstrated by a clear signature in the near-edge fine structure of the B and N EELS K edges but also that of C K edge of neighboring C atoms (fig. 1). Ab initio calculations are used to simulate experimental spectra (fig. 2) and to rationalize the experimental observations, thus providing further insight into the nature of bonding around the foreign species.

[1]A. K. Geim et al., Nat Mater 6, 183 (2007).
[2]K. Kim et al., Nature 479, 338 (2011).
[3]S. Casolo et al., Nanostructured Mater 115, 1 (2010).
[4]U. Bangert et al., Nano Lett 13, 4902 (2013).
[5]Q. M. Ramasse et al., Nano Lett 13, 4989 (2013).
[6]R. J. Nicholls et al., ACS Nano 7, 7145 (2013).


SuperSTEM is the UK National Facility for Aberration-Corrected STEM and is funded by the UK Engineering and Physical Sciences Research Council (EPSRC)

Fig. 1: HAADF STEM images of a) N and b) B implanted graphene showing single N and B substitutions in the graphene lattice, respectively. The images were low-pass filtered and colorised for clarity.

Fig. 2: EEL spectra from a) N and b) B implanted graphene samples, showing the near edge fine structure of the C K, N K edges and B K and C K edges of the N- and B-doped sample, respectively.

Type of presentation: Poster

MS-2-P-1807 Water-soluble multi-layered graphene nanosheets via high temperature acidic treatment of graphite oxide

Gevko P. N.1, Tur V. A.1, Okotrun A. V.1, Bulusheva L. G.1
1Nikolaev Institute of Inorganic Chemistry SB RAS, Novosibirsk, Russia
paul@niic.nsc.ru

In recent years graphene has attracted wide attention from the scientific community. Its exceptional conductivity, high specific surface area and mechanical strength can be used in the energy storage devices (such as supercapacitors, lithium-ion batteries), composites, sensors, as well as makes it promising as a substrate for the deposition of various particles (metal oxides and sulfides , metal nanoparticles, conductive polymers, etc.)
For some applications, it is necessary to obtain stable solutions (or dispersions) of graphene. From the viewpoint of practical application, aqueous solutions of graphene have great advantages compared with the organic solvents and the possibility of obtaining of these solutions is actively investigated recently.
Previously we have proposed the method of high-temperature treatment of graphite oxide in concentrated sulfuric acid. It was shown that such treatment leads to the formation of the product which contains a large number of holes in carbon layers and called as a «perforated graphite». It was noted that during the process the color of the liquid phase is changed to brown. Subsequent study of the composition of this phase showed that it is, in fact, a solution of multi-layered graphene nanosheets (MGNS). In this work, we perform the further development of method of water-soluble graphene nanoparticles obtaining in order to increase the yield of MGNS, as well as the comprehensive investigation of these particles. It was shown that treatment of graphite oxide in a concentrated sulfuric acid at a temperature of 200°C leads to the formation of MGNS with in-plane size of ~300 nm and small amount of oxygen in its structure in a form of various functional groups. To improve the yield of MGNS the mixture of concentrated sulfuric/nitric acids was used. It was find that initial graphite oxide almost totally decomposed with the formation of MGNS, but the resultant product has a larger amount of oxygen in its structure.
Thus, the present study demonstrates a simple high temperature acidic treatment of graphite oxide as a method of obtaining of easy water-soluble MGNS as well as data of comprehensive study of these nanosheets.


Fig. 1: AFM-images of MGNS on silicon substrate

Type of presentation: Poster

MS-2-P-1822 Relations between deficiencies in CVD deposited graphene and the lattice defects of the Ni (111) substrate

Fogarassy Z.1, Rümmeli M. H.2, 3, Gorantla S.4, Bachmatiuk A.2, 3, 4, Dobrik G.1, Kamarás K.5, Biró L. P.1, Havancsák K.6, Lábár J. L.1
1Hungarian Academy of Sciences, Research Centre for Natural Sciences, Institute for Technical Physics and Materials Science, Hungary, 2IBS Center for Integrated Nanostructure Physics, Institute for Basic Science (IBS), Republic of Korea, 3Department of Energy Science, Department of Physics, Republic of Korea, 4Leibniz Institute for Solid State Materials Research Dresden, Germany, 5Hungarian Academy of Sciences, Wigner Research Centre for Physics, Institute for Solid State Physics and Optics, Hungary, 6Department of Materials Physics, Eötvös Loránd University, Hungary
fogarassy@mfa.kfki.hu

In this work graphene layers were investigated that were deposited by chemical vapor deposition (CVD) on nickel (111) thin film substrates. The Ni (111) thin substrates themselves were grown previously on bulk sapphire (0001) substrates at 550°C with 33nm/min. growth velocity. The nickel layer grew epitaxially on the sapphire, as it is shown by the diffraction in figure 1. The Ni (111) planes were parallel to the sapphire (0001) planes. Twin boundaries were formed in the nickel layers. At some places grains were rotated with 30° about an axis perpendicular to the sapphire surface.
The graphene deposited on the Ni (111) was investigated by TEM, Cs-corrected HRTEM, STM, AFM and by SEM / EBSD. According to [1] above 600°C, attempts to grow graphene on Ni (111) only resulted in the appearance of multi-layer turbostratic graphite. Here we show how the orientation of the first graphene layer and the appearance of additional turbostratic layers are related to the quality and orientation of the Ni (111) substrate. The nickel substrates were annealed in hydrogen before the graphene deposition. The graphene was deposited at atmospheric pressure from a mixture of argon, methane and hydrogen gases. The sample was cooled in argon and hydrogen gas mixture. Large area continuous single-layer graphene formed on top of nickel (111). Above it both another turbostratic graphene partial layer (Fig. 2), and graphite flakes were formed at certain locations. Such flakes are seen in the SEM image of figure 3.a. The continuous single-layer graphene formed epitaxially on the nickel (111) and the epitaxial relation was corroborated by both STM and TEM analyzes.
The growth of the thinner and smaller flakes was suppressed by changing the gas concentration. SEM and EBSD studies showed that the thicker flakes grew above the incoherent twin boundaries and high-energy (30°) grain boundaries of nickel.
Our results show that the first atomic layer grew as a continuous and epitaxial graphene layer on nickel (111). The additional local turbostratic layers and graphite flakes grew above the incoherent twin boundaries or high-energy boundaries in the nickel substrate. No thicker flakes were observed above coherent twin boundaries.
[1] Patera L.L. et al., ACS Nano 7 (9) 7901–7912 (2013)


Fig. 1: SAED diffraction pattern recorded from the nickel and sapphire interface. The nickel layer grew epitaxially on the sapphire.

Fig. 2: Cs-corrected HRTEM image recorded from a local turbostratic graphene flake over the large area continuous graphene layer.

Fig. 3: Figure a) SEM image. The dark patches of different gray shades are from an additional turbostratic graphene layer and graphite flakes. Figure b) EBSD orientation map. The straight lines mark locations where twin boundaries reach the surface of nickel.

Type of presentation: Poster

MS-2-P-1848 Observation of Nanobubbles on Graphene with Atomic Force Microscopy

Ko H. C.1, Yang C. W.1, Hsu W. H.1, Hwang I. S.1
1Institute of Physics, Academia Sinica, Nankang, Taipei, 105, Taiwan, R.O.C.
hsienchenko@gmail.com

For the last 14 years, a variety of experimental studies have revealed the existence of nanobubbles at liquid/solid surface. Nanobubbles have attracted much scientific interest because of several highly disputed properties and many potential applications in fields from surface to nanofluidics [1]. Previous studies suggested that nanobubbles prefer to form on hydrophobic surfaces [2]. In this work, we prepare a substrate of a flat hydrophilic substrate with a small area covered by a hydrophobic material. The purpose is to see the preference of nanobubble formation on such an inhomogeneous substrate in water. Muscovite mica is a hydrophilic substrate that is strongly attracted to water, but graphene interacts weakly with water. Here we prepare a mica substrate with a small patch of mechanical exfoliated graphene layers. We inject water supersaturated with air on this sample and image the solid/water interface with atomic force microscopy (AFM). Figure 1a shows a height image taken with PeakForce mode. A high density of nanobubbles forms on graphene, but none is seen on the mica. Figure 1b is a higher-resolution image, which shows that nanobubbles can form on graphene of different thicknesses. Our observations demonstrate that the surface hydrophobicity has significant effect on nanobubble formation. Further study may help understanding the accumulation of gases at solid-water interfaces.

References

[1] Attard, P. Langmuir 12, 1693-1695 (1996)

[2] Agrawal, A. Nano Lett. 5, 1751-1756 (2005)


We thank support for this work from National Science Council of ROC (NSC96-2628-M-001-010-MY3 and NSC99-2112-M-001-029-MY3) and Academia Sinica

Fig. 1:  (a) Height image of graphene deposited on a mica surface in air-super-saturated water. (b) A higher-resolution image taken in the outlined region in (a).

Type of presentation: Poster

MS-2-P-1947 Pt-terminating Carbyne Observed by Aberration-Corrected TEM

Kano E.1, 2, Hashimoto A.2, Takeguchi M.2
1University of Tsukuba, Tsukuba, Japan, 2National Institute for Materials Science, Tsukuba, Japan
KANOU.Emi@nims.go.jp

  Carbyne is a one-dimensional (1D) single atomic linear chain, composed of sp-hybridized carbon atoms. Some low-dimensional carbon allotropes, such as zero-dimensional fullerenes, quasi-one-dimensional (quasi-1D) carbon nanotubes, and two-dimensional graphene, have recently been discovered and have attracted worldwide attention for their unique properties. Theoretical studies have predicted that carbynes have more remarkable properties than the other low-dimensional materials. However, these properties have not yet been proven experimentally because of the difficulties encountered in production and during observation.
  Here, we report a novel, reproducible method of carbyne formation using Pt atoms on graphene. The formation and dynamics of carbynes were observed on an atomic scale by aberration-corrected TEM (JEM-ARM200F, JEOL, Japan) operating at an accelerating voltage of 80 kV. The samples were obtained by transferring monolayer graphene membranes onto in situ heating chips (E-chips for Aduro, Protochips, USA). Pt was deposited by a plasma sputtering system. We observed independent Pt atoms that appeared on a clean graphene surface. Using independent Pt atoms is an important key to produce and stabilize carbyne chains. We observed the migrations of Pt and C atoms on the graphene surface at 400 °C with an in situ heating holder (Aduro heating holder, Protochips, USA).
  Figures 1A–C show TEM images of carbyne formation. Three Pt atoms captured some carbon atoms, resulting in the formation of a C-shape chain (Fig. 1A). The Pt atoms and the chain moved around freely for 82 s, and then the chain suddenly turned into a straight chain (Fig. 1C). Both ends of the chain were terminated by Pt atoms, and the chain remained motionless for more than 20 s. From Fig. 1C, the length between the Pt atoms at both ends was measured to be approximately 1.5 nm. It corresponds to a model composed of 11 carbons with two Pt atoms (Fig. 1D). Fig. 1E shows a simulated image using this model.


A part of this work was supported by “Nanotechnology Platform Project” of the Ministry of Education, Culture, Sports, Science and Technology (MEXT), Japan.

Fig. 1: Figure 1. (A to C) TEM images of carbyne chain, which exhibited a variety of ringed, curved, straight shapes. Scale bars are 1 nm. (D) Atomic model of straight carbyne. (E) Simulated TEM image.

Type of presentation: Poster

MS-2-P-1987 Active role of carbon during the formation of porous cerium oxide layer used as catalyst in fuel cell

Lavkova J.1,2, Khalakhan I.1, Chundak M.1, Vorokhta M.1, Potin V.2, Matolinova I.1, Matolin V.1
1Department of Surface and Plasma Science, Charles University, Czech Republic, 2Laboratoire Interdisciplinaire Carnot de Bourgogne, Université de Bourgogne, France
vpotin@u-bourgogne.fr

The global trend in renewable energy investment is developing the new energetic system using hydrogen; well know as a hydrogen economy. The key devices seem to be the fuel cells (FC) that convert chemical energy from hydrogen or hydrocarbons into kinetic or electrical energy. The most critical component of the FC is catalyst. The versatile element in catalysis is platinum (Pt) that efficiently mediating a multitude of chemical reaction. Unfortunately, Pt is rare element and its high price, exceeded that of gold, limits large-scale applications. Therefore, not surprisingly, the goal of reducing the amount of Pt is the major driving force in the catalysis research.

The Pt - CeO2 porous layers have been reported to be significantly active catalysts for CO oxidation, hydrogen production and oxidation of ethanol. Thin – film technology permits to produce large variety of hetero-materials with different composition (low concentration of platinum) and morphology (porous structure) of layers. The knowledge of the materials structure is fundamental for the best understanding of their physical and chemical properties.

The key role of carbon in the porosity creation of the catalyst layer is presented. The morphology of the CeO2 films prepared by magnetron sputtering on graphite foil was investigated by using microscopy tools – the Atomic Force Microscopy (AFM), the Scanning Electron Microscopy (SEM) and the Transmission Electron Microscopy (TEM). These studies show modification of carbon, confirmed by the Energy-Dispersed X-ray Spectroscopy (EDX) – see Fig. 1. The formation of cerium carbides crystals on the catalyst-substrate interface was observed using the High-Resolution TEM (HR-TEM). Moreover, the reduction of cerium as a result of the interaction with the carbon support was obtained by spectroscopies – the X-ray Photoelectron Spectroscopy (XPS) and the Electron-Energy Loss Spectroscopy (EELS). Finally, the structural model of the system is designed.


The authors acknowledge the support by the Czech Science Foundation under grant No. 13-10396S, ANR within IMAGINOXE project (ANR-11-JS10-001) and by EU FP-7-NMP-2012 project “chipCAT” under contract No. 310191. J.L. is grateful to the Conseil Regional de Bourgogne (PARI ONOV 2012).

Fig. 1: The modification of carbon after deposition of 20nm thick CeO2 layer on graphite foil – a) the material contrast obtained by the Scanning Transmission Electron Microscopy (STEM), b) the element map obtained by EDX.

Type of presentation: Poster

MS-2-P-2025 Dynamic Motion of Ru-Polyoxometalate Ions (POMs) on Functionalized Few-Layer Graphene

Ke X.1, Turner S.1, Quintana M.2, Hadad C.3, Montellano-Lopez A.3, Carraro M.4, Sartorel A.4, Bonchio M.4, Praro M.3, Bittencourt C.5, Van Tendeloo G.1
1EMAT, University of Antwerp, Groenenborgerlaan 171, B-2020 Antwerpen, Belgium, 2Instituto de Fisica, Universidad Autonoma de San Luis PotoSi, Manuel Nava 6, Zona Universitaria, 78290 San Luis Potosi, SLP, Mexico, 3CENMAT, INSTM, UdR di Triest, Dipartimento di Scienze Chimich e Farmaceutich, University of Trieste, Piazzale Europa 1, I-34127 Trieste, Italy, 4ITM-CNR, Department of Chemical Sciences, Unviersity of Padova, Via F. Marzolo 1, 35131 Padova, Italy, 5ChiPS,, Unviersity of Mons, Rue du Par 20, B-7000 Mons, Belgium
xiaoxing.ke@uantwerpen.be

Recent advances in state-of-art aberration-corrected transmission electron microscopy (AC-TEM) have demonstrated its power in resolving the atomic structure of nanomaterials and nanohybrids down to the limit. Particularly, the interfacial structures of nanohybrids have strong influence on the properties and performances of the materials and thus need to be understood at atomic level. Aberration-corrected imaging shows enhanced resolution and improved signal-to-noise ratio, which largely benefits the straightforward interpretation at nanohybrids interface, particularly for soft 2D nanomaterials. An example on the interface study of a water oxidation catalyst functionalized on graphene surface is demonstrated at ultra-high resolution in this abstract.

The as-studied water oxidation catalyst, a tetraruthenate oxo-cluster ([Ru4(H2O)4(μ-O)4(μ-OH)2(γ-SiW10O36)2]10-, referred to as polyoxometalate (POM) ions in this text, has been recently discovered to hold promising application in water splitting. The effective grafting of the Ru4POM catalysts on a conductive substrate is therefore crucial in order to promote its further application in nanodevices on the large scale. Few layer graphene (FLG) is a best candidate due to its superior electric and mechanic properties such as high carrier mobility etc. Exploring the behaviour of Ru4POM on graphene supports is therefore one of the key issues in further applications, where the selection of functional group is important in tailoring its performance. Thus in this study we demonstrate how the interactions between Ru4POM molecules and supporting graphene layers can be studied by aberration corrected transmission electron microscopy (AC-TEM) at low voltage (80kV). Under the 80 kV irradiation of the electron beam the Ru4POM demonstrates dynamic motion on the graphene surface. The motion of the Ru4POM is captured as a series of images and is shown to vary as a function of time under certain constraints in the Ru4POM rotation. The frequency and amplitude of rotation is found to be related to the nature of the functional group used, including a polyamine dendron and a N,N,N-trimethyl benzenaminium moiety in our study. Distortions in the Ru4POM structure are revealed as well, suggesting that the ions can stand instantaneous structural changes without losing their integrity. The stability of the Ru4POM-graphene hybrid during the imaging corroborates the long-term robustness of the material applied to multielectronic catalytic processes.


X. Ke and G. Van Tendeloo acknowledge funding from the European Research Council under the FP7 ERC Grant Nº 246791–COUNTATOMS.

Type of presentation: Poster

MS-2-P-2040 Spatial distribution of particles in the graphene-nanoparticles system

Mantlikova A.1, Pacakova B.1, Kalbac M.2, Repko A.3, Vejpravova J.1
1Institute of Physics of the ASCR, v.v.i., Na Slovance 2, Prague 8, 182 21, Czech Republic, 2J. Heyrovský Institute of Physical Chemistry of the ASCR, v.v.i., Doleškova 3, Prague 8, 182 23 Czech Republic, 3Faculty of Science, Charles University in Prague, Albertov 6, Prague 2, 128 43, Czech Republic
mantlikova@fzu.cz

Graphene (GN) has been in the focus of intensive research in material science and nanotechnology in recent years due to its unique electrical, optical, thermal and mechanical properties. Creation of the graphene-nanoparticles (GN-NPs) system could lead to the improvement of the physical properties of GN due to the change of its topography via creation of wrinkles.

We have focused on the basic characterization of the GN-NPs systems, especially on the study of the NP spatial distribution on the substrate, which has a significant influence on the GN wrinkling. The samples possessing different concentration of the CoFe2O4 NPs (6 nm) on the Si/SiO2 substrate covered by the GN layer were characterized by High Resolution Scanning Electron Microscopy (HRSEM) and Atomic Force Microscopy (AFM). Different concentration of the NPs for individual samples was confirmed both by the HRSEM and AFM measurements and creation of the GN wrinkles around the NPs below the GN monolayer and their dependence on the NP concentration has been observed (Fig. 1).

The real NP spatial distribution determined by the HRSEM and AFM was compared with that one obtained from simulation in Matlab as follows: the NPs were randomly distributed inside the square box of 1 µm edge length, divided to the regular lattice with inter-node distances equal to the NP diameter for prevention of the NP overlap. The mean interparticle distances were calculated in both cases (real and simulated NP spatial distribution) for the nearest neighbors using the triangulation procedure. The results of both NP spatial distributions clearly demonstrate decrease of the interparticle distances (Fig. 2). Moreover, the resulting interparticle distances obtained from the simulation correspond very well to those obtained from the real positions of NPs determined by the AFM, showing that the NPs are randomly distributed on the surface and the influence of the substrate corrugations on the NP distribution is negligible.


This work was supported by Czech grant agency, project number P204/10/1677.

Fig. 1: The SEM image of the GN-NPs sample with high (a) and low (b) concentration of NPs. The GN wrinkles are clearly visible on both images, the border between the GN and substrate could be found on right image.

Fig. 2: The concentration, c dependence of interparticle distance, d for GN-NPs samples (a). The simulated NPs spatial distribution for the least (1:10000) concentrated (b) and the most (1:1000) concentrated (c) samples.

Type of presentation: Poster

MS-2-P-2084 Projected potential of graphene estimated by off-axis electron holography

Geiger D.1, Biskupek J.1, Algara-Siller G.1, Kaiser U.1
1Electron Microscopy Group of Materials Science, University of Ulm, Albert-Einstein-Allee 11, 89081 Ulm, Germany
dorin.geiger@uni-ulm.de

Graphene [1] is one of the most investigated 2D-structures in the last decade. Some efforts were undertaken [2,3,4] to determine the projected potential (PP) of graphene for different layer thicknesses. Off-axis electron holography (EH) [5] allows to recover numerically the image wave and, for known residual aberrations of the TEM, to reconstruct the object exit wave. The accuracy is significantly increased by using a Cs-corrected TEM [6,7].
CVD graphene transferred to holy carbon was cleaned using active carbon or Al2O3 powder at ~200°C at the sample surface [8] (fig.1a). Because of the quite long time needed for TEM-adjustment in EH-mode, searching for adequate sample locations and positioning in EH, new contamination appears (fig. 1b,c). Consequently, the graphene layers are often not perfectly clean and the PP is generally higher than the ideal values, cleanliness of the graphene layers proves out to be essential for accurate results.
The illumination of our Cs-corrected FEI Titan 80-300 TEM with rotatable Möllenstedt biprism, was optimized for EH at 80kV [2], where beam damage of graphene is reduced. Taking the C1–C3 condenser lens setting [9], the elliptical illumination could be optimized and using a reduced extraction voltage by ~2kV, a decrease of the electron energy distribution was achieved. Finally the hologram contrast could be noteworthy improved [2]. The experiments till now show, presumably due to contamination, a quite large dispersion and a tendency to higher values than the calculated ones. Using the independent atom model (IAM) and/or the density functional theory (DFT), image simulations and calculations of the PP, were made using the programs QSTEM [C. Koch] and JEMS [P. Stadelmann].
To characterize the local thickness and the PP of graphene, we took holograms in high-resolution TEM. The analysis of the profile lines in the reconstructed object phase allows the determination of the local thickness variations (fig. 2). Contamination and the EH-restriction, to use only object areas next to vacuum, make difficult to find large ideal uniform sample areas. Phase jumps, related to thickness variations, show for most of the results up to now, phase shifts of <0.08 rad at a phase detection limit per single hologram of ~2π/70. To conclude our studies, additional results with better statistics will follow.

1. A. K. Geim, K. S. Novoselov, Nature Materials 6 (2007) p.183.
2. D. Geiger et al., MS.7.P199,  3. F. Börrnert et al., IM.5.P110,  4. L. Ortolani et al., MS.7.P211, Proc. MC 2013.   5. H. Lichte, UM 20 (1986), p. 293.    6. M. Haider et al., UM 75 (1998) p. 53. 

7. D. Geiger et al., Micr. Microanal. 14 (2008), p. 68.  8. G. Algara-Siller et al., submitted. 

9. F. Genz et al., IM.2.P046, Proc. MC 2013.


This work was supported by the DFG (German Research Foundation) and the Ministry of Science, Research and the Arts (MWK) of Baden-Württemberg in the frame of the Sub-Angstrom Low-Voltage Electron microscopy (SALVE) project.

Fig. 1: Specimen contamination state immediately after insertion in the TEM (a), about three hours later (b) and at the end of the TEM-session (c).

Fig. 2: Reconstructed phase image from the object exit wave of graphene layers at vacuum, showing some contamination (a) with the phase profile along the arrow (b). Holograms were taken with Cs-corrected FEI Titan 80-300 TEM at 80kV in HRTEM mode.

Type of presentation: Poster

MS-2-P-2099 Microstructure investigation of Ru/CNT hybrid synthesized by different methods

Zhang B.1, Pan X.1, Su D.1
1Shenyang National Laboratory for Materials Science, Institute of Metal Research, Chinese Academy of Sciences, Shenyang, China
bszhang@imr.ac.cn

Ruthenium (Ru) as catalysts or functional additive has drawn considerable interest, owing to its fascinating unique properties and amazingly high efficiency in various reactions.[1-2] The shape, size distribution and spatial dispersion of supported Ru nanoparticles (NPs) could significantly contribute to the final performance of catalysts, especially for many structure-sensitive heterogeneous reactions. For the purpose of simultaneously achieving excellent catalytic activity and high stability for practical applications, there have been many attempts to explore preparation routes, such as impregnation, the polyol process, metal organic chemical vapor deposition, deposition–precipitation, microwave irradiation, and so on. However, there are few reports about comparing the structures of supported Ru nanocomposite synthesized by these methods systematically. Herein, we selected Ru supported on carbon nanotube (Ru/CNT) as model hybrid material, and used advanced electron microscopy to reveal the detailed structural features of Ru/CNTs prepared by difference methods, such as the local and surface structures, particles size distribution (PSD), and thermal stability. For instance, Ru NPs supported on oxygen functionalized CNTs (Ru/OCNTs) synthesized by ethylene glycol (EG) reduction and conventional impregnation (IP) methods have been compared.[3] Transmission electron microscopy (TEM), scanning TEM (STEM), and X-ray diffraction (XRD) characterizations reveal that the sample prepared by EG reduction method with ultra-small size and highly dispersed Ru NPs onto OCNTs, is superior to that obtained by conventional impregnation method (Fig. 1). The result of Ru NPs evolution process investigated by in situ heating TEM (Fig. 2) indicates that the Ru/OCNTs prepared by EG reduction method has good thermal stability, which may lengthen catalyst service life availably. In addition, we also studied the structural rearrangements of Ru NPs supported on CNTs from twinned Ru NPs into Ru single nanocrystals under the microwave irradiation by aberration-corrected electron microscope.[1] Our work can be expected as an important reference for the design and fabrication of ultra-small metal NPs with optimal morphologies and high thermal stability for a variety of chemical reactions.

Reference
[1] B. Zhang, X. Ni, W. Zhang, L. Shao, Q. Zhang, F. Girgsdies, C. Liang, R. Schlögl, D.S. Su, Chem. Commun. 2011, 47, 10716-10718.
[2] X. Ni, B. Zhang, C. Li, M. Pang, D.S. Su, C.T. Williams, C. Liang, Catal. Commun. 2012, 24, 65-69.
[3] X. Pan, B. Zhang, B. Zhong, J. Wang, D.S. Su, Chem. Commun. 2014, in press (DOI: 10.1039/C3CC48710E).


We gratefully acknowledge the financial support provided by NSFC of China (21203215, 21133010, 51221264, 21261160487), MOST (2011CBA00504), Strategic Priority Research Program of the Chinese Academy of Sciences (No. XDA09030103) and the China Postdoctoral Science Foundation (2012M520652).

Fig. 1: HAADF-STEM images of Ru/OCNTs-IP (a) and Ru/OCNTs-EG (b). The insets in (a) and (b) are the corresponding histograms of PSD.[3]

Fig. 2: Typical TEM images of the Ru/OCNTs-EG sample treated using an in situ heating TEM holder at RT (a), 350 oC (b), and 600 oC (c) and the HAADF-STEM image after heating (d). Inset in Fig. 2d is the corresponding histograms of PSD.[3]

Type of presentation: Poster

MS-2-P-2102 Formation mechanism and bending properties of carbon nanotetrahedron/nanoribbon structures

Kohno H.1
1Kochi University of Technology
hideokohno@gmail.com

A carbon nanoribbon is formed when a carbon nanotube flattens in one direction. We have found that a switching in the flattening direction results in the formation of a carbon nanotetrahedron in the middle of a carbon nanoribbon [1] (Fig. 1). Our TEM and SEM observations suggest a model of its formation mechanism as follows. When a carbon nanotube is expelled from an Fe catalyst nanoparticle, the tube is forced to flatten, and there are two preferable directions of flattening, which we call the origami mechanism. When one direction is dominant, a nanoribbon is formed, while a nanotetrahedron is formed when a switching of the flattening direction occurs (see Ref. 1 for more details).

To reveal bending properties of our carbon nanotetrahedron/nanoribbon structures, they were examined using a micromanipulator in a TEM and their bending behavior was observed in-situ [2]. We have found that a nanotetrahedron/nanoribbon structure bent at a nanotetrahedron/nanoribbon junction, and that the bending was reversible and repeatable. The nanotetrahedron/nanoribbon structures kept their shape during being bent and did not expand to take a tubular form. Our results show that the nanotetrahedron/nanoribbon structures have excellent durability against bending. The nanotetrahedron/nanoribbon structures can be bent at nanotetrahedron/nanoribbon junctions sharply and do not break, therefore we expect that the nanotetrahedron/nanoribbon structures can be used for three-dimensional wiring.

[1] Hideo Kohno, Takuya Komine, Takayuki Hasegawa, Hirohiko Niioka, and Satoshi Ichikawa, Nanoscale 5 (2013) 570.
[2] Hideo Kohno and Yusuke Masuda, to be submitted.


This work was supported in part by Adaptable and Seamless Technology Transfer Program through Target-driven R&D, Japan Science and Technology Agency.

Fig. 1: TEM images and schematic illustration of carbon nanotetrahedron/nanoribbon structures (from Ref. 1).

Type of presentation: Poster

MS-2-P-2174 A new structural model for Graphene Oxide and Reduced Graphene Oxide as revealed by core EELS and DFT

Tararan A.1, Zobelli A.1, Benito A.2, Maser W.2, Stéphan O.1
1Laboratoire de Physique des Solides, Univ. Paris-Sud, CNRS UMR 8502, F-91405, Orsay, France, 2Department of Chemical Processes and Nanotechnology, Instituto de Carboquímica ICB-CSIC, C/Miguel Luesma Castán 4, E-50018 Zaragoza, Spain
anna.tararan@u-psud.fr

Graphite oxide (GO) is known since the middle of the XIX century. In the latest years it has attracted a renewed interest as a precursor for a cheap large-scale production of graphene. Indeed, GO conserves graphite layered structure with an expanded interlayer distance that facilitates exfoliation. A subsequent reduction yields a material whose properties are very similar to those of graphene but strongly depend on the local structure and stoichiometry. However, many questions remain still open about GO and reduced GO (RGO) chemical homogeneity and the functional groups effectively present.
In previous spectroscopy studies the oxygen content in GO ranges from 22% to 32%. However, TEM images revealed that GO is very inhomogeneous at the nanometer scale. Still, no spatially resolved spectroscopic studies have yet been reported and only average evaluations are provided in literature.1 EELS in a STEM could give access to the suitable scale but GO and RGO are highly sensitive to irradiation.
In this study we overcame this limitation by adopting an experimental set up combining a liquid nitrogen cooling system at the sample stage, a low accelerated electrons beam (60 keV) and a liquid nitrogen cooled CCD camera with a low read-out noise of three counts r.m.s. and a negligible dark count noise. Hyperspectral core EELS images have been acquired in a low dose mode (order of 105 e-nm-2) at a 10 nm spatial resolution.2
Chemical maps for GO and RGO (see figure) show regions within individual flakes with different oxidation levels. Whereas oxygen rates averaged over the whole area are in agreement with literature, we observe that the oxygen content can locally rise up to 60%.
Lower oxidized GO regions present a fine structure at the carbon K-edge similar to amorphous carbon, while highly oxidized regions show specific core EELS signatures. RGO samples display the well-known fine structure profile of graphite, proving an excellent restoration of the carbon network. Nevertheless regions characterized by residual oxygen exhibit an additional sharp peak.
These results have been combined with complementary DFT analysis of formation and binding energies for different oxygen functional groups and concentrations and EELS spectra simulations. This allowed us to provide a new structural model compatible with our experimental findings. We suggest a full functionalization with hydroxyl groups in the strongly oxidized regions, while in lower oxidized regions also epoxide groups are expected.

1K. A. Mkhoyan, A. W. Contryman, J. Silcox, D. A. Stewart, G. Eda, C. Mattevi, S. Miller, and M. Chhowalla, Nano Lett. 9, 1058 (2009).

2M. M. v. Schooneveld, A. Gloter, O. Stéphan, L. F. Zagonel, R. Koole, A. Meijerink, W. J. M. Mulder, and F. M. F. d. Groot, Nat Nano 5, 538 (2010).


The authors acknowledge support from the European Union in Seventh Framework Programme under Grant Agreement No. 312483 (ESTEEM2).

Fig. 1: EELS hyperspectral analysis of Graphene Oxide and Reduced Graphene Oxide: oxygen concentration maps, associated histograms and carbon K-shell EELS edges extracted from the selected regions.

Fig. 2:
Type of presentation: Poster

MS-2-P-2219 Surface formation of electrospun carbon nanofiber mats controlled by HRSEM

Zhigalina V. G.1, Ponomarev I. I.2, Razorenov D. Y.2, Ponomarev I. I.2
1Shubnikov Institute of Crystallography RAS, Moscow, Russia, 2Nesmeyanov Institute of Organoelement Compounds RAS, Moscow, Russia
v.zhigalina@gmail.com

Hydrogen fuel cells will play the leading role among alternative energy sources in the XXI century. Electrocatalytic and gas diffusion layers of a fuel cell consist of electroconductive materials and metal layers, which are the most critical components. Carbon nanofiber nonwoven materials are the most promising materials owing to their thermal and chemical resistance, higher sorption capacity, electrical conductivity and mechanical properties. Their properties determine the most important fields of their application: for purifying various liquid and gaseous media, as a reinforcing filler in composites for accumulating gaseous or condensed substances and also for creating catalysts with higher activity, selectivity and stability. Polyacrylonitrile is the most promising and the most used polymer for producing carbon nanofiber nonwoven materials. Recently electrospinning has been used to create new materials for various alternative power supplies [1]. By this method highly porous fiber mats can be molded from solutions of polymers. Due to various additives the properties and characteristics of the produced materials can differ widely.
The main problems are to improve the porosity of the carbon nanofiber mats, to reduce their electroresistivity and to decrease the precious catalytic metal concentration. The aim of the present work is to investigate the influence of different treatments on the morphology, metal particle distribution and surface structure of electrospun polyacrylonitrile mats.
The surface investigation of obtained mats was performed by a high resolution scanning electron microscopy (HRSEM) in a FEI Quanta 250 FEG and a FEI Helios 600 DualBeamTM with EDX analysis.
The obtained electrospun polyacrylonitrile fiber mats were 10-100 μm thick, which depends on the molding conditions and treatment temperature with the fiber diameters in the range of 50-400 nm. Most of the fibers had a characteristic diameter of 100-150 nm and a length of several tens of microns [2]. These mats with a smooth surface are shown in Fig. 1. High temperature annealing (at 1200 and additional 2800 oC) and chemical treatment (by polyvinylpyrrolidone and polyimide) led to significant changes in the morphology, length and surface condition (Fig. 2). The chemical treatment was performed for a better deposition of Pt particles because of the formation of cavities on the fibers’ surface [3]. As a result, a thick Pt nanoparticles coating was formed on their surface (Fig. 3).

1. Dong Z, Kennedy SJ, Wu Y Journal of Power Sources 2011 196 4886
2. Ponomarev II et al. Doklady Physical Chemistry 2013 448(6) 670
3. Zhigalina VG et al. Nanomaterials and Nanostructures - XXI Century 2012 4 36


The investigation was supported by RFBR grant № ofi-m-11-03-12115.

Fig. 1: Nontreated electrospun carbon fiber mats.

Fig. 2: Carbon fibers with a damaged surface after annealing at 1200 and 2800 oC.

Fig. 3: HRSEM image of carbon fibers coated by Pt (a) and corresponding EDX spectrum (b).

Type of presentation: Poster

MS-2-P-2376 Low-Energy Electron Diffractive Imaging of Graphene Based on a Single-Atom Electron Source

Hsu W. H.1, 2, Chang W. T.1, Lin C. Y.1, 3, Chen Y. S.1, Lai W. C.1, 3, Hwang I. S.1, 2
1Institute of Physics, Academia Sinica, Nankang, Taipei, Taiwan, 2Department of Materials Science and Engineering, National Tsing Hua University, Hsinchu, Taiwan, 3Department of Physics, National Taiwan University, Taipei, Taiwan
hsuwh@phys.sinica.edu.tw

    A single-atom tip (SAT) can be an ideal field emitter of electron beams. It has been shown that noble-metal covered W(111) SATs can be reliably prepared1,2 with several controllable profiles3. The growth of the faceted pyramidal tips is a thermodynamic process. These SATs are both physically and chemically stable and can be regenerated through a simple annealing in vacuum, ensuring a long operation lifetime. Both the brightness and spatial coherency of these single-atom electron sources are orders of magnitude better than those of the state-of-the-art electron sources used in current electron microscopes4.

    We have built a low-energy electron point projection microscope (PPM) combined with a retractable micro-channel plate detector (MCP), housed in an ultra-high vacuum (UHV) chamber, to image nano-objects. A schematic is displayed in Fig. 1. The PPM is a shadow microscope where an object is placed between the electron point source and a detector screen. The magnification of the image depends on the tip-sample distance (d) and the sample-detector distance (D). We record the high resolution bright-field image when D is large. On the other hand, the diffraction patterns of the object at large angles can be obtained when D is small.

    With the advantages of the high stability of single-atom electron source and high contrast due to low energy of the source, some characteristics on graphene can easily be observed. Fig. 2b shows ribbon-like patterns in each diffraction disk, which are also visible in the bright-field image in Fig. 2a. Figs. 3a, b, and c show the bright-field images of graphene with the same beam energy of 270 eV at 0 second, 75 seconds, and 150 seconds, respectively. Both growth (labeled with yellow arrows) and migration (labeled with red arrows) of particles can be observed in a continuous-time imaging.

Reference

1 H.-S. Kuo, I.-S. Hwang, T.-Y. Fu, J.-Y. Wu, C.-C. Chang, and T.T. Tsong , Nano Lett. 4(12), 2379 (2004).

2 H.-S. Kuo, I.-S. Hwang, T.-Y. Fu, Y.-C. Lin, C.-C. Chang, and T. T. Tsong, Jap. J. Appl. Phys. 45, 8972 (2006).

3 W.-T. Chang, I.-S. Hwang, M.-T. Chang, C.-Y. Lin, W.-H. Hsu, and J.-L. Hou, Rev. Sci. Instrum. 83, 083704 (2012).

4 C.-C. Chang, H.-S. Kuo, I.-S. Hwang, and T. T. Tsong, Nanotechnology 20, 115401(2009).


We acknowledge the financial support from Academia Sinica and National Science Council.

Fig. 1: Schematic of our electron point projection microscope with a retractable MCP. The single-atom tip is driven by a X-Y-Z piezo-manipulator for approaching and scanning of the tip to the sample. The magnification of the bright field image is M = (D + d)/d.

Fig. 2: (a) Bright-field image of suspended graphene sheet, recorded with 500 eV and sample-detector distance D = 130 mm. (b) Diffraction pattern of the same sample recorded with 480 eV and sample-detector distance D = 35mm.

Fig. 3: Bright-field images of graphene recorded with 270 eV at (a) t = 0 s, (b) t =75 s, and (c) t = 150 s.

Type of presentation: Poster

MS-2-P-2386 Implantation and atomic scale characterization of self-interstitials in free standing graphene

Lehtinen O.1, Vats N.1, Algara-Siller G.1, Knyrim P.1, Kaiser U.1
1Ulm University, Materials Science Electron Microscopy
ute.kaiser@uni-ulm.de

A surplus density of carbon atoms is introduced into free-standing graphene by means of low-energy implantation [1]. The implantation is conducted using an evaporating carbon coating apparatus, designed for depositing thin layers of amorphous carbon on non-conducting specimens for electron microscopy. By careful tuning of the deposition parameters, a low enough density of extra atoms is reached in order to produce isolated self-interstitial dimers in graphene. The structure of these defects is imaged at the atomic scale, employing aberration corrected high resolution transmission electron microscopy. The earlier predicted, completely sp2-hybridized structural configurations of ad-dimers in graphene [2,3] are experimentally verified. Additionally larger aggregates of extra atoms and edge dislocation dipoles incorporated in the graphene lattice are observed, and based on atomistic modeling, such structures are determined to be energetically favourable arrangements for the extra atoms. All of the adatom structures are predicted to strongly buckle out-of-plane. Such blister-like structures can be expected to have higher reactivity than pristine graphene, which can be advantageous when functionalization of graphene is desired. Further on, defect structures containing surpulus carbon atoms have been predicted to have exciting electronic and magnetic properties [4], and our experiment demonstrates that such structures can, in fact, be fabricated.

[1] Lehtinen, O., Vats, N., Algara-Siller, G., Knyrim, P. and Kaiser, U., (2014), in review

[2] Lusk, M. T. and Carr, L. D., Phys. Rev. Lett. 100 (2008) 175503.

[3] Lusk, M. T., Wu, D. T. and Carr, L. D., Phys. Rev. B 81 (2010) 155444.

[4] Lehtinen, P. O., Foster, A. S., Ayuela, A., Krasheninnikov, A. V., Nordlund, K. abd Nieminen, R. M., Phys. Rev. Lett. 91 (2003) 017202.

[5] Lehtinen, O., Kurasch, S., Krasheninnikov, A. V. and Kaiser, U., Nat. Commun. 4 (2013) 2098


This work was supported by the DFG (German Research Foundation) and the Ministry of Science, Research and the Arts (MWK) of Baden-Württemberg in the frame of the Sub-Angstrom Low-Voltage Electron microscopy (SALVE) project and by the DFG through the TR21 project.

Fig. 1: Self-interstitial dimers in graphene at 80 kV AC-HRTEM. First column raw images of inverse Stone-Thrower-Wales defect and its two polymorphs. Second column: same images after maximum filtering [5], Third column: wire frame models of the defects. Fourth column: relaxed atomic structures. Scale-bar 1 nm.

Type of presentation: Poster

MS-2-P-2423 Topologically induced dielectric response in multilayer graphene nanocones

Hage F. S.1, 2, 3, Kepaptsoglou D. M.1, Ramasse Q. M.1, Seabourne C. R.4, Scott A. J.4, Prytz Ø.3, Gunnæs A. E.3, Helgesen G.2, 3, Brydson R.4
1SuperSTEM Laboratory, SciTech Daresbury, Daresbury, U.K , 2Physics Department, Institute for Energy Technology, Kjeller, Norway, 3Department of Physics, University of Oslo, Oslo, Norway., 4Institute for Materials Research, SPEME, University of Leeds, Leeds, U.K
fshage@superstem.org

Among the multitude of known carbon nanostructures, graphene nanocones are quite unique. These multilayer cones are characterised by macroscopic apex angles (0˚, 112.9˚, 84.6˚, 60˚, 38.9˚ and 19.2˚), which correspond to the incorporation of zero to five pentagons (P=0-5) in a graphene sheet [1] (where P=0 corresponds to the case of a flat disc). Due to their topology, graphene cones are ideal for investigating the effect of pentagonal defects on local valence electron structure in multilayer carbon nanostructures. This was done here by recording the dielectric response at cone apices by means of valence electron energy loss spectroscopy in the aberration corrected dedicated scanning transmission electron microscope.

Fig. 1a shows a distinct feature in the loss spectrum at 1.5 eV, originating from the tip of a two pentagon cone (P=2, fig. 1b). From ab inito simulations this feature was attributed to the presence of the pentagonal defects themselves. This was explained by pentagons inducing topology specific localised low energy states, where the 1.5 eV feature arises as a sum over interband transitions involving such states. Upon extension, this indicates that multilayer graphene cones should show great promise as field emitters [2].

Localisation of collective modes was investigated by the momentum dependence (i.e. dispersion) of their associated loss peaks: where a vanishingly small dispersion corresponds to a localised state. Fig. 2a shows the π plasmon dispersion from the tip of multilayer cone with five pentagons at its apex (P=5, fig. 2b) compared to that of a flat multilayer discs (P=0). While a parabolic dispersion indicates significant ‘graphite-like’ plasmon propagation in the disc, the vanishing plasmon dispersion in the P=5 cone indicates a high degree of confinement at its apex [3]. The observed slightly negative cone plasmon dispersions will also be discussed. All data were acquired with a Nion UltraSTEM100 operated at 60kV, and all computational modelling was carried out using the CASTEP DFT code at the University of Leeds ARC1 facility.

[1] A. Krishnan et al., Nature, 388 (1997) 451
[2] F.S. Hage et al., Nanoscale, 6 (2014) 1833
[3] F.S. Hage et al., Physical Review B, 88 (2013) 155408


SuperSTEM is the national facility for aberration corrected STEM, supported by the UK Engineering and Physical Sciences Research Council. This work was supported by the Research Council of Norway under Contract No. 191621/V30.

Fig. 1: Figure 1 (a) Valence electron energy loss spectrum originating from the tip of the (P=2) cone shown in the high angle annular dark field (HAADF) image in (b).

Fig. 2: Figure 2 (a) Dispersions of the π plasmon of a cone with five pentagons at the apex (P=5) and a flat disc (P=0). (b) HAADF image of a five pentagon cone (P=5).

Type of presentation: Poster

MS-2-P-2542 Atomic resolution of nanocrystalline Ge and SbTe encapsulated inside carbon nanotubes

Marks S.1, Kashtiban R.1, Sloan J.1
1University of Warwick
s.r.marks@warwick.ac.uk

We have the ability to image individual atoms using high resolution transmission electron microscopy (HRTEM) yet for most materials this has only been achieved whilst looking at a section of a larger system. There has been relatively little work investigating nanocrystalline systems that have been formed inside extremely small growth spaces[1]. This is of interest due to the possibility that with this bottom up growth method we may see the formation of new structures[2].
Melt filling is a technique that allows us to encapsulate materials inside ≈1 nm carbon nanotubes (CNTs) therefore allowing us to image molecules formed inside confined spaces. This is achieved by grinding a combination of CNTs and the desired material. It is then heated to above its melting temperature, cooled and annealed, repeatedly, inside a furnace before being dispersed onto a lacy carbon grid. The dispersed filled carbon nanotubes were then imaged using a JEOL ARM200F TEM/STEM at 80kV accelerating voltage, a HRTEM is required as individual atomic resolution is essential. The CNTs were commercially acquired and had a range of diameters from 1.1 – 1.6 nm. Two systems were investigated, Germanium (Ge) and Antimony Telluride (SbTe). Once successfully filled the systems were then recreated using Crystalmaker and simulated using SimulaTEM in an attempt to reproduce the achieved image.
Ge was successfully filled into CNTs but rendered a small filling ratio due to the Ge reacting with the lab glass used during the melting phase. Even though the filling ratio was low, filled CNTs were captured with high resolution. The Ge was found to be growing in the [010] crystal orientation relative to the bulk with lattice spacings that are consistent with bulk Ge. The system captured in Figure 1 has been recreated and simulated to prove it has formed from bulk Ge. Streaking is visible on the atomic spots of Ge, this is due to the CNT being tilted rather than exactly perpendicular to the beam.
SbTe also successfully filled into the CNTs but again with a low filling ratio. Two different systems were observed within the SbTe sample; the first a 4 atom motif and the second a 5 atom (Figure 2). This is of high interest as a single atom variation will affect the quantum structure of the system. Both systems were recreated and found to have lattice spacing similar to that of the bulk with growth once again in the [010] direction, this is due to SbTe and Ge both having a FCC crystal structure.

[1] J Sloan et al. "Two layer 4: 4 co-ordinated KI crystals grown within single walled carbon nanotubes." Chemical Physics Letters 329 (2000): 61-65.
[2]R Carter et al. "Correlation of structural and electronic properties in a new low-dimensional form of mercury telluride." Physical review letters 96 (2006): 215501.


Fig. 1: Crystallographic simulation of Ge encapsulated inside a CNT and an image of the encapsulated Ge with an overlay of the simulated system.

Fig. 2: Crystallographic simulations of encapsulated SbTe, below is the captured structure with the simulation superimposed. Both 5 layer and 4 layer SbTe systems are included.

Type of presentation: Poster

MS-2-P-2592 Structural transformations in electrospun Pt-decorated carbon nanofibers

Zhigalina O. M.1, Zhigalina V. G.1, Ponomarev I. I.2, Razorenov D. Y.2, Ponomarev I. I.2, Kiselev N. A.1
1Shubnikov Institute of Crystallography RAS, Moscow, Russia, 2Nesmeyanov Institute of Organoelement Compounds RAS, Moscow, Russia
zhigal@crys.ras.ru

It is a well-known fact that good support material must have a high surface area which disperses the nanoparticles over, be porous enough to transmit gas flow and have a good electrical conductivity. Therefore nowadays carbon nanofibers (CNFs) are considered to be very promising support materials [1]. Recently electrospinning, a simple, inexpensive technique, has attracted significant attention in the preparation of CNFs [2]. Thin catalyst layers and their supporting substrates are the most critical components and up to now the problem of their optimisation is far from trivial.
In this investigation we have studied a possibility of electrospun CNFs carbonization by different methods such as high temperature annealing, chemical treatment and carbonization initialized in the metal particles presence. Pt-decorated CNFs were obtained using different catalytical coating and treatment techniques. To improve the CNF porosity and the process of catalytic nanoparticles deposition the CNFs were obtained on a base of polyvinylpyrrolidone (PVP) and polyimide (PI) polymer mixture and annealed at T = 1200 and 2800 °С [3].
The specimens for TEM investigations were prepared by the dispersion of CNF mats in acetone using an ultrasonic bath to get single fibers and to separate the bundles. These suspensions were dropped onto the Cu lacey carbon grids. The samples structure was characterized by a high resolution transmission electron microscopy (HRTEM) in a FEI Tecnai G2 30ST with SAED, EDX analysis and HAADF STEM detector and a FEI Titan 80-300 with probe Cs corrector at an accelerating voltage of 300 kV.
It was shown that PVP/PI treatment with annealing at 250 оС or at 1200 оС stimulated the fiber carbonization process without changes to their morphology or surface destruction, but the carbonization was incomplete. High temperature annealing at Т = 1200 и 2800 оС led to full fiber carbonization. In the structure of fibers the straight graphene planes were mainly observed (Fig. 1). In the case of combined high temperature (1200 оС) and chemical (PVP+PI) treatment the carbon nanofibers consisted of “curly” graphene planes within the whole fiber space (Fig. 2). In this case for the fibers carbonized in the Fe particles presence the same structure was revealed as well. As a result of this process the fibers’ surface became porous which promoted platinum nanoparticles to form a thick layer on the fibers’ surface (Fig. 3-4).


The investigation was partially carried out using IC RAS Research Centre and NRC “Kurchatov institute” equipment and supported by RFBR grant ofi-m-11-03-12115.

Fig. 1: HRTEM image of a carbonized nanofiber with straight graphene planes.

Fig. 2: TEM image of a carbonized nanofiber with “curly” graphene planes.

Fig. 3: TEM image of CNFs covered by Pt nanoparticles.

Fig. 4: HRTEM image of Pt nanocrystals.

Type of presentation: Poster

MS-2-P-2607 Nitrogen-Doped Graphene/Carbon Nanotube Hybrids: In-Situ Formation on Bifunctional Catalysts and Their Superior Electrocatalytic Activity for Oxygen Evolution/Reduction Reaction

Tian G.1, Zhao M.1, Yu D.2, Kong X.1, Huang J.1, Zhang Q.1, Wei F.1
1Beijing Key Laboratory of Green Chemical Reaction Engineering and Technology Department of Chemical Engineering, Tsinghua University, Beijing, China, 2School of Chemical and Biomedical Engineering, Nanyang Technological University, Singapore
tian-gl10@mails.tsinghua.edu.cn

There is a growing interest in oxygen electrode catalysts for oxygen reduction reaction (ORR) and oxygen evolution reaction (OER), as they play a key role in a wide range of renewable energy technologies such as fuel cells, metal-air batteries, and water splitting. Nevertheless, the development of highly-active bifunctional catalyst at low cost for both ORR and OER still remains a huge challenge. Herein, we report a new N-doped graphene/single-walled carbon nanotube (SWCNT) hybrid (NGSH) material as an efficient metal-free bifunctional electrocatalyst for both ORR and OER. NGSHs were fabricated by in situ doping during chemical vapor deposition growth on layered double hydroxide derived bifunctional catalysts. Our one-step approach not only provides simultaneous growth of graphene and SWCNTs, leading to the formation of three dimensional interconnected network, but also brings the intrinsic dispersion of graphene and carbon nanotubes and the dispersion of N-containing functional groups within a highly conductive scaffold. Thus, the NGSHs possess a large specific surface area of 812.9 m2 g-1 and high electrical conductivity of 53.8 S cm-1. Despite of relatively low nitrogen content (0.53 at%), the NGSHs demonstrate a high ORR activity, much superior to two constituent components and even comparable to the commercial 20 wt% Pt/C catalysts with much better durability and resistance to crossover effect. The same hybrid material also presents high catalytic activity towards OER, rendering them high-performance cheap catalysts for both ORR and OER. Our result opens up new avenues for energy conversion technologies based on earth-abundant, scalable, metal-free catalysts.


This work was supported by National Basic Research Program of China (973 Program, 2011CB932602) and Natural Scientific Foundation of China (No. 21306102).

Type of presentation: Poster

MS-2-P-2694 TEM Characterization of ALD-grown TiO2 on CNT

Zhang Y.1, Utke I.2, Erni R.1
1Electron Microscopy Center, EMPA, Swiss Federal Laboratories for Materials Science and Technology, Überlandstrasse 129, CH-8600 Dübendorf, Switzerland, 2Laboratory of Mechanics of Materials and Nanostructure, EMPA, Swiss Federal Laboratories for Materials Science and Technology, Feuerwerkstrasse 39, CH-3602 Thun, Switzerland
yucheng.zhang@empa.ch

A thriving field in nanotechnology is to develop synergetic functions of nano-materials by taking full advantages of unique properties of each component. In this context, combining TiO2 nano-crystals and carbon nanotubes (CNTs) offers enhanced photo-sensitivity and improved photo-catalysis efficiency, which is crucial to achieving sustainable energy and preventing environment pollution and hence has aroused a tremendous research interest. Despite progress in synthesis and performance of the material system, further research is required to understand some fundamental aspects of the material system, such as how TiO2 nucleates and grows on CNTs, and what is the bonding at the TiO2-CNT interface. Answers to these questions also help to design nano-composites based on CNTs and metal/metal-oxides with novel functionalities.
In this work an atomic layer deposition (ALD) technique has been adopted to grow TiO2 nano-particles on multiwall-CNTs (MW-CNTs). Control of the crystallinity, particle size and morphology of TiO2 can be obtained through deposition parameters adopted in ALD and a surface pre-treatment of MW-CNTs using O2 plasma. Transmission electron microscopy (TEM) has been very useful to characterize the ensemble structurally, chemically and electronically. In particular, electron energy loss spectroscopy (EELS) in the scanning TEM (STEM) mode has been employed to study C-K and Ti-L2,3 edge fine structures in TiO2, CNTs and their interface, in order to shed light on the mechanism of nucleation and growth of TiO2 on CNTs, as well as the interfacial bonding of the ensemble.


The author Dr. Y Zhang would like to thank Marie Curie Cofund action for financial support

Fig. 1: TEM micrographs of TiO2 deposited on CNT at 200°C for various ALD numbers of cycles: (a), (c) and (e) are after 20, 200 and 750 cycles respectively, without plasma pre-treatment; (b), (d) and (f) are 20, 200 and 750 cycles respectively, with the CNT subjected to O2 plasma pre-treatment. The insets show the corresponding diffraction patterns.

Fig. 2: Core-loss EELS spectra show the Ti_L2,3 and O_K edges of TiO2 on CNT after 20 ALD cycles with and without Oplasma pre-treatment. Difference in the near edge fine strucuture indicates different crystallinity of TiO2.

Type of presentation: Poster

MS-2-P-2699 Multi-scale investigations of nitrogen doping in graphene

Alloyeau D.1, Riccardi E.1, Lagoutte J.1, Ricolleau C.1, Wang G.1, Gallais Y.1
1Laboratoire Matériaux et Phénomènes Quantiques, Université Paris Diderot - CNRS, Paris, France
damien.alloyeau@univ-paris-diderot.fr

The electronic, thermal, and mechanical properties of graphene are exceptionally sensitive to lattice imperfections, surface functionalization and doping. Therefore, atomic scale structural and electronic investigations in this material are critically important for understanding these properties. Graphene samples produced by CVD method were doped with nitrogen by plasma exposure. We have exploited the complementarities of aberration-corrected TEM, Scanning Tunnelling Microscopy (STM) and micro-Raman spectroscopy to investigate the link between the structural and the electronic properties of N-doped graphene. Our experimental protocol allows applying these characterization techniques on the same samples in order to reduce the gap between micro and atomic scales investigation.

STM and HRTEM were used to characterize the nitrogen-induced single-point defects in graphene and the charge redistribution due to chemical bonding. As previously reported [1], the charge redistribution due to the insertion of nitrogen atoms in graphene that is easily detected by STM, allows the detection of such a low-contrast defect by HRTEM (Fig. 1). Our study highlights two important structural information about N-doped graphene and doping process. At first, Cu-supported graphene during the plasma exposure are more likely to be N-doped than suspended graphene. Secondly, the high variability of the C/N ratio on the same graphene sample reveals that nitrogen doping is not spatially homogeneous. This latter result pushed us to combine HRTEM and micro-Raman investigations on same micron-large areas of the samples, in order to provide a deeper understanding of the Raman spectrum as a function of the structure (holes, number of layers) and the Nitrogen doping rate of graphene (Fig. 2). 

[1] Meyer et al. Nature materials, 10, 209 (2011)


We are grateful to Region Ile-de-France for convention SESAME E1845, for the support of the JEOL ARM 200F electron microscope installed at the Paris Diderot University.

Fig. 1: HRTEM image of single layer graphene before nitrogen doping (a). Atomic scale analysis of nitrogen insertion in graphene: the charge redistribution due to chemical bonding is observed by aberration-corrected TEM (b) and STM (simulation of the structure and charge distribution in insert) (c).

Fig. 2: Raman spectra of the G-band optical phonon in two different spots of a Nitrogen-doped suspended graphene sample: the shift of the phonon peak energy (ωG) and its broadening show a significant variation of the chemical potential due to different Nitrogen doping.

Type of presentation: Poster

MS-2-P-2786 Bilayer graphene structures formed by passage of current through graphite: HRTEM and HAADF-STEM studies

Harris P.1, Slater T.2, Haigh S.2, Hage F.3, Kepaptsoglou D.3, Ramasse Q.3
1Electron Microscopy Laboratory, J.J. Thomson Building, University of Reading, Reading, RG6 6AF, UK, 2School of Materials, The University of Manchester, Manchester, M13 9PL, UK, 3SuperSTEM Laboratory, SciTech Daresbury, Keckwick Lane, Daresbury, WA4 4AD, UK
thomas.slater-5@postgrad.manchester.ac.uk

The subject of this paper is a new form of carbon which can be formed by passing an electric current through graphite [1,2]. This new carbon apparently consists of hollow graphitic shells bounded by curved and faceted planes, typically made up of two graphene layers. We describe studies of this carbon using high resolution transmission electron microscopy (HRTEM) and high-angle annular dark-field scanning transmission electron microscope imaging (HAADF-STEM). These studies appear to confirm that the bilayer graphene structures are 3-dimensional.
The carbon was prepared in an arc-evaporator which is normally used for coating specimens for scanning electron microscopy. Following evaporation, a small deposit was formed in the area where the two graphite electrodes made contact, and it was this deposit which contained the “transformed” carbon.
Some conventional TEM images of the transformed carbon are shown in Fig. 1. In the low magnification image (Fig. 1(a)), it can be seen that the outline of the structure is much more irregular than in normal graphite, with many curved and unusually-shaped features. Higher magnifications images, such as Fig. 1(b), show that the transformed carbon consists largely of bilayer graphene.
In order to determine the 3-dimensional shapes of the graphene structures we have used HAADF-STEM imaging. Both individual images and tilt sequences have been analysed. Individual HAADF-STEM images were recorded on an aberration-corrected Nion UltraSTEM100, operated at 60kV. Figure 2(a) shows a HAADF-STEM image of a region in which a nanotube is joined to a larger bilayer structure. The contrast in this image, in combination with a quantitative analysis of the near edge fine structure of the C K EELS edge [4], indicate that the edges of the structure are highly curved. This is consistent with the 3-dimensionality of this material.
Tilt series were recorded using an FEI Titan microscope operated at 80kV. A typical tilt sequence is shown in Fig. 2(b). This appears to show a 3-dimensional particle with the shape of a flattened cone.
Structural transformation of graphite as a result of the passage of an electric current has been observed by other groups [e.g. 5,6]. These groups have discussed the process in terms of the sublimation and edge reconstruction of flat graphene. However, as argued here, there are good reasons for believing that the transformed carbon is in fact 3-dimensional. If this is correct, this new carbon may have a number of possible applications, for example in supercapacitors.

[1] PJF Harris, J. Phys.: Condens. Matter 21 (2009), 355009.
[2] PJF Harris, Carbon 50 (2012) p.3195.
[3] PJF Harris et al., in preparation.
[4] XT Jia et al., Science 323 (2009) p.1701.
[5] JY Huang et al., PNAS 106 (2009) p. 10103.


The authors gratefully acknowledge funding from EPSRC, HM Government and the USA Defense Threat Reduction Agency.

Fig. 1: Conventional HRTEM images showing structure of carbon following passage of current.

Fig. 2: HAADF-STEM images of structures in transformed carbon. (a) Image showing junction between bilayer nanotube and larger region, (b) tilt sequence of approximately conical structure.

Type of presentation: Poster

MS-2-P-2825 Band Gap Expansion and Low-Voltage Induced Crystal Oscillation in Low-Dimensional Tin Selenide Crystals

Sloan J.1, Carter R.2, Suyetin M.3, Dyson M. A.1, Trewhitt H.1, Liu Z.4, Suenaga K.4, Giusca G.5, Kashtiban R. J.1, Bell G.1, Bichoutskaia E.3
1Department of Physics, University of Warwick, Coventry, CV4 7AL, UK, 2Department of Materials, University of Oxford, South Parks Road, OX1 3PH, UK, 3School of Chemistry, University of Nottingham, Nottingham, NG7 2RD, UK, 4National Institute of Advanced Industrial Science and Technology (AIST), Tsukuba, 305-8565, Japan, 5National Physical Laboratory, Teddington, TW11 0LW, UK
j.sloan@warwick.ac.uk

SnSe forms all-surface two atom-thick low dimensional crystals when encapsulated within single walled nanotubes (SWNTs) with diameters < 1.4 nm. Density Functional Theory (DFT) studies indicate that low-dimensional SnSe crystals typically undergo band-gap expansion. In slightly wider diameter SWNTs (~1.4-1.6 nm), we observe that three atom thick low dimensional SnSe crystals undergo a previously unobserved form of a shear inversion phase change resulting in two discrete strain states in a section of curved nanotube (not shown here). Under low-voltage (i.e. 80-100 kV) imaging conditions in a transmission electron microscope, encapsulated SnSe crystals undergo longitudinal and rotational oscillations, possibly as a result of the increase in the inelastic scattering cross-section of the sample at those voltages. Initial AC-TEM images were obtained at 100 kV using a JEOL 2010F transmission electron microscope fitted with a CEOS aberration corrector for which Cs was tuned to ~0.001 mm. Additional AC-TEM images were obtained at 80 kV were obtained on a JEOL JEM-ARM200F fitetd with an imaging corrector for which Cs was tuned to ~0.001 mm. 

Inside the narrower 1-1.4 nm SWNTs, we observed bilayer 2×2 SnSe nanocrystals (Fig. 1) and obtained images effectively viewed parallel to <001> relative to an ideal 2x2 rocksalt structure fragment. Systematic measurements of the lateral spacings of these encapsulated SnSe nanorods (Fig. 1(b)) relative to the centre point of the SWNT wall indicate that the obtained microstructure is undistorted and does not deviate significantly from the idealised 2x2 structure. A 2x2x6 atomic layer Sn12Se12 cluster based on average lattice spacings from the AC-TEM images (Fig. 1(b)) was DFT optimised and the resulting model used to generate a mutlislice image simulation from this model embedded in a (9,9) SWNT (Fig. 1(c)). DFT computed densities of states (DOS) for both bulk and 2x2 SnSe nanocrystals revealed that the latter have an expanded band gap of ca. 1.41 eV relative to the corresponding bulk bandgap of the rocksalt form (i.e. 0.68 eV). We had intended to perform higher resolution studies on the embedded SnSe nanocrystals using exit wave reconstruction from focal series of images but found that at 80 and 100 kV accelerating voltages, the embedded SnSe nanocrystals oscillate inside the encapsulating SWNTs (Fig. 2(a) and (b)).    

[1] Carter et al. Dalton Trans. 2014, published online DOI: 10.1039/c4dt00185k. 


J.S. and R.J.K. acknowledge the Warwick Centre for Analytical Science (EP/F034210/1). Z.L. and K.S. acknowledge JST-CREST and MEXT (19054017). E.B. acknowledges an ERC Starting (Consolidator) Grant, an EPSRC Career Acceleration Fellowship, and an EPSRC Research Leaders Award (EP/G005060).

Fig. 1: (a) AC-TEM image of (2 × 2)SnSe@SWNT. (b) Enlargement with dots indicating the centres of the Sn–Se columns and SWNT wall. (c) Second enlargement from (a) with overlaid multislice image simulation. (d) Side-on view and (b) end-on view of the experimental structure model.

Fig. 2: (a) This sequence of images obtained over ~12s at reveals two different modes of oscillation of 2 × 2 SnSe in a ~1 nm diameter SWNT. (b) Simulations and models of different rotational states of a 2 × 2 SnSe fragment in an (8,8) SWNT relative to an “ideal” <100> orientation (i.e. bottom right).

Type of presentation: Poster

MS-2-P-2850 Influence of charge carriers density on flexural phonon spectrum in graphene measured by electron diffraction

Kirilenko D. A.1
1Ioffe Institute
zumsisai@gmail.com

Graphene is a specific form of matter – an electronic membrane [1,2,3]. It is known that free-standing graphene undergoes severe corrugation [4]. And the main reason for this is flexural phonons. At the same time, structural fluctuations in graphene are influenced by charge carriers, which are easily generated in gapless graphene. It gives rise to a complicated phenomenon, so that electronic properties of graphene are defined by the interplay between charge carries and lattice distortions. This interplay enormously complicates the theory of transport in graphene [5]. For instance, full calculations of thermal dependence of graphene’s resistivity are still to be completed [6].

Previously, a technique for measuring of the flexural phonon spectrum basic parameters in free-standing graphene was presented [7], which was later expanded to measuring of the full spectrum profile. The technique uses electron diffraction obtained in transmission electron microscope (TEM) to scan the reciprocal lattice of graphene that gives information on its structural distortions. The feature of the electron diffraction imaging is that provides information on rapidly varying structural distortions, what is inaccessible by most of other techniques. It is remarkable, that electron beam of TEM does generate charge carriers in graphene. Variation of the electron beam intensity changes the generated charge carriers density. This allows measuring the changes of the flexural phonons spectrum related to the influence of charge carriers.

A significant dependence of the small-wavevector (long undulations) part of the flexural phonons spectrum has been found. Whereas, the large-wavevector part seems to avoid the charge carriers influence. In the accessible range of electron beam densities and, thus, generated charge carriers densities, the amplitude in the left part of the spectrum is being successively suppressed with increasing charge carriers density. That is, graphene becomes noticeably smoother at large scale. It must influence graphene’s resitivity at increased densities (or bias voltage in graphene-based devices).

Finally, influence of the charge carriers density on corrugation of suspended graphene has been measured and degree of the specific electron-phonon coupling estimated.

1. E.-A. Kim and A. H. Castro Neto, Europhysics Letters 84 (2008), 57007.
2. D. Gazit, Phys. Rev. B 80 (2009), 161406(R).
3. P. San-Jose, J. Gonzalez and F. Guinea, Phys. Rev. Lett. 106 (2011), 045502.
4. J.C. Meyer, A.K. Geim et al., Nature 446 (2007), p. 60.
5. M. Gibertini, A. Tomadin et al., Phys. Rev. B 85 (2012), 201405(R)
6. S. Das Sarma, S. Adam et al., Rev. Mod. Phys. 83 (2011), p. 407.
7. D.A. Kirilenko, A.T. Dideykin and G. Van Tendeloo, Phys. Rev. B. 84 (2011), 235417.


This work was supported by Russian Foundation for Basic Research.

Type of presentation: Poster

MS-2-P-2878 Imaging and Spectroscopy of Graphene/Hexagonal Boron Nitride Lateral Heterostructure Interfaces

Basile L.1,2, Liu L.3, Gu G.3, Vlassiouk I.2, Lupini A. R.2, Unocic R. R.2, Idrobo J. C.2
1Escuela Politécnica Nacional, Quito, Ecuador, 2Oak Ridge National Laboratory, Oak Ridge, USA, 3The University of Tennessee, Knoxville, USA.
lbasilec@gmail.com

Boundaries or in-plane interfaces in two-dimensional (2D) materials will play a critical role in future device applications. For example, electronic and mechanical properties are affected by structure, chemistry, morphology, and location of a grain boundary [1,2]. By using microscopic tools, we recently demonstrated lateral coherence in an in-plane heterostructure of graphene and hexagonal boron nitride (BN) on samples staying on the growth substrate [3]. An atomic-resolution STEM can be an ideal tool to image the presumably atomically sharp graphene-BN interface, but contaminations introduced during the transfer process hinder its direct observation.

In this study, we examined the contaminant-covered graphene-BN boundary using EELS in an aberration-corrected STEM, Nion UltraSTEM 100, equipped with a cold field emission electron source, a corrector of third and fifth order aberrations, and a Gatan Enfina spectrometer [4]. To avoid graphene and BN knock-on damage we operated the microscope at an acceleration voltage of 60 kV. A convergence semi-angle of 30 mrad, and 54 to 200 mrad collection semi-angles were used to obtain the medium angle annular dark field (MAADF) images. The EEL spectrum maps were collected with an energy resolution of ~350 meV.

Fig. 1(A) shows an experimental MAADF image of the graphene-BN boundary. Fast Fourier Transform of the areas shown in (A) indicates that the BN is aligned with the graphene monolayer. Evidence of a sharp interface is provided by the chemical map shown in (E), where the boron K-edge clearly defines a sharp graphene-BN interface.

Fig. 2 shows intensity profiles along the yellow lines of Fig 1. The left panel of Fig 2 shows that the graphene-BN boundary is composed of monolayer graphene and monolayer BN. The right panel shows a transition width of 0.5 nm between graphene and BN as determined from the boron K-edge signal. The transition width is defined from 25% to 75% of the values of the boron K-edge signal across the graphene-BN interface.

Direct observation of a boundary at atomic resolution requires a reliable method to free the graphene-BN interface of contaminants. We will discuss our current efforts on removing contaminants by in-situ annealing, thus revealing the buried graphene/BN interface. Our preliminary results indicate that regions of thousands of nanometer squares of clean graphene are produced during in-situ annealing. The method opens the door for the study of the long-range structure of 2D lateral heterostructure interfaces at the atomic scale.
[1] Adam W. Tsen, et al., Science 336 (2012), 1143
[2] Gwan-Hyoung Lee, et al., Science 340 (2013), 1073
[3] L Liu et al, Science 343 (2014), 163
[4] OL Krivanek et al, Ultramicroscopy 108 (2008), 179


National Secretariat of Higher Education, Science, Technology and Innovation of Ecuador (LB), NSF, the Defense Advanced Research Projects Agency (LL & GG), Office of Basic Energy Sciences, U.S. Department of Energy (DOE) (ARL), Center for Nanophase Materials Sciences, Office of Basic Energy Sciences, U. S. DOE (RRU, JCI).

Fig. 1: (A) MAADF image of a buried graphene-BN boundary. (B) and (C) fast Fourier transforms of the highlighted areas shown in (A). (D) and (E) simultaneously acquired STEM image and EEL spectrum map of the region shown in (A), respectively. The blue/white dashed lines indicate the graphene-BN boundary. Scale bars are 5 nm. Adapted from Ref. [3].

Fig. 2: (A) Intensity profile along the yellow line in Fig. 1(A). The intensity profile indicates that the boundary is composed of monolayer graphene and BN. (B) Boron K-edge intensity profile along the yellow line in Fig. 1(E). The boundary is sharp with a transition width of 0.5 nm. The spatial resolution is 0.5 nm. Adapted from Ref. [3].

Type of presentation: Poster

MS-2-P-2888 Wetting behavior of ionic liquid on a carbon nanotube

Imadate K.1, Hirahara K.1
1Osaka University, Osaka, Japan
hirahara@mech.eng.osaka-u.ac.jp

It is interesting question how a single nanomaterial such as carbon nanotube gets wetting by liquid. Wettability of materials is generally determined by the balance of interface tensions acting at air-liquid, liquid-solid, solid-air boundaries, but nanometer-scale morphologies of solid surfaces often cause anomalous wetting behavior. Regarding carbon nanotubes, they have extremely high curvature surfaces due to their cylindrical shapes with nanometer scale diameters. In this study, we investigated wetting behavior on a single CNT by in-situ electron microscopy. Prior to the experiment, CNT probe was prepared by using TEM-STM holder for nanomanipulation in a transmission electron microscope (TEM, JEM-2500SE) at 90kV acceleration voltage. A multi walled CNT was attached to the tip of cantilever probe used for scanning probe microscopy. On the other hand, ionic liquid was used as a liquid specimen, since it is rather stable in vacuo due to extremely low evaporation pressure. It was supported as the liquid level was parallel to the incident beam direction on a specimen stage of nanomanipulator. Tip of the CNT was then approached to the ionic liquid from the normal direction to the liquid level, and the series of images were recorded as movies at the moment when the tip touched to the ionic liquid. As the result, meniscus formed at the contact region, and a thin film with 3nm thickness simultaneously formed to cover entire the CNT. The contact angle measured at the meniscus was almost zero. These results indicate that CNT shows autophobic wetting, although macroscale droplet on a plane graphite surface shows about 25˚ contact angle. In addition, similar experiments were performed in a scanning electron microscope, and attractive wetting forces were measured on the basis of Wilhelmy method for CNTs with 5~15 nm diameters. Measured values indicated a tendency to be greater than the expected values from Wilhelmy equation representing the correlation of the force, tube diameter and surface tension. Instead, fitted curve to the experimental data showed the increment of effective diameter of cylindrical sample. The corrected value was 2.84nm, which is consistent to the thickness of liquid film formed on CNT. Accordingly, the wetting behavior observed in the present study can be explained by considering that the liquid film acts as a part of a solid cylinder, which suggests a possibility that liquid molecules are rather strongly constrained on the CNT surface.


Ionic liquid used in this study was provided by T. Tsuda in Osaka University.

Fig. 1: SEM images of a CNT probe before and after contacting to the ionic liquid. We can see that the CNT is pulled into the ionic liquid due to the attractive wetting force. For this case, the force was measured by 1.2nN.

Type of presentation: Poster

MS-2-P-2898 Measurement of the tensile force applied to a carbon nanotube during the axial shrinking deformation

Hirahara K.1, Nishiyama Y.1
1Osaka University, Osaka, Japan
hirahara@mech.eng.osaka-u.ac.jp

We have studied on deformation process of a bridged carbon nanotube (CNT) during Joule heating by in-situ transmission electron microscopy (TEM). Many papers reported that a CNT got thin or cut at the central portion due to or cutting and reconnecting of bonds or sublimation of carbon atoms during the Joule heating [1~4]. In this study, another type of deformation was observed, namely shrinking deformation along the axial direction of the CNT. We found that such a shrinking was observed when the CNT was bridged between rather frexible electrodes, namely the CNT could change its length during the heating. This result suggested that these deformation process strongly depended on how release the tensile stress applied to the CNTs caused by sublimation of carbon atoms during the Joule heating. Therefore we measured the stress loaded to the CNT during the Joule heating. For the measurement, a cantilevered probe for scanning probe microscopy was used as the flexible electrode, which spring constant was 0.02~0.41 N/m. A single or double wall CNT is bridged between the cantilevered probe and a Pt/Si substrate by operating a nanomanipulation holder (TEM-STM system, nanofactory) in TEM (JEM-2500SE, 90kV). Current is then applied to the CNT, and its deformation process was recorded. When the CNT began shrinking, cantilevered probe underwent deflection and made balance to the tensile force, so that the tensile force applied to the CNT were able to be measured by monitoring the deflection of cantilevered probe. In this system, tensile stress applied to the CNT gets increase as the degree of shrinking increased, and the CNT finally cut or detached from the electrode. Experimental results revealed that the shrinking deformation of CNTs occurred with loading tensile stress under 1.9 N/m2. It is also suggested that such a shrinking deformation was promoted when topological defects formed; carbon atoms may selectively evaporate from such defective site.

[1] H. Maruyama, et al., Appl. Phys. Ex. 3, (2010) 025101.
[2] T. D. Yuzvinsky, et al., Nano Lett. 6 (2006) pp. 2718-2722.
[3] J. Y. Huang, et al., Nature 439 (2006) pp. 281.
[4] K. Hirahara et al., Appl. Phys. Lett. 97 (2010) 051905.


Fig. 1: A series of TEM images showing shrinking deformation of a carbon nanotube. We can see that the cantilevered probe underwent deflection by pulling. Initial length of  this nanotube was 239nm, and shrank about 100 nm at the botom image.

Type of presentation: Poster

MS-2-P-2925 Probing the dynamics of structure defects and chemical dopants in monybdenum disulfide monolayer at elevated temperature by Cs-corrected STEM

Jin C.1, Lv D.1, Hong J.1
1State Key Laboratory of Silicon Materials, Key Laboratory of Advanced Materials and Applications for Batteries of Zhejiang Province, Department of Materials Science and Engineering, Zhejiang University, Hangzhou, Zhejiang, PR China.
chhjin@zju.edu.cn

As a representative family member of the emerging two-dimensional transition metal dichalcogenides (TMDCs), atomically thin molybdenum disulfide has attracted intensive research efforts owing to its unique structural and electronic properties that has promised a wide application in future nanoelectronic and optoelectronic devices [1-3]. Since defects plays an important role on tailoring the physical and chemical properties of any semiconductors, molybdenum disulfide is not an exception. Therefore it is very important to resolve the structure defects and figure out the impact of these defects on the physical properties of molybdenum disulfide. [4-6]

In this talk, we will present our latest progress on studying the atomic defects and dopants in molybdenum disulfide monolayers by aberration-corrected STEM. Furthermore, with the assistance of MEMS-heating technique, the adsorption, migration and coalescence of structural defects at elevated temperatures can be directly observed in situ. The absorption sites, diffusion pathways and the associated activation energy are experimentally determined experimentally, which are further supported by the DFT calculations.

Reference:

[1] K. F. Mak, C. Lee, J. Hone, J. Shan, and T. F. Heinz, Physical Review Letters 105, 136805 (2010).

[2] A. Splendiani et al., Nano Letters 10, 1271 (2010).

[3] Q. H. Wang, K. Kalantar-Zadeh, A. Kis, J. N. Coleman, and M. S. Strano, Nature Nanotechnology 7, 699 (2012).

[4] P. Komsa, J. Kotakoski, S. Kurasch, O. Lehtinen, U. Kaiser, and A. V. Krasheninnikov, Physical Review Letters 109, 035503 (2012).

[5] A. M. van der Zande et al., , P. S. Huang et al., Nature Materials 12, 554 (2013).

[6] W. Zhou et al., Nano Letters 13, 2615 (2013).


The work on microscopy is done in the EM Center of ZJU. This work is financially supported by the NSFC (51222202,), the National Basic Research Program of China (2014CB932500), the Program for Innovative Research Team in University of Ministry of Education of China (IRT13037) and the Fundamental Research Funds for the Central Universities (2014XZZX003-07).

Type of presentation: Poster

MS-2-P-2934 Probing band structures of atomically thin molybdenum disulfide by EELS

Hong J.1, Li K.2, Jin C.1, Zhang X.2, Yuan J.3, Zhang Z.1
1State Key Laboratory of Silicon Materials, Key Laboratory of Advanced Materials and Applications for Batteries of Zhejiang Province, Department of Materials Science and Engineering, Zhejiang University, Hangzhou, Zhejiang 310027, P. R. China., 2Advanced Nanofabrication, Imaging and Characterization Core Lab, King Abdullah University of Science and Technology (KAUST), Thuwal 239955, Kingdom of Saudi Arabia, 3Department of Physics, University of York, Heslington, York, YO10 5DD, United Kingdom
jinhuahong436@gmail.com

In recent years, semiconducting MoS2 has attracted much public attention because of its hexagonal structure, proper bandgap (1.3~1.8eV) and potential application in nanoelectronics and valleytronics. Currently a clear and complete picture of bandgap transition (<2eV), higher interband transition (~5eV) and plasmon resonance (~23eV) associated with the thickness-dependent electronic structure is still lacking. In this talk, we will present the EELS study on the electronic structures of atomically thin MoS2.
        We use a spherical aberration corrected TEM (FEI Titan Cube) to conduct angle resolved EELS measurement. This microscope is equipped with a monochromator providing an energy resolution of 0.14eV which help us resolve fine structures of low loss EEL spectrum and obtain band structure of MoS2. A transition from indirect to direct gap is illustrated as the thickness decreases down to monolayer. Other strong interband transition peaked at 3.1eV and 4.5eV and high-energy π+σ Plasmon excitation at 23eV are also presented as a function of thickness and momentum transfer q. These excitations (not easily accessible by conventional optical characterization) in atomically thin MoS2 are reported for the first time. Their energy redshift with the decreasing thickness and monotonically-increasing linewidth dispersion with q indicate the spilling-out effect and Landau damping, respectively, in this low dimensional electron gas system. Our investigation provides a successful paradigm to depict the electronic structures of any other novel transition metal dichalcogenides (TMDs).


This work on microscopy was carried out in the Imaging and Characterization Core Lab of KAUST in Saudi Arabia. This work is financially supported by the National Science Foundation of China (51222202,), the National Basic Research Program of China (2014CB932500), the Program for Innovative Research Team in University of Ministry of Education of China (IRT13037).

Fig. 1: Figure 1 (a) The primitive cell of monolayer MoS2 with lattice basic vector a1 and a2. Purple atom: Mo; yellow: S2. (b) Reciprocal lattice (electron diffraction pattern)  and first Brillouin zone. (c) The scattering geometry of angle resolved EELS . (d) Corresponding angle-resolved spectrum profile with qy along ΓM direction in (c). 

Type of presentation: Poster

MS-2-P-2938 HRTEM studies of Bamboo-like nanotubes found in the carbonaceous chondrite Allende meteorite.

Rendon L.1, Cocho G.1, Cruz H.1, Ortega F.2, Reyes M.2, Buhse T.3, Garibay V.4, Santiago P.1
1Instituto de Física, Universidad Nacional Autónoma de México; Circuito de la Investigación Científica s/n. Ciudad Universitaria. C.P. 04510, México D.F, 2Instituto de Geología, Universidad Nacional Autónoma de México. Circuito de la Investigación Científica s/n. Ciudad Universitaria. C.P. 04510, México D.F., 3Facultad de Ciencias, Universidad Autónoma del Estado de Morelos. Av. Universidad 1001, Col. Chamilpa, 62209 Cuernavaca, Morelos, México., 4Instituto Mexicano del Petróleo. Eje Central Lázaro Cárdenas Norte 152 Col. San Bartolo Atepehuacan, C.P 07730, México.
paty@fisica.unam.mx

In February 8, 1969 a large carbonaceous chondrite meteorite fell in Allende Chihuahua, Mexico. Carbonaceous chondrites meteorites are very important because of their organic compounds and peculiar composition. Allende meteorite has large and abundant chondrules (mm-sized) in olivine matrix, large refractory inclusions, a low degree of aqueous alteration and graphitized carbon. Its large carbon content has represented an interesting source to the study of evolution and lineage of carbon chemistry, from nebular to current ages and has been related with prebiotic Earth, because its collisions and impacts with early Earth formation represent an organics extraterrestrial input. The organic composition in carbonaceous chondrites is diverse, and it is possible to mention as example kerogenic material macromolecular, sugar alcohols, ketones, amines, and amino acids. The nanostructures in this type of meteorites have been identified as fullerenes, carbon onions and the possible presence of carbon nanotubes (CNTs) has been suggested since 2006. Actually, inorganic serpentine nanotubes were described by Zega and co-workers in meteoritic matrix of carbonaceous chondrites, ranking in ~20nm diameter tube.

At high temperatures, carbon precursors are decomposed or evaporated and then condensed to build the sp2 graphite networks of CNTs. High temperatures are normally obtained from external heating, which is highly energy-consumed. Theoretically, such a problem can be solved by employing hugely exothermic reaction systems like the conditions of the early sun.

The sample was obtained from the collection of the Geology Institute, at the Autonomous National University of Mexico. In order to avoid the contamination of the meteorite sample, we drilled a hole with a steel laboratory spatula in the small piece of the meteorite. The powder obtained was supported in a microscopy glass slide and grinded with a second slide. The powder was supported in an electron microscopy grid to be observed in a JEM-2200FS TEM.
By HRTEM we observed bamboo-carbon nanotubes (BCNTs) in the meteorite sample as well as polyhedrical carbon structures. BCNTs can be thought as coaxial graphene sheets built of sp2 bonded. The tubes are concentric and coaxial. They also are highly defective and several bounds are broken, this fact promotes active bonds to act as chiral templates of other organic molecules.


Authors acknowledge the financial support from DGAPA-UNAM, through grant IN113411.

Fig. 1: Figure 1. a) BCNTs of about 20 nm wide. A highly defective structure is shown in the tubes indicating a chirality. b) A closer view of the coaxial structures.

Type of presentation: Poster

MS-2-P-3037 Topography mapping of ultrathin layered crystals

Dolle C.1, Niekiel F.1, Mittelberger A.1, Butz B.1, Spiecker E.1
1Center for Nanoanalysis and Electron Microscopy (CENEM), Friedrich-Alexander-Universität Erlangen-Nürnberg
christian.dolle@ww.uni-erlangen.de

Quasi-2D crystalline materials are widely investigated to explore their optical, electronic and mechanical properties. The most prominent examples in this class of materials are graphene, hBN, and recently dichalcogenites, e.g. MoS2.

The third dimension of those crystals may not be neglected since freestanding membranes are thermodynamically forced to form intrinsic ripples. Moreover, the resulting topography of such materials is expected to have a severe influence on their properties. For example, strain, which can alter the band structure, is caused by any change of the surface inclination. In graphene corrugations have already been confirmed by Meyer et al. applying electron diffraction [1]. Inspired by that study we developed a method to determine the topography of freestanding membranes by diffraction-contrast TEM imaging. The procedure is based on dark field (DF) tilt series using at least two independent g-vectors oriented perpendicular to the respective tilt axis. In those DF images, the measured intensity directly depends on the local excitation condition and thus on the local inclination of the membrane. By tilting, the reciprocal lattice rod (relrod) is scanned simultaneously in each sub-region of the DF images. To determine the inclination of each sub-region with respect to the specific g-vector, the maximum of the tilt-angle dependent intensity distribution is fitted. This is done for two different tilt series and the obtained data are used to calculate the absolute inclination of each sub-region and thus to determine the membrane topography.

We applied the procedure to freestanding membranes from high-quality epitaxial graphene on SiC [2]. Fig. 2a) depicts a representative {11-20} DFTEM image of such a few-layer graphene membrane. The local mean image intensity represents the number of graphene layers as proven by rocking curves. The sharp dark lines in the DF image are due to basal-plane partial dislocations, which have an additional impact on the local topography [3]. To demonstrate the strongly different intensity distributions along different directions Fig. 2e)-f) show exemplary DF images at 0 tilt for 3 independent {11-20} directions. It can be recognized that, while the wavy topography leads to strong, almost parallel contrast variations in the (11-20) and (1-210) images, the (2-1-10) DF image (with g perpendicular to the wave-direction) is less influenced.

While in the used example the basal-plane partial dislocations have a severe influence on the topography of the material, it will be shown that even the choice of the TEM support has a strong impact on the topography of defect-free membranes.

1Meyer et al., Solid State Commun. 2007, 143, 101

2Waldmann et al., ACS Nano 2013, 7, 4441

3Butz et al., Nature 2014, 505, 533


We acknowledge financial support by the Cluster of Excellence: Engineering of Advanced Materials and SFB 953: Synthetic Carbon Allotropes.

Fig. 1: a) Model of inclined membrane: Inclinations non-parallel to the used g-vector show intensity variations as indicated by the dark and light gray areas, b) Ewald sphere construction, c), d) enlargement for almost flat and strongly inclined membrane area

Fig. 2: a) Graphene membrane with 2-, 3- and 4-layer areas (scale bar 500 nm), b)-d) rocking curves extracted from the areas indicated (2, 3, 4 layers), e)-f) 3-layer graphene DF images obtained with the 3 indicated g-vectors, dotted line shows tilt axis orthogonal to the reflection used for imaging.

Type of presentation: Poster

MS-2-P-3075 Graphene-based materials and breast cancer cells

Ponticelli G. S.1, Colone M.1, Rago I.2, Sarto M. S.2, Stringaro A.1
1Italian National Institute of Health, 2Research Center on Nanotechnology Applied to Engineering Sapienza University of Rome
gennaro.ponticelli@guest.iss.it

Recent discoveries on graphene, a two-dimensional, crystalline allotrope of carbon, stimulated research on related structures, such as Graphite NanoPlatelets (GNPs), a 1-15 nm thick flake, constituted of 3-48 layers of graphene, obtained starting from Intercalated Graphite Compounds (GIC) via thermochemical exfoliation. These novel nanomaterials are providing fascinating opportunities for biotechnological development because of their unique structures, properties and possible applications.
Graphene and its derivatives are promising candidates for important biomedical applications because of their versatility. Due to the expanding applications of nanotechnology, human and environmental exposures to graphene-based nanomaterials are likely to increase in the future. However, the prospective use of graphene-based materials in a biological context requires a detailed comprehension of their toxicity.
Herein, we report on the interaction of stable and evenly dispersed exfoliated GNPs obtained using an ultrasonic bath for different times (30 min, 50 min and 70 min) with human breast adenocarcinoma cells (SKBR3 and MDA-MB-231) for 24 h. Biocompatibility of nanoplatelets has been evaluated by MTT (Fig. 1) while cell viability has been detected using Trypan Blue assays (Fig. 2). GNPs particles were more cytotoxic in SKBR3 than MDA-MB-231 cells suggesting a cell phenotype-dependent effect.
Furthermore, light microscopy observations (Fig. 3 and 4) and scanning electron microscopy analysis (data not shown) were used to gain understand on the mechanism of cell-nanoplatelets interaction. The bright-field images showed GNPs particles on SKBR3 and MDA-MB-231 cellular surfaces (see arrows).
Our results lead us to expect that efforts with interdisciplinary approaches among chemistry, biology, and engineering will accelerate mechanistic understanding of graphene-based platforms for bio and nanomedicine applications.


Fig. 1: GNPs (30 min, 50 min and 70 min) biocompatibility on SKBR3 and MDA-MB-231 cell lines by MTT test after 24 hrs of incubation.

Fig. 2: SKBR3 and MDA-MB-231 cell viability evaluation by Trypan blue assay after incubation with GNPs (30 min, 50 min and 70 min) for 24 hours.

Fig. 3: Bright-field microscopy image of SKBR3-nanoplatelets interaction. Cells were incubated for 24 hours with GNPs (arrows).

Fig. 4: Bright-field microscopy image of MDA-MB-231-nanoplatelets interaction. Cells were incubated for 24 hours with GNPs (arrows).

Type of presentation: Poster

MS-2-P-3097 TEM and EELS studies of structures obtained under the different conditions of the thermobaric treatment of C60+CS2

Tyukalova E. V.1,2, Perezhogin I. A.1, Kulnitskiy B. A.1,2, Blank V. D.1,2, Popov M. Y.1,2, Alekseev M. V.1,2
1Technological Institute for Superhard and Novel Carbon Materials, Troitsk, Moscow, Russian Federation , 2Moscow Institute of Physics and Technology State University, Dolgoprudny, Moscow Region, Russian Federation
elizavetatyukalova@gmail.com

In our work we carry out the transmission electron microscopy (TEM) and electron energy loss spectroscopy (EELS) studies of the samples obtained in a series of experiments on the thermobaric treatment of C60 with addition of CS2 in a diamond anvil cell with shear deformation ability. The main goal of the research is to obtain and characterize the new structures based on C60, because C60 is a very promising precursor for production of superhard materials. TEM study was done by a JEM-2010 microscope with GIF Quantum attachment for EELS. Features of radial distribution of pressure and shear in diamond anvil cell resulted in the coexistence of two types of structures in the samples: the crystalline and disordered phases. The crystalline phase represents itself slightly distorted fragments of the original face centered cubic (FCC) lattice fullerene (fig. 1). Apparently, these distortions appear due to the polymerization of C60 molecules.
Exact structure of disordered phases (fig. 2) has not been established by us, but according to our high-resolution images, it has inherited some elements of symmetry from the FCC lattice of fullerenes. The interplanar spacing in fig. 2 is about 0.35 nm, while the fragments of the “lattice” are deformed and disoriented one relatively to another. The microdiffraction and Fourier analysis have shown that two systems of fringes seen in fig. 2 intersect composing different angles in a range from 70° to 85°.
Figure 3 shows EELS spectra of the obtained specimens. The spectrum of C60 obtained under the pressure of 12-17GPa with shear deformation (fig. 3, (a)) correspond to the structures shown in fig. 1, and it is very close to the spectrum of the FCC lattice of the original C60 (fig. 3, (b)). The spectrum taken from the sample obtained at 25-30 GPa with shear deformation (fig. 3, (c)) and that from the sample obtained under 5 GPa at a temperature of 973° C (fig. 3, (d)) without shear, are almost identical.
At the same time the peaks of the all spectra in fig. 3 have the same positions, but the relative intensities of these peaks in (c) and (d) is different from those in (a) and (b). For example, the initial fullerene in (a) (and sample in (b)) has an absolute maximum at 300 eV, while in (c) and (d) it is about 292 eV. According to the literature data the small peaks at about 287eV correspond to the presence of molecular C60. Therefore, basing on our EELS data, we assume that the molecular C60 is present in the structures of both types in our samples, but the structure shown in fig. 2 is significantly different from the traditional fullerene FCC lattice.


Fig. 1: HRTEM image of the distorted FCC lattice of fullerene [110] zone axis. The interplanar spacing is slightly distorted in different areas, and the angle between planes in such fragments is not always exactly 70.5°.

Fig. 2: HRTEM image of the disordered phase. The interplanar spacing is about 0.346 nm, and the angles between the intersecting planes differ within the range from 70° to 85°.

Fig. 3: EELS spectra of: a) sample obtained at 12-17 GPa b) initial fullerene; c) sample obtained at 25-30 GPa with shear deformation; d) sample obtained at 5 GPa at temperature of 973° C

Type of presentation: Poster

MS-2-P-3149 Nitrogen Doped Single-Walled Carbon Nanotubes: Experimental and Theoretical Atomic-Resolved EELS Studies

Arenal R.1, 2, March K.3, Ewels C. P.4, Rocquefelte X.4, Kociak M.3, Loiseau A.5, Stephan O.3
1Lab. Microscopias Avanzadas (LMA), Inst. Nanociencia Aragon (INA), U. Zaragoza, Spain., 2ARAID Fundation, Zaragoza, Spain., 3Lab. Physique Solides (LPS), CNRS-U. Paris Sud, Orsay, France., 4Institut Matériaux Jean Rouxel (IMN), CNRS-U. Nantes, Nantes, France., 5Lab. Etude Microstructures (LEM), CNRS-ONERA, Châtillon, France.
arenal@unizar.es

Having access to the chemical environment at the atomic level of a dopant in a nanostructure is crucial for the understanding of its properties. A very good example in this context is the case of notably nitrogen-doped carbon nanotubes (CNx-NT) because their properties are significantly affected by the atomic arrangement of the dopant atoms in such nanostructures [1-4]. Thus the knowledge of this information requires precision measurements, combining high spatial resolution and high spectroscopic sensitivity. In order to achieve these goals, we have developed, for the first time, atomically-resolved EELS allowing us to detect individual N dopants in single-walled (SW) carbon nanotubes. These results have been compared with first principles calculations.

The STEM-EELS-experiments were performed in a NION UltraSTEM 200, operated at 60 kV. In parallel, HRTEM imaging studies have been performed using an imaging aberration-corrected FEI Titan-Cube microscope working at 80 kV.

Figure 1 displays a HAADF image of a CNx-SWNT where an EEL spectrum-image (SI) has been recorded in the red marked area of the image. Three single EEL spectra, extracted from this spectrum-image, in the marked positions/pixels of Fig. 1 (b) (spectra labelled (i), (ii) and (iii)), the 4th spectrum is the sum of (i) and (ii). The C-K edge is visible in the three spectra. In only two of the spectra of the whole dataset (1755 spectra), the nitrogen signal is also detectable. The nitrogen 1s (N1s) ELNES, expanded in Fig. 1 (c), show a strong peak at ~401 eV, with very little signal at energies above this. Comparing the spectra to density functional theory (DFT) ELNES calculations of possible single nitrogen defects, there is excellent agreement with the spectrum for substitutional nitrogen (Fig. 1 (c)(iii) and atomic model, Fig. 1(d)) across the range of π* and σ* bands. We have also investigated other more complex configurations that we will present and discuss in this contribution [5]. In summary, these studies elucidate a crucial question concerning the nature of the nitrogen atomic configuration of CNx-NTs. In fact, this detailed knowledge of how nitrogen atoms are incorporated in the carbon lattice as well as precisely control of their incorporation are required for the use of these NTs for future technological applications.

[1] R. Arenal, X. Blase, A. Loiseau, Advances in Physics 59, 101 (2010).
[2] P. Ayala, R. Arenal, A. Rubio, A. Loiseau, T. Pichler, Rev. Mod. Phys. 82, 1843 (2010).
[3] P. Ayala, R. Arenal, M. Rummeli, A. Rubio, T. Pichler, Carbon 48, 575 (2010).
[4] C.P. Ewels, M. Glerup, J. Nanosci. Nanotech. 5, 1345 (2005).
[5] R. Arenal, K. March, C.P. Ewels, et al., submitted.


The research leading to these results has received funding from the EU 7th Framework Program under Grant Agreement 312483-ESTEEM2 (I3) and from the French CNRS (FR3507).

Fig. 1: Fig. 1 (a) HAADF image of a CNx-SWNT. An EELS-SI has been recorded in the red area. (b) Selection of EEL spectra extracted from the SI, pixels outlined in the inset HAADF image acquired simultaneously with the SI. Each curve corresponds to a single spectrum from the SI, except the black, which is a sum of previous ones.

Fig. 2: Fig. 2(a) Simulated N1s ELNES (iii) & N partial DOS calculations ((i)purple=pz π*-states, (ii)green=px-y σ*-states) for substitutional N, compared to the experimental spectrum (iv) (Fig. 1(b-iii)). These simulations allow unambiguous assignment of the peak at ~401eV to a substitutional configuration shown in the DFT optimized structure, Fig. 2(b).

Type of presentation: Poster

MS-2-P-3164 STEM and EELS investigation of graphene nanoribbon epitaxially grown over SiC

Gloter A.1, Palacio I.2, Celis A.3, Nair M.2, Zobelli A.1, Sicot M.4, Malterre D.4, Nevius M. S.5, Berger C.5, de Heer W. A.5, Conrad E. W.5, Taleb-Ibrahimi A.1, Tejeda A.1,2
1Laboratoire de Physique des Solides, Université Paris-Sud, CNRS, UMR 8502, F-91405 Orsay Cedex, France, 2UR1 CNRS/Synchrotron SOLEIL, Saint-Aubin, 91192 Gif sur Yvette, France, 3Synchrotron SOLEIL, L’Orme des Merisiers, Saint-Aubin, 91192 Gif sur Yvette, France, 4Université de Lorraine, UMR CNRS 7198, Institut Jean Lamour, BP 70239, F-54506 Vandoeuvre-lès-Nancy, France, 5School of Physics, The Georgia Institute of Technology, Atlanta, Georgia 30332-0430, USA
gloter@lps.u-psud.fr

Graphene nanoribbons grown on the (1-10n) and (-110n) facets of SiC have demonstrated exceptional electronic properties as ballistic transport along their long direction and a band gap in the small direction [1]. In order to understand these electronic properties, we have performed (S)TEM (HAADF, LAADF, EELS) investigation in combination with STM and ARPES measurements. The (S)TEM have been performed on X-section sample as it can be schematically seen in the figure 1. Using Cs corrected STEM at 60 keV voltage, the structural aspect of the graphene can be maintained for high resolution investigation and EELS spectromicroscopy (Figure 2). These electronic properties (i.e. linear dispersion, no gap and Dirac point at the Fermi level) are precisely observed by angle-resolved photoemission on these ribbons at the (1-107) facet [2] and this will be discussed in term of curvature effect, quantum confinement or presence of sp3 bonding with respect to the STEM-EELS investigation [3].

[1] "Exceptional ballistic transport in epitaxial graphene nanoribbons," J. Baringhaus, M. Ruan, F. Edler, A. Tejeda, M. Sicot, A. Taleb-Ibrahimi, A.-P. Li, Z. Jiang, E.H. Conrad, C. Berger, C. Tegenkamp, and W.A. de Heer, Nature 506, 349 (2014).
[2] “A wide band gap metal-semiconductor-metal nanostructure made entirely from graphene”
J. Hicks, A. Tejeda, A. A. Taleb-Ibrahimi, M.S. M.S. Nevius, F. F. Wang, K. K. Shepperd, J. J. Palmer, F. Bertran, P. Le Fèvre, J. Kunc, W.A. de Heer, C. Berger, E.H. Conrad, Nature Physics 9, 49 (2013).
[3] “The origin of the gap in armchair sidewall nanoribbons: a structural study” I. Palacio, A. Celis, A. Gloter, A. Zobelli, M. Sicot, D. Malterre, M.S. Nevius, C. Berger, W.A. de Heer, E.W. Conrad, A. Taleb-Ibrahimi and A. Tejeda, in preparation.


Fig. 1: Figure 1. General overview of graphene nanoribbons grown on SiC and SiC facets (sidewall ribbons). a) Scheme of the localization of the ribbons on the samples. b) STM image showing the regions with [0001] normal. Plateaus width of 50 nm c) Cross sectional TEM image of the array of ribbons in another sample. Plateaus width of 300 nm.

Fig. 2: Figure 2. STEM-HAADF view of a graphene-SiC interface.

Type of presentation: Poster

MS-2-P-3218 TEM electron diffraction analysis of few-layer Black Phosphorus

Vicarelli L.1, Castellanos-Gomez A.1, Van der Zant H. S.1, Zandbergen H. W.1
1Kavli Institute of Nanoscience, Delft University of Technology, Delft, The Netherlands
l.vicarelli@tudelft.nl

Black phosphorus (BP) is an allotrope of Phosphorus characterized by a layered structure. It has been recently shown [1] that, similarly to graphene, it can be mechanically exfoliated to isolate atomically thin layers which have very interesting electrical and photonic properties. Single-layer BP is in fact an intrinsic semiconductor with a direct bandgap (~2 eV) and it has been employed in the fabrication of field-effect transistors with large current on-off ratios and high mobilities (100-3000 cm2/Vs) [1].

Given the rising interest in this layered material, an extensive TEM analysis of few-layer BP was performed [2].
We have investigated the Electron Diffraction (ED) pattern of few-layer black phosphorus transferred on a holey Silicon Nitride membrane with 1 µm holes diameter (see Figure 1(a)). An HRTEM image from a multilayer area of the sample is shown in Figure 1(b). The uniformity in this image indicates that the lattice contains no extended defects (single vacancies cannot be detected). We found that electron diffraction patterns depend on the number of layers and thus ED can be employed to determine the thickness of the BP flakes. We simulated electron diffraction patterns finding that the ratio between the 101 and 200 reflections depends on the number of black phosphorus: in particular this ratio is > 1 for single layer BP and decreases rapidly with the number of layers. The table shown in Figure 2 summarizes the simulated 101/200 intensity ratios for different number of layers, together with the experimental data acquired. Figure 3(a) and 3(b) show an ED taken from a thin region and a thick region of the flake, with 101/200 intensity ratios of 0.4 and 0.01, respectively.
We also noticed the presence of “forbidden” reflections (h+l = 2n+1) in the thin sample, which was not accounted in our simulations. This could be explained by the presence of adatoms on the surface of the black phosphorus layer or a slight distortion of the lattice.

References:
[1] Li, L.; Yu, Y.; Ye, G. J.; Ge, Q.; Ou, X.; Wu, H.; Feng, D.; Chen, X. H.; Zhang, Y.
Preprint at arXiv:1401.4117 (2014)

[2] “Isolation and characterization of few-layer black phosphorus”, Castellanos-Gomez, Andres; Vicarelli, Leonardo; Prada, Elsa; Island, Joshua O.; Narasimha-Acharya, K. L.; Blanter, Sofya I.; Groenendijk, Dirk J.; Buscema, Michele; Steele, Gary A.; Alvarez, J. V.; Zandbergen, Henny W.; Palacios, J. J.; van der Zant, Herre S. J. Preprint at arXiv 1403.0499 (2014)


The research leading to these results has received funding from the European Research Council, ERC Project n. 267922

Fig. 1: (a) Optical image of a black phosphorus flake transferred onto a holey silicon nitride membrane. (b) High resolution transmission electron microscopy image of the multilayered region of the flake (~ 13-21 layers).

Fig. 2: Thickness dependence of the electron diffraction patterns. We display the thickness dependence of the intensity ratio between the 101 and 200 reflections. The experimental data acquired on two spots of the thin flake and one spot of the thicker area has been included for comparison.

Fig. 3: (a) and (b) are the electron diffraction patterns acquired with a 400 nm spot on the thin (~ 2 layers) and on the thick (~ 13-21 layers) region of the flake, respectively.

Fig. 4:
Type of presentation: Poster

MS-2-P-3242 In situ growth of layered carbon

Kling J.1, Hansen T. W.1, Wagner J. B.1
1Center for Electron Nanoscopy (DTU Cen), Technical University of Denmark, Kgs. Lyngby, Denmark
jenk@cen.dtu.dk

Nanostructured carbon materials are predicted to play a major role in future electronic applications. Cheaper and smaller components with improved or new functionality and lower power consumption are necessary, where conventional materials reach their limitations. Layered carbon materials, such as graphene or multilayer graphene, can be used for extremely compact devices with outstanding performance [1],[2]. A cheap way to synthesize such materials on a large scale is chemical vapor deposition (CVD) growth on catalysts like copper or nickel [3],[4]. However, the understanding and control of such growth processes are still in their infancy.

Here we present in situ transmission electron microscopy (TEM) experiments in a FEI Titan 80-300 Environmental TEM (ETEM) for studying the growth of layered carbon materials on Ni and Cu catalysts. The ETEM allows imaging with controlled gas environments around the sample up to a few mbar. In combination with a MEMS-based heating holder, growth of layered carbon materials is systematically studied at the atomic level using various carbon sources and growth temperature.

As an example, growth of few layer graphene from C2H2 on a Ni catalyst is shown in Fig. 1-4. NiO particles in the size range up to a few hundred nm are reduced in the microscope under H2 at 500-600°C in order to form a catalytically active Ni surface. Introducing C2H2 at about 650°C leads to growth of layered carbon (Fig. 1-4). By following the appearance of carbon layers, the growth rate dependence on various parameters can be determined directly from the ETEM observations.

[1] K. S. Novoselov, S. V. Morozov, T. M. G. Mohinddin, L. a. Ponomarenko, D. C. Elias, R. Yang, I. I. Barbolina, P. Blake, T. J. Booth, D. Jiang, J. Giesbers, E. W. Hill, and a. K. Geim, Phys. Status Solidi 244, 4106 (2007).
[2] F. Schwierz, Proc. IEEE 101, 1567 (2013).
[3] X. Li, W. Cai, J. An, S. Kim, J. Nah, D. Yang, R. Piner, A. Velamakanni, I. Jung, E. Tutuc, S. K. Banerjee, L. Colombo, and R. S. Ruoff, Science 324, 1312 (2009).
[4] X. Li, W. Cai, L. Colombo, and R. S. Ruoff, Nano Lett. 9, 4268 (2009).


Financial support of the 7th Framework project “GRAFOL” is gratefully acknowledged. The A.P. Møller and Chastine Mc-Kinney Møller Foundation is acknowledged for their contribution toward the establishment of the Center for Electron Nanoscopy in the Technical University of Denmark. Thanks to Søren B. Simonsen and Quentin Jeangros for providing the NiO samples.

Fig. 1: Three layers grown shortly after introduction of C2H2.

Fig. 2: Multiple layers grown 79.2s after Fig. 1; the arrow marks next growing layer close to the metal particle surface.

Fig. 3: Multiple layers grown 80s after Fig. 1; the arrows mark next growing layers close to the metal particle surface.

Fig. 4: Multiple layers grown 80.8s after Fig. 1; the arrows mark next growing layers close to the metal particle surface.

Type of presentation: Poster

MS-2-P-3370 Novel 3-dimensional nanocomposite of covalently interconnected multiwalled carbon nanotubes using Silicon as an atomic welder

Pulickal Rajukumar L.1, Belmonte M.2, Roman B.2, Slimak J.1, Elías A. L.1, Cruz-Silva E.2, Perea-López N.1, Morelos-Gómez A.3, Terrones H.4, Miranzo P.2, Terrones M.1,3
1The Pennsylvania State University, University Park, United States, 2 Institute of Ceramics and Glass (ICV-CSIC), Madrid, Spain, 3Shinshu University, Nagano, Japan, 4Rennselaer Polytechnic Institute, Troy, United States
lzp130@psu.edu

There is a growing interest in synthesizing three-dimensional (3-D) carbon nanotube structures with multi-functional characteristics. Here, we report the fabrication of a novel composite material consisting of 3-D interconnected multi-walled carbon nanotubes (MWNTs) with Silicon Carbide (SiC) nano- and micro-particles. The materials were synthesized by a two-step process involving the chemical coating of MWNTs with Silicon oxide, followed by Spark Plasma Sintering (SPS). SPS enables the use of high temperatures and pressures that are required for the carbothermal reduction of silica and for the densification of the material into a 3-D composite block. Covalent interconnections of MWNTs are facilitated by a carbon diffusion process resulting in silicon carbide formation as silica coated MWNTs are subjected to high temperatures. The presence of SiC in the sintered composite has been confirmed through Raman spectroscopy, which shows the characteristic peak close to 800 cm-1 and also Energy Filtered Transmission Electron Microscopy maps. X-ray Diffraction, Scanning Electron Microscopy, Energy Dispersive X-Ray Spectroscopy and High Resolution Transmission Electron Microscopy have also been used to characterize the produced material. Interestingly, the thermal property measurements of the sintered composite reveal a high thermal conductivity value (16.72 W/mK) for the material. From the electrical point of view, a 3-D variable range hopping (VRH) electron hopping was observed in the composite.


Fig. 1: High Resolution Transmission Electron Microscopy images of SiC/MWNT composite prepared by Spark Plasma Sintering.

Fig. 2: (a) Raman spectrum of the SiC/MWNT sample showing characteristic D, G and G' peaks for MWNTs and the SiC peak at 800 cm-1. (b) X-Ray diffraction data of SiC/MWNT composite. (c) Raman mapping of G peak position and (d) SiC peak  position within a 30 µm x 30µm area.

Type of presentation: Poster

MS-2-P-3372 Imaging of carbon nanostructures by low energy STEM below 5 keV

Pokorná Z.1, Knápek A.1, Jašek O.2, Prášek J.3, Majzlíková P.3
1Institute of Scientific Instruments of the ASCR, v. v. i., Královopolská 147, Brno, Czech Republic, 2Masaryk University, CEITEC - Central European Institute of Technology, Kamenice 753/5, 625 00 Brno, Czech Republic, 3Brno University of Technology, CEITEC - Central European Institute of Technology, Technická 3058/10, 616 00 Brno, Czech Republic
zuzana.pokorna@isibrno.cz

Our work deals with the imaging of nanostructures composed of light biogenic elements, such as carbon nanotubes, by low energy scanning transmission electron microscopy (STEM). Compared to imaging at the voltages commonly used for TEM and STEM, low energy electrons seem very promising in terms of specimen damage that is caused by a number of elastic and inelastic collisions [1]. In carbonaceous materials, the most problematic is probably the knock-on damage, where the structure can be impaired by carbon atom displacement. To avoid this problem with structures composed of light elements, a reduction in beam voltage going down to 5 keV has recently been proposed [2]. The range below 5 keV has not been explored yet for this purpose, although electron scattering in matter is lower for these energies, which allows achieving a higher spatial resolution [3]. We aim to demonstrate that additional reduction of incident electron energy may yield interesting contrast features.

We used a FEI Magellan 400L microscope capable of high resolution imaging even at low and very low incident electron energies, equipped with a multi-segment, retractable STEM detector. Carbon specimens were prepared e.g. by depositing a solution of commercial Sigma Aldrich nanotubes, with dimethylformamide used as a solvent, on Agar S147 holey carbon mesh grids. Contrast features were recorded by secondary electron (SE), bright field (BF) and dark field (DF) detectors, including high-angle annular dark field (HAADF).

We have studied the aspects influencing the image information, such as incident electron energy, electron dose, sample thickness, the presence of the ubiquitous hydrocarbon contamination layer and other. The results were also tested using Monte Carlo simulations.

References:

[1] INADA, H., et al. Atomic imaging using secondary electrons in a scanning transmission electron microscope: experimental observations and possible mechanisms. Ultramicroscopy, 2011, 111.7: 865-876.

[2] BEYER, Y.; BEANLAND, R.; MIDGLEY, P. A. Low voltage STEM imaging of multi-walled carbon nanotubes. Micron, 2012, 43.2: 428-434.

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The financial support of the Czech Ministry of Education, Youth and Sports through projects LO1212 and CZ.1.05/1.1.00/02.0068, and of the Academy of Sciences of the Czech Republic through projects AVČR L100651304 and AVČR L100651402, is acknowledge