Friday 02 September
Time Amphithéâtre Salle Bellecour 1,2,3 Salle Prestige Gratte Ciel Salle Gratte Ciel 1&2 Salle Tête d'or 1&2 Salon Tête d'Or Salle Gratte Ciel 3
09:00
09:00-10:00
Added to your list of favorites
Deleted from your list of favorites

PL6
Plenary Lecture 6

Plenary Lecture 6

09:00 - 10:00 Plenary Lecture 6 - Visualizing Crystal Growth Processes using Liquid Cell Transmission Electron Microscopy. Frances ROSS M. (Yorktown Heights, USA)

10:30
10:30-12:45
Added to your list of favorites
Deleted from your list of favorites

MS1-II
MS1: Structural materials, defects and phase transformations
SLOT II

MS1: Structural materials, defects and phase transformations
SLOT II

Chairmen: Patricia DONNADIEU (ST MARTIN D'HERES CEDEX, FRANCE), Randi HOLMESTAD (Trondheim, NORWAY), Simon RINGER (Sydney, AUSTRALIA)
10:30 - 11:00 #8312 - MS01-S67 Recent investigations of small-scale plasticity mechanisms in 3D and small-sized systems using advanced in-situ TEM nanomechanical testing.
Recent investigations of small-scale plasticity mechanisms in 3D and small-sized systems using advanced in-situ TEM nanomechanical testing.

Recently, the development of a new generation of advanced instruments for in-situ TEM nanomechanical testing has allowed establishing a one-to-one relationship between load-displacement characteristics and stress-induced microstructure evolution in the transmission electron microscope (TEM). In the present work, it will be demonstrated that a step forward in the investigation of structural defects and small-scale plasticity mechanisms can be made by combining commercial [1] and in-house developed lab-on-chip [2] nanomechanical testing techniques with advanced TEM techniques.

High resolution aberration corrected TEM and STEM as well as automated crystallographic orientation and phase mapping in TEM (ACOM-TEM) have been used to reveal the nanoscale plasticity mechanisms controlling the mechanical responses of nanocrystalline (nc) metallic Al and Pd freestanding thin films. Special attention was paid to the strain rate sensitivity of nanostructured metallic materials involving thermally activated plasticity mechanisms [3] (Figure 1). Furthermore, the microstructure of nc Pd thin films subjected to hydriding cycles has been investigated in order to unravel the interaction mechanisms of hydrogen with dislocations and interfaces [4] as well as the influence of hydrogen cycling on the mechanical properties of the Pd films. An original method combining the measurement of dislocation mobility using in-situ TEM nanomechanical testing and dislocation dynamic (DD) simulations has also been used to investigate the plasticity of olivine small-sized crystals at low temperature. It demonstrates for the first time the possibility of characterizing the mechanical properties of specimens, which could be available in the form of micron-sized samples only [5].

More recently, quantitative nanobeam electron diffraction (NBED) was used to investigate the relationship between the local atomic order and the activation of shear transformation zones (STZs) in nanostructured ZrNi metallic glasses freestanding thin films. These films exhibit outstanding mechanical properties involving very large homogenous plastic deformation and giant ductility without the observation of mature shear bands until fracture. The basic principle of NBED is shown in Figure 2, consisting of a coherent electron beam with diameter of around 0.4 nm in order to produce two-dimensional diffraction patterns from atomic clusters with comparable size. Furthermore, high resolution HAADF-STEM and EELS revealed a heterogeneous microstructure with Ni-rich and Zr-rich regions exhibiting different atomic densities with characteristic length of 2-3 nm. The role of such behaviour in the absence of shear bands and the delay of fracture in the ZrNi thin films is discussed.

 

 

References

[1] H. Idrissi, A. Kobler, B. Amin-Ahmadi, M. Coulombier, M. Galceran, J-P Raskin, S.Godet, C. Kübel, T. Pardoen, D. Schryvers. Applied Physics Letters. 104 (2014) 101903

[2] H. Idrissi, B. Wang, M.S. Colla, J.P. Raskin, D. Schryvers, T. Pardoen. Advanced Materials. 23 (2011) 2119

[3] M.S. Colla, B. Amin-Ahmadi, H. Idrissi, L. Malet, S. Godet, J.P. Raskin, D. Schryvers, T. Pardoen. Nature communications. 6 (2015) 5922

[4] B. Amin-Ahmadi, D. Connetable, M. Fivel, D. Tanguy, R. Delmelle, S. Turner, L. Malet, S. Godet, T. Pardoen, J. Proost, D. Schryvers, H. Idrissi. Acta Materialia. 111 (2016) 253.

[5] H. Idrissi, C. Bollinger, F. Boioli, D. Schryvers, P. Cordier. Science Advances. 2 (2016) e1501671.

Hosni IDRISSI (Antwerpen, BELGIUM)
Invited
11:00 - 11:15 #4913 - MS01-OP200 In situ deformation of nanocrystalline Al2O3 thin films at room temperature.
In situ deformation of nanocrystalline Al2O3 thin films at room temperature.

Introduction

Recent TEM in situ mechanical experiments on single alumina nanoparticles have shown unexpected plasticity in room temperature alumina [1, 2]. These results push the theoretical boundaries of ceramics mechanical ductility towards comparable levels with metals. The important questions for materials science now are: (i) whether the plastic behaviour can be transferred into polycrystalline systems; (ii) what is the microstructure of such plastic polycrystalline system and (iii) what is the mechanism behind the hypothetical plasticity of the polycrystalline system. Relatively cheap and abundantly available engineering ceramic, such as alumina, with room temperature plasticity would be a breakthrough in the engineering ceramics field.

We report the findings of our study of polycrystalline alumina thin films, produced by pulsed laser deposition, with crystal size of < 5 nm using TEM and in situ TEM. Pulsed laser deposition is an extreme fabrication method where the deposition material is transformed into plasma by a short laser pulse. As the plasma quickly expands into vacuum or background gas the nucleation and growth of nanoparticles is rapidly quenched. Alumina produced this way has an exotic, nanocrystalline microstructure, and is a strong candidate for having the capability for room temperature plasticity. The more conventional TEM studies are focused on determining the as-received state of the material, grain size, morphology, crystal structure, grain boundary structure and whether any structural defects pre-exist since they have major impact on the mechanical response of the material. In situ TEM studies are focused on analysis of the material’s mechanical response (strain, dislocation activity, fracture etc.) to compression and indentation forces and look for evidence of the mechanism behind the mechanical response.

Experimental

Pulsed laser deposition (PLD) of Al2O3 thin films was done on various substrates including silicon, sapphire and sodium chloride using PLD coating equipment (Nano2Energy Laboratory, Italian Institute of Technology and Coldabtm PLD coating system, Picodeon Ltd Finland).

Two techniques were used to prepare TEM characterization samples from the PLD alumina coatings. First, TEM samples were prepared using a focused ion beam (FIB) lift-out technique and second, NaCl crystals coated with PLD thin film alumina were dissolved in water and the free-standing alumina film was deposited on a TEM grid. Figure 1 shows a TEM image of the microstructure of the PLD alumina film prepared using FIB lift-out method. Figure 2 shows a selected area electron diffraction pattern taken from the Figure 1 site indicating the presence of polycrystalline gamma-Al2O3.

For in situ TEM mechanical testing, R-plane sapphire substrates were used. Sapphire substrate was prepared using broad ion milling (Ilion II, Gatan Inc.) to produce an electron transparent, roughly 20° edge on the sapphire substrate. Furthermore a part of the edge was modified with FIB to produce electron transparent anvils with flattened tip in order to quantify the area of compression. The produced edge and anvils were either directly PLD coated or a PLD film separated from the NaCl substrate was transported on the sapphire edge or anvils.

The in situ tests were conducted using Nanofactorytm and Hysitron® PI 95 in situ TEM sample holders with JEOL 2010F and FEI Titan microscopes. In the test the PLD alumina film was compressed between the sapphire substrate and a diamond tip and the deformation process was filmed in situ together with synchronized strain and force measurement.

Acknowledgments
The authors thank the Centre LYonnais de Microscope (CLYM) for access to the electron microscopy equipment.

[1]           E. Calvié et al. Journal of European ceramic society, Vol. 32, No. 10, p. 2067-2071, 2012

[2]           E. Calvié, et al. Materials Letters, Vol. 119, p. 107-110, 2014

Erkka FRANKBERG (VILLEURBANNE CEDEX), Lucile JOLY-POTTUZ, Francisco GARCIA, Turkka SALMINEN, Thierry DOUILLARD, Bérangère LE SAINT, Ville KEKKONEN, Saumyadip CHAUDHURI, Jari LIIMATAINEN, Fabio DI FONZO, Erkki LEVÄNEN, Karine MASENELLI-VARLOT
11:15 - 11:30 #5791 - MS01-OP202 In Situ TEM Study of Fatigue Crack Growth of Cu Thin Films Using a Modified Nanoindentation System.
MS01-OP202 In Situ TEM Study of Fatigue Crack Growth of Cu Thin Films Using a Modified Nanoindentation System.

Material fatigue is often the limiting factor for many engineering cases. Repeated cyclic loading, even at stresses well below the monotonic yield stress of the material, leads to the accumulation of microstructural damage, crack initiation, crack growth, and eventual failure of the device. In terms of the number of loading cycles, the high cycle regime of >104 loading cycles is often of interest, although there are cases in which fatigue lifetimes may reach >>107 cycles. Cyclic loading experiments with bulk specimens can determine probabilistic fatigue lifetimes, however, a more fundamental understanding of the crack initiation and growth regimes is desired. These, early stage I, fatigue processes are by their very nature of limited size. Therefore, in situ nanomechanical testing in the transmission electron microscope is a good match in terms of both size and time scales.  However, to date no general capability for fatigue loading has been made available. For this set of experiments, an in situ nanomechanical system from Hysitron, Inc. was modified to use a sinusoidal loading function, capable of frequencies from 1 to 300Hz, where the resulting displacement amplitude and phase is measured using a lock-in amplifier. Thus it is now possible to reach the 106 cycle threshold within one hour. Here, observations of microstructural changes near propagating cracks in nanocrystalline Cu are presented.

Magnetron sputtered Cu films deposited on single crystal, then floated onto Push-to-Pull MEMS devices. These devices provide a micromechanical test frame converting an indentation into a tensile force. Focused ion beam milling was employed to remove excess film, leaving a tensile specimen [1]. Mechanical loading experiments were performed at frequencies from 1 to 200 Hz with mean loads of ≈100 µN, and load amplitudes of ≈50 µN. The total number of accumulated cycles typical exceeded 105 cycles. Local microstructural change was characterized by bright field TEM video during loading. Testing was periodically paused to collect additional still micrographs, as well as data collection by a precession electron diffraction-assisted automated orientation mapping technique [2]. Observable changes in contrast preceded crack nucleation. These cracks then stably propagated prior to rapid unstable failure. Localized microstructural changes occurred near the propagating crack including larger grains oriented differently from the surrounding matrix, figure 1, indicative of possible fatigue-induced grain growth [3].

References:

[1] A Kobler, et al, Ultramicroscopy 128 (2013), p.68-81.

[2] EF Rauch, et al, Zeitschrift Für Kristallographie 225 (2010), p.103-109.

[3] BL Boyce and HA Padilla, Metallurgical and Materials Transactions a-Physical Metallurgy and Materials Science 42A (2011), p.1793-1804.

[4] Work performed by K.H., B.L.B., and D.C.B. was fully supported by the Division of Materials Science and Engineering, Office of Basic Energy Sciences, U.S. Department of Energy. Work by W.M. was performed, in part, at the Center for Integrated Nanotechnologies, an Office of Science User Facility operated for the U.S. Department of Energy (DOE) Office of Science under proposal #U2014A0026. Sandia National Laboratories is a multi-program laboratory managed and operated by Sandia Corporation, a wholly owned subsidiary of Lockheed Martin Corporation, for the U.S. Department of Energy’s National Nuclear Security Administration under contract DE-AC04-94AL85000.

Daniel BUFFORD, Douglas STAUFFER (Minneapolis, USA), William MOOK , S.a. Syed ASIF, Brad BOYCE, Khalid HATTAR
11:30 - 11:45 #5898 - MS01-OP205 Investigation of plasticity/fatigue mechanisms at interfaces in Ni using ex-situ and in-situ SEM/TEM micro/nano-mechanical testing.
MS01-OP205 Investigation of plasticity/fatigue mechanisms at interfaces in Ni using ex-situ and in-situ SEM/TEM micro/nano-mechanical testing.

The present work focuses on the fundamental plasticity/fatigue mechanisms operating at interfaces in micro/nano scale Ni samples. In-situ SEM fatigue tests have been performed on FIB prepared single and bi-crystal micropillars with well-known orientations as revealed by EBSD. Careful characterizations of the nature and the distribution of deformation dislocations, the character and the local structure of the interface as well as the mechanisms controlling the interaction between these defects under cyclic loads were performed using ex-situ TEM techniques including diffraction contrast imaging, automated crystallographic orientation and nanostrain mapping in TEM (ACOM-TEM) as well as electron tomography.

The primary TEM results obtained on single crystal micropillars after fatigue tests revealed the presence of dislocation walls structure as shown in Figure 1(a) and (b). ACOM-TEM revealed local changes of crystal orientation around 1-2 degrees at the position of the dislocation walls. Furthermore, systematic contrast analysis of dislocations in these areas confirmed that only slip systems with the highest Schmid factor have been activated. The analysis of micropillars with GBs subjected to fatigue tests has shown the accumulation of dislocations at the GBs (Figure 1(c) and (d)) without slip transfer or localized plasticity. Electron tomography has been used to investigate the 3D distribution and the interaction of deformation dislocations within the dislocation walls.  

In order to directly observe the plasticity mechanisms, quantified in-situ TEM tensile tests were performed on both single and bi-crystal samples using the Pi 95 picoIndenter instrument and a MEMS device called ‘’Push-to-Pull’’ (PTP) from Hysitron.Inc (Fig 2). In order to minimize the effect of FIB on the in-situ tensile samples, an original sample preparation method combining twin jet electro-polishing and FIB was used, see figure (2). Nucleation-controlled-plasticity has been observed with defects induced by FIB at the edges of the sample acting as preferential sources for the nucleation of deformation dislocations. Furthermore, the in-situ TEM nanotensile experiments revealed the elementary mechanisms controlling the interaction between dislocations and pre-selected GBs from the electro-polished thin foils, as the nucleation of dislocations from GB, see figure (3). These results indirectly shed light on the micropillar’s behaviour as the root cause of the deformation is connected to the FIB preparation and the dislocation/GB interaction mechanisms.

Vahid SAMAEEAGHMIYONI (Antwep, BELGIUM), Jonas GROTEN, Hosni IDRISSI, Ruth SCHWAIGER, Dominique SCHRYVERS
11:45 - 12:00 #6911 - MS01-OP215 Assessing chemical and microstructural evolution at interfaces of γ’- strengthened superalloys at high temperature by in situ TEM heating experiments.
Assessing chemical and microstructural evolution at interfaces of γ’- strengthened superalloys at high temperature by in situ TEM heating experiments.

Single crystal Ni-base superalloys exhibit outstanding high temperature properties and are used as turbine blade material in advanced gas turbines, where they withstand high temperatures and mechanical loads in harsh environments [1]. Their high temperature properties are associated with their unique two phase γ/γ' microstructure consisting of cuboidal γ’ precipitates, exhibiting the ordered L12 crystal structure (vol. frac. ca. 80%) ,  which are coherently embedded in a solid solution γ fcc matrix. To ensure precipitation hardening at temperatures relevant to gas turbine applications, 700-1100 °C, the stability of the γ/γ’ phases is of fundamental importance. At high temperatures the γ' - precipitates start to dissolve until a stable γ' - volume fraction is reached [2][3].

In this analysis the microstructural evolution of the γ and γ’ phase towards a new thermodynamic equilibrium, corresponding to a selected temperature, is investigated. In situ TEM studies with chip-based heating systems (by DENS solution) are performed on a single crystal Ni-based superalloy, ERBO1. Figure 1 presents a typical ERBO1 lamella, fixed in between the Pt spirals of a high temperature heater chip. The sample is heated to 950°C for different  time periods and is subsequently quenched (see heating profile on the right of Figure 1). The elemental distribution was measured by EDXS at room temperature, whereby the sample was tilted close to the [001] zone axis to minimize projection effects. Figure 2 shows the γ/γ' microstructure of the ERBO1 alloy with a black rectangle indicating the region where the ChemiSTEM EDXS measurements were conducted after each heating sequence. The Cr concentration profiles show that the transition towards a new thermodynamic state can indeed be resolved, with progressing time the Cr concentration decreases in the γ phase and increases in the γ’ phase. Such data contain valuable information on the kinetics of interdiffusion in a real superalloy microstructure. Quantitative evaluation of diffusion profiles will enable to simultaneously determine interdiffusion coefficients of various alloying elements. Finally the experimental results will be validated by complementary ex situ experiments on bulk samples employing quenching in a vertical oven. Furthermore, the results will be compared with thermodynamic and kinetic data from theoretical calculations.

[1] R.C.Reed, Superalloys Cambridge University Press, (2006)

[2] A.Royer et al., Acta Materialia, (1998) Vol46

[3] Y.M.Eggeler et al., Proc. International Microcopy Congress IMC18 2016, Prag, Czech 2014

Acknowledgement:  This work has been carried out within the framework of the SFB-TR 103 "Single Crystal Superalloys".

Yolita EGGELER (Erlangen, GERMANY), Daniel ENGE, Erdmann SPIECKER
12:00 - 12:15 #6242 - MS01-OP209 Electron beam induced in situ writing and recovering of vacancy layers in Ge2Sb2Te5 crystal lattice.
Electron beam induced in situ writing and recovering of vacancy layers in Ge2Sb2Te5 crystal lattice.

Phase change materials (PCMs), such as Ge-Sb-Te-based alloys, are of high interest due to their technologically eminent optical and electronic properties. Among the Ge-Sb-Te based PCMs, Ge2Sb2Te5 (GST225) is a well-known compound and the most used PCM. GST225 is utilized in optical and electronic data storage devices [1]. Technological relevant phases of GST225 are an amorphous phase, a metastable (cubic) phase containing 20% of intrinsic vacancies and a stable trigonal phase. Despite comprehensive studies on amorphisation-crystallisation processes of GST225, the phase transformation from the cubic to the trigonal phase has not been experimentally investigated up to now. In this work, the phase transformation is investigated in-situ in a Cs-corrected STEM by exposing distinct layers of GST225 crystal lattice to repeated line scanning of a focused electron beam [2].

The cubic GST225 phase was prepared by laser irradiation of amorphous GST225 thin films [2]. No planar defects such as vacancy layers (VLs) [3] were observed in the cubic GST225 phase (Fig. 1(a)) [2]. However, VLs can be intentionally produced in the GST225 phase by repeated line scanning of the focused electron beam along individual mixed GeSb/V layers, while no such defects can be created by scanning the electron beam along the Te layers (Fig.1(b)). Most notably, these VLs disappeared after repeated scanning of the focused electron beam over a scanning window covering these defects (Fig. 1(d)). Moreover, the Te-Te distance in the [001] direction is reduced in these newly formed VLs [2]. The observed ordering of vacancies into layers is due to energy transfer by inelastic interactions of the electron beam with the GeSb layer and takes place by diffusion of Ge and Sb atoms towards vacancies located in the nearest neighbouring GeSb/V layers [2]. Previous DFT calculations indicated that the ordering of vacancies into layers is energetically favourable [4]. It was also shown that the structural transition to the layered trigonal GST225 phase is driven by the ordering of vacancies in the cubic GST225 phase. However, the vacancy ordering necessarily occurs before complete formation of the vacancy planes, the process of which could be observed in this work.

From the above presented results and those reported in Refs. [3,4], a transformation mechanism between the cubic and trigonal GST225 phases can be proposed. The stable GST225 consists of nine layers alternatingly containing Te and GeSb in one unit cell, e.g. -Te-GeSb-Te-GeSb-Te-VL-Te-GeSb-Te-GeSb-, with intrinsic VL between adjacent Te layers (Fig. 2(a)). The stacking sequence in cubic GST225 is -Te-GeSb/V-Te-GeSb/V-Te-GeSb/V-Te-GeSb/V-Te-GeSb/V- (Fig. 2(b)) and the trigonal GST225 can therefore be derived from the cubic GST225 by removing the vacancies from the sublattice and accumulating them in the VLs. However, movement of Ge and Sb atoms in the cubic GST225 lattice along direction and subsequent shift of newly formed building blocks against each other are required to complete the trigonal lattice. Thus, this phase transition is driven by local short-distance movements of Ge and Sb atoms towards vacancies without long range diffusion and without change in composition of the parent cubic GST225 phase. Consequently, the phase change between the cubic and trigonal GST225 phases is a diffusionless transformation process similar to martensitic transformation [5].

[1] S.J. Feinleib et al., Appl. Phys. Lett. 18, 254 (1971).

[2] A. Lotnyk et al., Acta Mater. 105, 1 (2016).

[3] U. Ross, A. Lotnyk et al., Appl. Phys. Lett. 104, 121904 (2014).

[4] W. Zhang et al., Nature Mater. 11, 952 (2012).

[5] W. Zhang et al.,  Adv. Eng. Mater. 10, 67 (2008).

Andriy LOTNYK (Leipzig, GERMANY), Sabine BERNÜTZ, Xinxing SUN, Ulrich ROSS, Martin EHRHARDT, Bernd RAUSCHENBACH
12:15 - 12:30 #5813 - MS01-OP203 Surface Atomic Structure and Growth Mechanism of {1 0 0}-Faceted Perovskite Oxide Nanocubes.
MS01-OP203 Surface Atomic Structure and Growth Mechanism of {1 0 0}-Faceted Perovskite Oxide Nanocubes.

Monodisperse faceted nanocrystals, with controllable shapes and sizes, have been becoming increasingly important for applications in catalysis, gas sensing, and energy conversion. Such highly shape sensitive and selective physical and chemical properties inherently stem from the atomic and electronic structures on the faceted surfaces. For elemental nanocrystals, the atomic structure on the surfaces is determined by the geometric shape itself. However, for compound materials such as alloys and complex oxides, the compositional segregation and different terminating lattice planes on the surfaces have to be taken into account. In order to understand the unique property and growth mechanism of these nanocrystals, atomic details on the faceted surfaces need to be studied on the atomic level.

 

Strontium titanate (SrTiO3), strontium zirconate (SrZrO3) and their solid solutions (SrTi1−xZrxO3) are important members in the class of perovskite structures with a general formula ABO3 (Figure 1a). These materials are of great technological and fundamental importance not only because of their interesting properties, but also because of their ability to combine and to adjust these properties by chemical substitution with a wide variety of cations. However, despite the success of the synthesis of the {1 0 0}-faceted BaTiO3, SrTiO3, and Ba1−xSrxTiO3 nanocubes, whether the {1 0 0} facets of the nanocubes are terminated with AO (SrO) or BO2 (TiO2) is a question which still remains open for speculation and investigation. A comprehensive understanding of the growth mechanisms of these faceted nanocubes has not been achieved. Direct experimental evidence for the atomic structure on these nanocube surfaces has become one of the key steps in exploring the growth mechanisms.

 

In this work, we report on detailed studies of monodisperse {1 0 0}-faceted nanocubes of SrTi1−xZrxO3 (x = 0.25 to 0.5) which were synthesized using the oil-water two-phase solvothermal method. The surface atomic structure of the monodisperse faceted nanocrystals is determined by means of aberration-corrected high-angle annular dark-field scanning transmission electron microscopy (HAADF-STEM). On the basis of the structural features on the faceted surfaces, a deeper insight into the growth mechanisms could be obtained.

 

References and Acknowledgements

 (1) Du, H.; Jia, C.-L.; Mayer, J. Surface Atomic Structure and Growth Mechanism of Monodisperse {1 0 0}-Faceted Strontium Titanate Zirconate Nanocubes. Chem. Mater. 2016, 28 (2), 650–656.

(2)  Du, H.; Wohlrab, S.; Weiß, M.; Kaskel, S. Preparation of BaTiO3 Nanocrystals Using a Two-Phase Solvothermal Method. J. Mater. Chem. 2007, 17, 4605–4610.

(3) This work has been supported in parts by the Deutsche Forschungsgemeinschaft (SFB 917).

Hongchu DU (Aachen, GERMANY), Chun-Lin JIA, Joachim MAYER
12:30 - 12:45 #5892 - MS01-OP204 Investigation of solid state dewetting phenomena of epitaxial Al thin films on sapphire using electron microscopy.
MS01-OP204 Investigation of solid state dewetting phenomena of epitaxial Al thin films on sapphire using electron microscopy.

Solid state dewetting[1] is a topic of current research. Besides targeted patterning, research focuses on the mechanisms and its prevention to avoid degradation or failure of e.g. microelectronic devices. While several studies have addressed solid state dewetting of bare metallic films the focus of this study is laid on Al thin films covered with a native surface oxide layer. In order to simplify the complexity of the film microstructure we grew thin Al films by molecular beam epitaxy on (0001) single crystalline sapphire (α-Al2O3) substrates.

The microstructure and epitaxial orientation relationships of the Al films were analysed by scanning electron microscopy (SEM) and transmission electron microscopy (TEM) methods including electron backscatter diffraction (EBSD), selected area electron diffraction, high resolution TEM (HRTEM) and atomic resolved scanning TEM (STEM). The as-deposited Al films form two orientation relationships (ORI and ORII) both containing two twin-related growth variants: {111} Al || (0001) α-Al2O3 with ±<1‾10> Al || <101‾0> α-Al2O3 (OR I) and {111} Al || (0001) α-Al2O3 with ±<21‾1‾> Al || <101‾0> α-Al2O3 (OR II). The “±” indicates the twin related variants which differ by a 180° rotation around the surface normal. Cs-corrected high angle annular dark field (HAADF) HRSTEM micrographs (Fig. 1) of cross-sectional as-deposited samples indicated strain at the twin boundaries of OR I. In addition, a translation in at the twin boundary by 0.91±0.13 Å (HRTEM[5]: 0.84±0.17 Å) compared to an ideal, non-relaxed sigma three twin boundary was revealed.

After annealing for 1 to 45 hours below the melting point of aluminum (660°C) at 600°C, instead of Al islands[3] dark appearing features are observed in plan-view SEM micrographs (Fig. 3). This has been observed in literature for different model systems (Ni films on Al2O3[2] and Al films on Al2O3[4]). Two different models[1][2][4] exist to explain their formation. The capillary energy driven retraction of thin films can be described by classical solid state dewetting and would lead to holes of bare substrate surrounded by a rim slightly higher than the original film thickness.[1][2][3] In contrast, in the second model only volume and grain boundary diffusion can take place due to the formation of a thin oxide layer on top of the Al film. Film retraction below the oxide layer would result in drum-like voids.[4] Site-specific cross-sections prepared by focused ion beam and investigation by SEM, TEM and EDS revealed the presence of voids in the Al film with a thin cover layer (Fig. 2). Electron energy loss spectroscopy (EELS) of the surface layer revealed a phase transformation from amorphous alumina (as-deposited state) to γ-Al2O3 (after annealing) as proposed in literature from glancing incidence X-ray diffraction measurements.[4] Although solid state dewetting was done at 600°C, facetted single crystalline sapphire ridges form at the Al/sapphire/void triple phase boundary as a consequence of the capillary energy force component acting perpendicular to the film/substrate interface. The thickness of the Al film increases locally in the region of the sapphire ridge compared the original film thickness (Fig. 2). The drum-like features possess distinct facets and reflect the hexagonal symmetry of the basal plane of the sapphire substrate (Fig. 3). The EBSD investigations indicate that the grain boundaries act as initial points of void formation.

References

[1] C. V. Thompson, Annu. Rev. Mater. Res., 42, 399-434 (2012).

[2] E. Rabkin et al., Acta Mater., 74, 30-38 (2014).

[3] W. Kaplan et al., J. Mater. Sci., 48, 5681-5717 (2013).

[4] S. Dutta et al., J. Am. Ceram. Soc., 95, 823-830 (2012).

[5] G. Dehm et al., Acta Mater., 50, 5021-5032 (2002).

Stefan HIEKE (Düsseldorf, GERMANY), Gerhard DEHM, Christina SCHEU

10:30-12:45
Added to your list of favorites
Deleted from your list of favorites

IM8-III
IM8: Spectromicroscopies and analytical microscopy
SLOT III

IM8: Spectromicroscopies and analytical microscopy
SLOT III

Chairmen: Gerald KOTHLEITNER (Graz, AUSTRIA), Anders MEIBOM (Lausanne, SWITZERLAND), Bénédicte WAROT-FONROSE (Toulouse, FRANCE)
10:30 - 11:00 #8360 - IM08-S55 Quantitative Nanoplasmonics in the TEM.
Quantitative Nanoplasmonics in the TEM.

Quantitative Nanoplasmonics in the TEM

 

Michel Bosman1

 

1 Institute of Materials Research and Engineering, A*STAR (Agency for Science, Technology and Research), 2 Fusionopolis Way, #08-03 Innovis, Singapore 138634. michel.bosman@gmail.com

 

Scanning TEM (STEM)-based surface plasmon characterization will be discussed in this paper, and it will be argued that some plasmon properties can be measured quantitatively in the STEM, by carefully analyzing the local spectral response. Examples will be given that demonstrate the unique capability of STEM-based plasmon analysis in comparison with other experimental techniques.

Some very impressive experimental results have recently been published in which plasmons on subwavelength metal nanostructures were mapped with 1-100 fs time resolution. These include time-domain techniques based on photoelectron emission microscopy (PEEM) [1, 2], time-resolved scanning near-field optical microscopy (SNOM) [3], ultrafast TEM [4] and even plasmon mapping with free electron lasers [5].

STEM-based electron energy-loss spectroscopy (EELS) is performed in the energy/frequency domain, and it has a long history as an experimental technique for surface plasmon characterization [6-17]. However, since surface plasmon resonances are damped harmonic oscillators, it is possible to interpret the EELS plasmon spectra as Fourier-transforms of oscillations in the time-domain [18].

This paper will explore the implications of this approach, and will apply it to quantify local dynamic materials properties [19].

 

Acknowledgements:

The National Research Foundation (NRF) is kindly acknowledged for supporting this research under the CRP program (award No. NRF-CRP 8-2011-07). This work results from close collaborations with the groups of Joel Yang (SUTD, Singapore), Christian Nijhuis (NUS, Singapore), Erik Dujardin (CE-MES, France). Antonio Fernández-Domínguez (UAM, Spain), Wu Lin & Bai Ping (IHPC, Singapore).

 

[1] MI Stockman et al. Nature Photon. 1, 539-544 (2007).

[2] E Mårsell et al. Nano Lett. 15, 6601-6608 (2015).

[3] Y. Nishiyama et al. J. Phys. Chem. C 119, 16215-16222 (2015).

[4] A Yurtsever et al. Science 335, 59–64 (2012).

[5] SE Irvine et al., Phys. Rev. Lett. 93 (18) 184801 (2004).

[6] C Powell and JB Swan, Phys. Rev. 115, 869–875 (1959).

[7] PE Batson, Phys. Rev. Lett. 49, 936–940 (1982).

[8] ZL Wang, JM Cowley, Ultramicrosopy 21, 347–366 (1987).

[9] F Ouyang, PE Batson and M Isaacson, Phys. Rev. B 46, 15421–15425 (1992).

[10] J Nelayah et al., Nature Phys. 3, 348 – 353 (2007).

[11] M Bosman et al., Nanotechnology 18, 165505 (2007).

[12] B Schaffer et al., Phys. Rev. B 041401 (2009).

[13] B. Ögüt et al. ACS Nano 5 (8) 6701-6706 (2011).

[13] H Duan et al. Nano Lett. 12, 1683–1689 (2012).

[14] M Bosman et al. ACS Nano 6 (1) 319-326 (2012).

[15] D Rossouw and GA Botton, Phys. Rev. Lett. 110, 066801 (2013).

[16] S Raza et al., Optics Express 21 (22) 27344 (2013).

[17] A Teulle et al. Nature Materials 14, 87-94 (2015).

[18] M Bosman et al. Scientific Reports 3 1312 (2013).

[19] M Bosman et al. Scientific Reports 4 5537 (2014).

Michel BOSMAN (, SINGAPORE)
Invited
11:00 - 11:15 #5798 - IM08-OP143 Localized surface plasmon resonance mapping on aluminium voids with three-dimensional nanostructures.
Localized surface plasmon resonance mapping on aluminium voids with three-dimensional nanostructures.

     Today’s nanotechnology has enabled the fabrication of metallic nanoparticles with a variety of geometries, greatly advancing the research field of plasmonics. The complementary system of the inverted nanostructures such as nano-voids, however, has so far been limited to either 2D holes or spherical voids, owing to the difficulty in creating voids with well-defined 3D geometries. Here we present the first localized surface plasmon resonance (LSPR) study, both experimentally and theoretically, on aluminium nano-voids in the shape of truncated octahedra.

    Nano-voids were made in high-purity aluminium using an annealing and quenching treatment.1 To characterize LSPRs of voids, we used electron energy-loss spectroscopy (EELS) in a scanning transmission electron microscope (STEM) – an aberration-corrected FEI Titan operating at 80 kV. The lower accelerating voltage mitigates irradiation damage on aluminium, which is critical to perform reliable STEM-EELS mapping on voids. To confirm experimental results, electrodynamic EELS calculations were carried out based on electron-driven discrete dipole approximation (e-DDA).2

    Our results show that these aluminium nano-voids exhibit strongly localized field enhancements, with the LSPR energies 10.7 - 13.3 eV (116 - 93 nm), well beyond the conventional LSPR spectrum range. The LSPR tunability was demonstrated by tailoring the shape of nano-voids using controlled electron irradiation. Furthermore, owing to the simplicity of the nano-void system which is free of aluminium oxidation and supporting substrates, we demonstrate that the intrinsic LSPR properties of pure Al nanoparticles can be revealed from nano-voids characterization using the sum rule for the complementary systems. Combined with the low cost and mass producibility of Al, our results indicate that both the Al nano-voids and nanoparticles can effectively expand the available plasmonic spectrum to the extreme UV region (≤ 124 nm), which opens possibilites for applications such as plasmon-enhanced UV photoemission spectroscopy and photoionization.

 

Acknowledgements

This work was supported by the Australian Research Council (ARC) grant DP110104734 and DP150104483 and a Monash University IDR grant. The FEI Titan3 80-300 S/TEM at Monash Centre for Electron Microscopy was funded by the ARC Grant LE0454166.

 

References:

[1] Z. Zhang, T. Liu, A. E. Smith, N. V. Medhekar, P. N. H. Nakashima and L. Bourgeois, submitted. (2016).

[2] N. W. Bigelow, A. Vaschillo, V. Iberi, J. P. Camden and D. J. Masiello, ACS Nano. 6, 7497-7504 (2012).

Ye ZHU, Philip NAKASHIMA, Alison FUNSTON, Laure BOURGEOIS, Joanne ETHERIDGE (Melbourne, AUSTRALIA)
11:15 - 11:30 #5907 - IM08-OP145 Surface plasmon coupling revisited with electron energy loss spectroscopy.
Surface plasmon coupling revisited with electron energy loss spectroscopy.

In the last fifteen years or so, a significant amount of research activities took place in the field of metal nanoparticle plasmonics probed by fast electron beam. Local electron probe techniques like electron energy-loss spectroscopy (EELS) and cathodoluminescence (CL) have advanced very fast during this time. Both these techniques helped us to gain considerable insight into the plasmonic properties of metallic nanostructures [1,2].  One of the many interesting nanostructures are the metal nanoparticle dimers. They can confine a huge amount of field in the gap. As the modes can be tuned precisely by changing the separation between the nanoparticles, they bear the promise to be used as sensors. It is well known that as the two individual particles approach each other, bonding and antibonding modes appear. Bonding and antibonding corresponds to in-phase and out-of-phase interaction of the individual particle plasmons. Analogous to the molecular orbital theory, the bonding mode appears at lower energy and the antibonding mode at higher energy with respect to the individual particle plasmon mode [3]. With electron beam we can probe the precise location of the particle dimer to see which mode is excited at what location. Even though there is a wealth of literature on the plasmonic modes of individual metallic nanostructure, a systematic experimental study of coupled plasmons is deficient.

The central idea of the current work is to develop a deeper understanding on the coupling of the nanoparticle plasmons when they are brought close to each other (a few nanometers). For this purpose, we have chosen a cross shaped metal nanoparticle (Figure 1). We choose this structure for easy tuning of plasmon modes along the length of the rods by changing the rod length and because it is easy to model numerically. To have a detailed idea on the underlying physics, we start with the basic building block of the cross, i.e. a single nano rod. Then we increase the complexity of the structure by adding two rods perpendicularly to make a cross, and finally bringing two crosses close together to make a dimer. We make EELS on all of these structures and see the evolution of the modes. In this way, we will be able to explore the exact formation of different modes and the coupling between them.

To realize our idea, we have performed electron beam lithography to make silver nanostructures on Silicon Nitride substrate. EELS experiments are performed using a scanning transmission electron microscope (STEM) fitted with homemade EELS detection system [4]. To gain insight about the plasmonic modes, we perform 3D boundary element method (BEM) simulations [5] and compare with the experimental data. A representative EELS spectrum and the 2D plasmon maps has been shown in Figure 2.

In the conference, we will present and discuss our experimental and simulation results. This will provide new insight into the physics of plasmon coupling.

References:

[1] M. Kociak and O. Stéphan, Chem. Soc. Rev. 2014, 43, 3865.
[2] F. J.  García de Abajo, Rev. Mod. Phys. 2010, 82, 209.
[3] E. Prodan et al. Science, 2003, 302, 419.
[4] J. Nelayah, et al. Nat. Phys. 2007, 3, 348.
[5] U. Hohenester and A. Trugler, Comput. Phys. Commun. 2012, 183, 370.

Pabitra DAS (ORSAY), Hugo LOURENÇO MARTINS, Luiz TIZEI, Mathieu KOCIAK
11:30 - 11:45 #6352 - IM08-OP153 Tailoring the wave function of electron probes for the selective detection of plasmonic modes.
Tailoring the wave function of electron probes for the selective detection of plasmonic modes.

Electronic spectroscopies are important in the study of localised surface plasmon resonances of metallic nanostructures, allowing to detect and image the strong spatial variations in the electrical field of the induced resonances of a single nanoparticle [1].

These techniques do however present some drawback when compared to their optical counterparts. While optical spectroscopies can make use of polarisation to directionally probe the response of a nanoparticle, an electronic beam can’t discriminate between energy-degenerate eigenmodes and is also blind to optical activity and dichroism.

Here we present a way to expand the applicability of EELS to the characterisation of plasmonic resonances by exploiting the recently developed methods of electron beam shaping through phase manipulation [3].

This radically new approach is based on the idea of tailoring the electronic probe to fit the properties under investigation that can then be selectively detected, allowing to perform new measurements that were previously impossible [3].

In particular, we first show how the phase in the electron beam's complex wave function couples to the electric potential of the plasmonic excitation (see fig1), allowing to selectively detect localised plasmonic excitations that possess the same symmetry as the electron probe [4].

While this concept is entirely general and potentially applicable to any plasmonic resonance, we decide to focus a first experimental demonstration on detecting the dipolar mode of a nanorod.

The ideal probe for this purpose is formed by two intensity lobes opposite in phase as shown in figure 1a, which we successfully generate in the TEM by applying state of the art phase manipulation techniques.

Finally, we show experimental proof of the method’s effectiveness on purposefully made test sample, demonstrating the viability of this new approach and opening to a new generation of plasmon oriented TEM experiments, that will expand in parallel with the availability of wave manipulation methods.

Acknowledgments:

GG, AB and JV acknowledge funding from the European Research Council under the 7th Framework Program (FP7), ERC Starting Grant No. 278510-VORTEX.

References:

[1] J. Nelayah et al., Nat. Phys. 3, 348 (2007).

[2] F.-P. Schmidt et al., Nano Lett. 12, 5780 (2012).

[3] G. Guzzinati et al., Ultramicroscopy 151, 85 (2015).

[4] H. Lorenço-Martins, M. Kociak, in preparation.

Giulio GUZZINATI (Antwerpen, BELGIUM), Hugo LOURENÇO-MARTINS, Armand BÉCHÉ, Mathieu KOCIAK, Jerôme MARTIN, Jo VERBEECK
11:45 - 12:00 #6805 - IM08-OP160 Monochromated Low-Voltage EELS of Optical Resonances in Quantum Materials.
Monochromated Low-Voltage EELS of Optical Resonances in Quantum Materials.

We have imaged and produced EELS spectra of Vacancy (NV) centers in diamond with our Zeiss Libra TEM [1] with a monochromated electron source and in-column energy filter, Fig. 1. With monochromated electron energy loss spectroscopy (EELS) we measure the amount of energy loss that an electron undergoes; this includes optical resonances and inter and intra band transitions but also of course the Cherenkov radiation background.  The low acceleration voltage of 40 kV directly reduces the background noise of the Cherenkov radiation.  According to the photoluminescence spectra, we find a resonance at 1.9 eV (corresponding to 638 nm wavelength). We have also demonstrated an unexpectedly strong surface-plasmonic absorption at the interface of silver and high-index dielectrics based on electron and photon spectroscopy [2]. The measured bandwidth and intensity of absorption deviate significantly from the classical theory. Our density-functional calculation well predicts the occurrence of this phenomenon. It reveals that due to the low metal-to-dielectric work function at such interfaces, conduction electrons can display a drastic quantum spillover, causing the interfacial electron-hole pair production to dominate the decay of surface plasmons.  This finding can be of fundamental importance in understanding and designing quantum nanoplasmonic devices that utilize noble metals and high-index dielectrics.

Depending on the composition, Quantum Materials may act as conductors, insulators, semiconductors or even as superconductors. Combinations of different quantum materials are of high interest to explore new phenomena and act as the foundation for future electronic devices at the nanometer scale. Our quantum materials research is widely spread, reaching from defect formation in graphene to the characterization of hybrid quantum materials. We present our work utilizing Low-Voltage Monochromated EELS and Low-Voltage High-Resolution Electron Microscopy (LV HREM). Together, these often improve the contrast to damage ratio obtained on a large class of samples, such as Quantum Materials. 

Fe3Sn2 is a rare metallic Kagome ferromagnet, which synthesis as a single crystal has not previously been reported (Fig. 2). We study this single crystal as well as other topological insulators with the particular interest in the correlated behavior in topologically non-trivial materials.  The (S)TEM images mapped with  low voltage EELS show the atomic structure of the layered material(Fig. 2), and the magnetic force microscopy measurements reveal the magnetic anisotropy of the crystal on the surface.  Monochromated EELS is the key to all this work, due to increased signal to noise and background and/or ZLP reduction.

References

  1. D.C. Bell, C.J. Russo and D. Kolmykov, “40 keV Atomic Resolution TEM”, Ultramicroscopy. 114, pp 31-37 (2012)
  2. D. Jin, Q. Hu, D. Neuhauser, F. von Cube, Y. Yang, R. Sachan, T.S. Luk, D.C. Bell, and N.X. Fang. “Quantum-Spillover-Enhanced Surface-Plasmonic Absorption at the Interface of Silver and High-Index Dielectrics”. Phys. Rev. Lett., 115(19), p.193901 (2015)
  3. This work was supported by the STC, Center for Integrated Quantum Materials, NSF Grant No. DMR-1231319

David BELL (Cambridge, USA), Felix VONCUBE, Peter REZ, Toshihiro AOKI
12:00 - 12:15 #6958 - IM08-OP162 Electron energy loss spectroscopy (EELS) fingerprints of p- and n-type doping in graphene.
Electron energy loss spectroscopy (EELS) fingerprints of p- and n-type doping in graphene.

Doping is a great way to make nanomaterials suitable for use in nanoscale devices and electronics. Dopant atoms supply mobile charge carriers in the form of electrons (n-type) or holes (p-type) whilst leaving the material neutral, and the physical interface between p- and n-type materials - the pn junction diode - is a fundamental building block of electronic circuits. There is a general consensus that moving beyond the (mainly) Si-, Ge- and GaAs-based technology of the last seventy years is highly desirable, so there is currently lots of interest in finding superior replacements for next-generation electronics. Graphene's discovery [1] lies at the root of this renewed interest, and it now seems likely that graphene, or nanomaterials inspired by graphene, will find their way into consumer and industrial electronics in some significant form quite soon.

 

To make progress, we need information about the electronic structure of the material of interest, and how doping affects that electronic structure. Atomic-resolution electron energy loss spectroscopy (EELS) is a great way to achieve this because it can reveal the bonding around an individual dopant atom. This insight can then be directly compared with theoretical electronic structure calculations in the form of density functional theory (DFT) to yield a detailed understanding of the electronic structure and the potential implications for use in nanoscale devices.

 

In this talk I shall present atomic-resolution K-edge EEL spectra for the case of substitutional B and N dopants in graphene synthesised using low-energy ion implantation, and I shall explain how a careful comparison with theoretical DFT calculations indicates that the EELS data is in fact the first direct experimental evidence of p- and n-type doping in graphene. [2] This approach demonstrates how potentially very lucrative information can be extracted from joint studies of experimental and theoretical EELS, and it could be readily extended to other nanomaterials.

[1] 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, 666 (2010)

[2] D. Kepaptsoglou, T. P. Hardcastle, C. R. Seabourne, U. Bangert, R. Zan, J. A. Amani, H. Hofsäss, R. J Nicholls, R. M. D. Brydson, A. J. Scott, Q. M. Ramasse ACS Nano 9, 11398 (2015)

Demie KEPAPTSOGLOU, Trevor HARDCASTLE (Leeds, UK), Che SEABOURNE, Ursel BANGERT, Recep ZAN, Julian AMANI, Hans HOFSÄSS, Rebecca NICHOLLS, Rik BRYDSON, Andrew SCOTT, Quentin RAMASSE
12:15 - 12:30 #6113 - IM08-OP151 Near Band Edge excitation in 2D materials by Transmission Electron Microscopy.
Near Band Edge excitation in 2D materials by Transmission Electron Microscopy.

In this work, we report on the characterization of near band edge excitation by electron energy loss spectroscopy (EELS). This technique is operated in a Transmission Electron Microscope and allows to rely the structure of a material obtained by HR-TEM with its chemical and physical properties deduced from EELS. Indeed, when the energy transfered by a transmitted electron remains below 50 eV, it is possible to have access to the electronic structure of the material and more precisely to its dielectric function. [1] In other words, we are able to obtain informations such as, plasmon resonances, interband transitions and band gap measurements.

 

We used a Libra 200 equipped with an electrostatic monochromator operating at 80 kV. Thanks to the in-column filter, energy loss signal is recorded on a CCD camera with a spectral resolution of 150 meV. The sharp cut-off of the omega filter allows to probe the dielectric properties of semiconducting materials down to 1eV losses. We are able to determine bandgaps in several 2D materials and rely them to the number of layers. For instance, we can see the blue shift of the “optical absorption” from several MoS2 layers (1.4 eV) down to a single layer (1.8 eV).

 

Recently, thanks to dedicated operating modes [2,3], we have been able to obtain additional informations on the plasmons and interband transitions over the Brillouin Zone in hexagonal Boron Nitride (hBN). Energy Filtered scattering patterns have been recorded in the TEM to have access to the symmetries of the dipole matrix elements involved in the observed transitions. Moreover, by dispersing the energy along specific crystallographic directions, we accessed to the related dispersion of plasmons and interband transitions with the so-called ω-q maps [2] as representated on fig 1. We show that, due to a strong electron-hole interaction, the observed dispersion is related to the one of the exciton [4]. The experimental results are in good agreements with inelastic X-ray scattering experiments [5] and calculations [6] as shown on fig 3.



[1] R.E. Egerton, Electron Energy-Loss Spectroscopy in the Electron Microscope, 3rd edition, Springer (2011)

[2] P. Wachsmuth et al., Phys. Rev. B 88, 075433 (2013)

[3] G. Radtke, G. Botton, J. Verbeeck, Ultramicroscopy 106, 1082 (2006)

[4] F. Fossard, L. Schué, F. Ducastelle, J. Barjon, A. Loiseau, in preparation

[5] S. Galambosi, L. Wirtz et al., Phys. Rev. B 83, 081413(R) (2011)

[6] G. Fugallo, M. Aramini, J. Koskelo, K. Watanabe, T. Taniguchi, M. Hakala, S. Huotari, M. Gatti, and F. Sottile, Phys. Rev. B 92, 165122 (2015)

Frédéric FOSSARD (Chatillon), Léonard SCHUÉ, Etienne GAUFRÈS, Amandine ANDRIEUX, François DUCASTELLE, Annick LOISEAU
12:30 - 12:45 #6064 - IM08-OP150 Vibrational Spectroscopy in the Electron Microscope.
Vibrational Spectroscopy in the Electron Microscope.

Vibrational spectroscopy in the scanning transmission electron microscope (STEM) was introduced two years ago [1, 2], and it has made much progress since.  It has opened a new window on the world of materials, in which nothing is quite like it was before.  

 

The main vibrational modes occur at energies of 0-500 meV, and exploring them requires a monochromated STEM-EELS system with an energy resolution

 

The energy of vibrational modes is given by ∆E = ħ √(k/m), where k is the force constant of the atomic bond and m the effective mass of the vibrating nucleus.  Strongly bonded light atoms give the highest vibrational energies, starting with hydrogen, an element that is nearly invisible in traditional electron microscopy.  Fig. 1(a) shows a vibrational spectrum of Ca(OH)2 [3], in which the peak at 452 meV is due to O-H stretch, and Fig. 1(b) shows the particle from which the spectrum was recorded.  Fig. 1(c) shows how the strength of the vibrational peak varied with the distance from the particle: the signal decayed only gradually outside the particle, and was still 50% strong 35 nm away.

 

Fig. 2 shows an EEL spectrum of guanine compared to an IR spectrum from the same specimen [4].  The agreement between the two types of spectra is very good. EELS has worse energy resolution (~10 meV), but much better spatial resolution than regular IR.  As is typical of vibrational spectroscopies, the different peaks can be assigned to different types of bonds and vibration modes (see the inset in Fig. 1).  

 

In order to minimize radiation damage, both the OH and guanine spectra were acquired in an “aloof” mode, with the electron beam parked just outside the sample [1, 3-5].  Aloof spectroscopy makes it possible to select the maximum energy of the beam-sample interaction, simply by adjusting the beam-sample distance [4,5].  Its great import to vibrational EELS is that the vibrational signal can be excited even when the interaction energy is limited so that ionization damage of the sample cannot occur.  It may even be possible to spatially map the vibrational features of a beam-sensitive sample by “coarse step (leapfrog) scanning”: scanning with a discrete pixel increment of 10-100 nm, so that even though the area that the beam traverses in each new position is essentially destroyed, large parts of the sample are not touched by the beam and remain in a pristine state [6]. 

 

We gratefully acknowledge the use of facilities within the LeRoy Eyring Center for Solid State Science at ASU and IAMDN at Rutgers, and grants NSF MRI-R2 #959905, DE-SC0004954, and DE-SC0005132.

 

References

[1] O.L. Krivanek et al., Nature 514 (2014) 209.

[2] T. Miata et al., Microscopy 63 (2014) 377.

[3] P.A. Crozier et al., Microsc. and Microan. 21 (Suppl 3, 2015) 1473.

[4] P. Rez et al., Nature Comm. 7 (2016) DOI: 10.1038/ncomms10945.

[5] R.F. Egerton, Ultramicroscopy 159 (2015) 95.

[6] R.F. Egerton et al., submitted to Microsc. and Microan. 22 (Suppl 3, 2016).

Ondrej KRIVANEK (Kirkland, USA), Toshihiro AOKI, Philip BATSON, Peter CROZIER, Niklas DELLBY, Raymond EGERTON, Tracy LOVEJOY, Peter REZ

10:30-12:45
Added to your list of favorites
Deleted from your list of favorites

IM3-III
IM3: Near-field, photon beam and unconventional microscopies
SLOT III

IM3: Near-field, photon beam and unconventional microscopies
SLOT III

Chairmen: Emmanuel BEAUREPAIRE (Palaiseau, FRANCE), Christian COLLIEX (Orsay, FRANCE), Jörg ENDERLEIN (Göttingen, GERMANY), Andreas ENGEL (Delft, THE NETHERLANDS), Ernst H.K. STELZER (Professor) (Frankfurt am Main, GERMANY)
10:30 - 11:00 #7320 - IM03-S40 Three-dimensional nanomechanical spectroscopy of soft matter-liquid interfaces.
Three-dimensional nanomechanical spectroscopy of soft matter-liquid interfaces.

This contribution is devided in two sections. The first section is devoted to examine some relevant issues regarding force microscopy imaging of biomolecules such as spatial resolution, molecule deformation and quantitative mapping of mechanical properties. Specifically, we present a method to obtain the stress-strain curve of a single protein in liquid. The second section is devoted to present an advanced AFM method to genearte three dimensional maps of  solid-liquid interfaces. We develop a force microscope method to map the 3D structure of solid-liquid interfaces. The maps provide atomic-scale spatial resolution images of the formation of hydration layers and the adsorption of ions on solid-liquid interfaces.  Some applications include to resolve the atomic structure of hydration layers generated from alkali chloride solutions on an atomically flat mica surface. 

 

D. Martin-Jimenez, E. Chacon, P. Tarazona, R. Garcia, Submitted    

E.T. Herruzo, A.P. Perrino and R. Garcia, Nature Commun. 5, 3126  (2014)     

R. Garcia, E.T. Herruzo, Nat. Nanotechnol. 7, 217-226 (2012)

                                                                                                                                                                                                                                                                                                         

 

Ricardo GARCIA (Madrid, SPAIN)
Invited
11:00 - 11:15 #7065 - IM03-OP102 Assessment of doping profiles in semiconductor nanowires by scanning probe microscopy: Study of p- type doping in ZnO nanowires.
Assessment of doping profiles in semiconductor nanowires by scanning probe microscopy: Study of p- type doping in ZnO nanowires.

Methods to measure and quantitatively determine the doping profile in semiconductor nanowires ( NW)  are strongly requested for understanding the doping incorporation in such  one-dimensional  structures and so for developing technology using them. In the last two decades, scanning capacitance microscopy (SCM) and scanning spreading resistance microscopy (SSRM) based on atomic force microscopy, has emerged as promising tools for two-dimensional high resolution carrier/dopant profiling. In SCM, the capacitance change providing by an alternating bias applied between the tip/sample system under a DC bias to alternately accumulate and deplete carriers within the semiconductor underneath the local tip is dependent on the local carrier concentration of the semiconductor. In SSRM, the local resistivity is determined via the resistance measurement at the tip/sample system allowing the determination of the doping concentration. These two techniques need of an accurate calibration method for a quantitative doping analysis.

In this communication, we present first a calibration method based on cross-sectional scanning of multilayers samples with different Ga doping concentration allowing the quantitative measurement of n-type ZnO doping by SCM and SSRM. Then, to study ZnO NWs, we have developed a methodology of sample preparation, based on dip-coating filling of NWs field. The dip-coating parameters as coating solution, removal rate and NW field morphology have been controlled by SEM, ellipsometry and atomic force microscopy topography in order to optimize the filling and polishing process.  

One important results has been to be able to measure using SCM and SSRM,  the non-intentionally n-type doping (nid)  of the ZnO nanowires, well estimated at 21018cm-3, explaining  the difficulty to turn these NWs  into p-type  during p-type doping experiments, a crucial problematic in ZnO.  Using antimony (Sb) doping in nid ZnO core/ Sb ZnO shell NW structures, we have successfully determined the decrease of carrier concentration with respect to the nid core ZnO, which can be ascribed to the formation of Sb-related acceptors compensating the native donors. The understanding of this electric compensation mechanism is the clear signature of  p-type Sb doping  feasibility in ZnO NW.  This important result opens the way to succeed in the p-type doping in ZnO.

The generalization of this doping profiling methodology to other semiconductors NW could be pointed out.

Georges BREMOND (Villeurbanne), Lin WANG, Jean-Michel CHAUVEAU, Corine SARTEL, Vincent SALLET
11:15 - 11:30 #5301 - IM03-OP085 Cathodoluminescence microscopy of biological samples for correlative light and electron microscopy (CLEM) using organic fluorophores.
Cathodoluminescence microscopy of biological samples for correlative light and electron microscopy (CLEM) using organic fluorophores.

A CLEM study in biology aims at providing knowledge on the identity and localisation of cellular constituents of interest within a cell's ultrastructure. The key limitation of such experiments, however, is the registration precision between light and electron microscopic data. Precise registration is usually achieved using fiducial markers, visible in both imaging modalities [1].

Integrated microscopes, i.e. light and electron microscopy are performed in one machine [2], in contrast, offer an inherent high-precision correlation. In such systems, photon emission from the sample, is either triggered by a light source (e.g. laser) or the electron beam. A glass objective lens, mounted below the sample collects the emitte light, that is ultimately detected by a photomultiplier tube or a camera [3].

This geometry is particularly advantageous for detecting a sample's cathodoluminescence (CL) signal, since roughly 80% of the emitted photons are emitted into the forward direction [3], and it allows for simultaneous, unobstructed secondary electron (SE) imaging.

We evaluate a CL detector similar to the one described by Narváez et al. [3], with respect to its applicability to CL imaging of small synthetic fluorophores. As these fluorophores are readily damaged by the electron beam [4], we study the CL signal of these molecules dependent on primary energy and beam current. Based on the results of these experiments we image the CL signal of fluorescently labelled, resin embedded biological specimen. Our data indicate the feasibility of CL imaging for CLEM of biological specimen.

The CL signal of 200 nm sized fluorescently labelled polystyrene beads, scales with primary energy (figure 1). Increasing the beam current increases the CL intensity (figure 2A) of individual images. Comparing image series taken at different beam currents (240 pA (240 frames), 480 pA (120 frames) and 1440 pA (40 frames)), a dose rate effect on the cumulative retrievable intensity (figure 2B) is observed. The signal-to-noise ratio can be maximized by either repeated scanning of the same area on the sample at low beam currents or by short pixel dwell times at higher beam currents. In both cases the final image is obtained by averaging the acquired frames.

Analyses of the influence of beam current and primary energy on the CL signal were performed at a pixel dwell time of 6.4 µs. Cumulative intensities in figure 2 are presented as mean ± standard deviation of 4 different areas on the sample per condition.

Having established that CL imaging of organic fluorophores is feasible, we imaged DNA stained with 1 µM Sytox® Green in mammalian cells. Figure 3 shows CL and SE images of 100 nm sections, of LR White embedded Hela cells before and after staining, deposited on ITO (Indium Tin Oxide) coated cover slips. CL signal of stained cell nuclei was detected at different magnifications. Following the imaging guidelines established for fluorescently labelled beads, CL of stained DNA was recorded by several fast scans (100 ns pixel dwell time) of the same position and averaging of the individual frames. Sections were imaged at 1 kV (unstained cells) and 2 kV (stained cells) primary energy, respectively, and a beam current of 480 pA.

Currently we investigate, whether low temperatures (120 K) increase the beam stability of the CL signal. Another factor we currently address is the background signal in images of resin embedded samples (Figure 3), which results from photon emission upon electron beam excitation of the resin or the glass substrate.

References:

[1] Kukulski et al., J. Cell Biol. 192 (2011), p. 111

[2] Peddie et al., Ultramicroscopy 143 (2014), p. 3

[3] Narváez et al., Opt. Express 21 (2013), p. 29968

[4] Niitsuma et al., J Electron Microsc 54 (2005), p. 325

[5] Our work is funded by the German Federal Ministry of Education and Research (13GW0044).

[6] We thank Marina van Ark for technical assistance.

Christopher SCHMID (Heidelberg, GERMANY), Klaus YSERENTANT, Lucian STEFAN, Dirk-Peter HERTEN, Rasmus SCHRÖDER
11:30 - 12:00 #8365 - IM03-S41 Quantitative coherent Raman scattering microscopies: new tools for the material and life sciences.
Quantitative coherent Raman scattering microscopies: new tools for the material and life sciences.

Unlike optical microscopies that are based on fluorescence detection, Raman-based micro-spectroscopies provide vibrational signatures that themselves represent quantitative measures of the sample’s molecular composition and structures, which for example can be successfully exploited as an intrinsic vibrational contrast of endogenous biomolecular species for label-free tumour diagnostic imaging [1]. In particular, by exploiting the coherent driving and detection of Raman modes in coherent anti-Stokes Raman scattering (CARS) and stimulated Raman scattering (SRS), coherent Raman scattering (CRS) microscopy allows the point-by-point chemical mapping of molecular compounds, which is often difficult to attain by conventional fluorescence and incoherent vibrational microscopy techniques. Here, we will review on two CRS modalities that provide quantitative molecular information [2]: (i) high-speed stimulated Raman loss (SRL) imaging at video-rates and (ii) hyperspectral CARS imaging that provides access to the full wealth of chemical and physical structure information of an a priori unknown molecular sample. We will discuss their underlying principles, their state-of-the-art experimental realizations, and demonstrate exemplifying applications for the label-free and noninvasive 3D visualization of chemical composition as well as of molecular structure properties of (bio)molecular components in heterogeneous and complex materials, ranging from polymers to living cells.

Particular emphasis will be given to the combination of coherent Raman scattering spectroscopy with optical microscopy, which has emerged as a highly sensitive and chemically selective tool for the extraction of quantitative molecular structure information from purely imaginary hyperspectral data cubes of the sample’s complex third-order nonlinear susceptibility, χ(3)(ν,x,y,z), as obtained by fast hyperspectral CARS imaging in conjunction with spectral phase retrieval algorithms. We will introduce a novel concept based on the wavelet prism decomposition and the maximum entropy method (MEM) for the fast and robust reconstruction of the pure vibrational response of the molecular sample inside a sub-femtoliter probe volume in the presence of experimental artefacts, which may obstruct the accurate phase retrieval from the experimental normalized CARS pixel spectra [3]. Furthermore, we will present exemplifying applications (see Fig. 1) for the quantitative 3D mapping of chemical composition in living cells, the intracellular chemical structure analysis of biologically relevant lipids, and the physical 3D structure analysis in polymers.

[1] P. Piredda, M. Berning, P. Boukamp, A. Volkmer, Subcellular Raman Microspectroscopy Imaging of Nucleic Acids and Tryptophan for Distinction of Normal Human Skin Cells and Tumorigenic Keratinocytes, Anal. Chem. 87 (2015) 6778−6785.

[2] A. Volkmer, Chapter 6: Coherent Raman scattering microscopy, in Emerging Biomedical and Pharmaceutical Applications of Raman Spectroscopy Eds. P. Matousek and M. Morris, 2010 (Springer-Verlag) 111-152.

[3] Y. Kan, L. Lensu, G. Hehl, A. Volkmer, E.M. Vartiainen, Wavelet prism decomposition analysis applied to CARS spectroscopy: a tool for accurate and quantitative extraction of resonant vibrational responses, Opt. Express 24 (2016) 11905-11916.

Acknowledgements. I would like to extent a special thanks to my present and past co-workers G.Hehl, S. Gomes da Costa, H. Barbosa de Aguiar, P. Piredda, R. Venkatnarayan, A. Kovalev, and N. Nandakumar at the University of Stuttgart, as well as to my collaborators  E.M. Vartiainen of Lappeenranta University of Technology,  P. Boukamp of the German Cancer Research Center (DKFZ), M. Schmitz of the University Hospital Regensburg, and G. Nikolaeva of the RAS General Physics Institute for their essential contributions to the presented work. Financial support from the Deutsche Forschungsgemeinschaft (DFG: VO 825/1-2, 1-3, 1-4), the European Union (HEALTH-F5-2008-200820 CARS EXPLORER), and from the German Federal Ministry of Education and Research (BMBF) of the projects MEDICARS (13N10776) and MIKROQUANT (13N11074) is gratefully acknowledged.

Andreas VOLKMER (Stuttgart, GERMANY)
Invited
12:00 - 12:15 #6059 - IM03-OP092 Soft X-ray tomography enhancement by focal series projections.
Soft X-ray tomography enhancement by focal series projections.

Soft X-ray tomography (SXT) is one of the most recent structural biology techniques for quantitative three-dimensional (3D) analysis of whole cells.  When preserved in cryo-conditions, whole cells can be imaged by Fresnel zone plate (FZP) lenses with spatial resolution in the few tens of nanometers. Information from these images can then be combined by tomographic techniques to generate 3D maps of the specimen.  X-ray photon-energies in the so-called water window (between 284 and 543 eV) are especially useful, since high contrast images of cryopreserved whole cells of ~10 microns can be obtained without use of staining agents. However, SXT is far from currently being delivering its promise to its full extend, mostly due to the intrinsic limited depth of field (DOF) of the microscope optical system and to the over simplicity of the standard reconstruction algorithms being used. Indeed, this limited DOF produces images where the contribution of each plane of the specimen is different, leading to three-dimensional (3D) maps where certain parts of the specimen are blurred. We analyze the SXT image formation process in detail, providing a new mathematical theory allowing for the quantitative SXT inversion using focal series projections, images acquired at different defocus (Figure 1). Following these ideas, the SXT microscope at the Spanish synchrotron ALBA has been modified so as to automatically acquire the necessary additional information in a fast and reliable manner. Examples of the use of the new formulation on several biological systems using images acquired at ALBA are provided.

Here we apply our quantitative approach to SXT inversion to the study of an experimental biological system: a rotavirus infected cell tomogram (Figure 2).

Joaquin OTON (MADRID, SPAIN), Eva PEREIRO, Jose Javier CONESA, Carlos Oscar S. SORZANO, Roberto MARABINI, José L. CARRASCOSA, Jose María CARAZO
12:15 - 12:45 #8636 - IM03-S42 Beating Time and Space resolutions in Ultrasound for disruptive innovations in Medical Imaging.
Beating Time and Space resolutions in Ultrasound for disruptive innovations in Medical Imaging.

In the last fifteen years, the introduction of plane or diverging wave transmissions rather than line by line scanning focused beams has broken the conventional barriers of ultrasound imaging. By using such large field of view transmissions, the frame rate reaches the theoretical limit of physics dictated by the ultrasound speed and an ultrasonic map can be provided typically in tens of micro-seconds (several thousands of frames per second). Interestingly, this leap in frame rate is not only a technological breakthrough but it permits the advent of completely new ultrasound imaging modes, including shear wave elastography, electromechanical wave imaging, ultrafast doppler, ultrafast contrast imaging, and even functional ultrasound imaging of brain activity (fUltrasound) introducing Ultrasound as an emerging full-fledged neuroimaging modality.

At ultrafast frame rates, it becomes possible to track in real time the transient vibrations – known as shear waves – propagating through organs. Such "human body seismology" provides quantitative maps of local tissue stiffness whose added value for diagnosis has been recently demonstrated in many fields of radiology (breast, prostate and liver cancer, cardiovascular imaging, ...). Today, first clinical ultrafast ultrasound scanners are available in the clinical world with such real time imaging of tissue elasticity. This is the first example of an ultrafast Ultrasound approach now widely spread in the clinical medical ultrasound community.

 

For blood flow imaging, ultrafast Doppler permits high-precision characterization of complex vascular and cardiac flows. It also gives ultrasound the ability to detect very subtle blood flow in very small vessels. In the brain, such ultrasensitive Doppler paves the way for fUltrasound (functional ultrasound imaging) of brain activity with unprecedented spatial and temporal resolution compared to fMRI.

Combined with contrast agents, our group demonstrated that Ultrafast Ultrasound Localization could provide a first in vivo and non-invasive imaging modality at microscopic scales deep into organs.

Many of these ultrafast modes should lead to major improvements in ultrasound screening, diagnosis, and therapeutic monitoring.

Mickael TANTER (PARIS CEDEX 5)
Invited

10:30-12:45
Added to your list of favorites
Deleted from your list of favorites

MS0-III
MS0: Nanoparticles: from synthesis to applications
SLOT III

MS0: Nanoparticles: from synthesis to applications
SLOT III

Chairmen: José CALVINO (Cadiz, SPAIN), Goran DRAZIC (Head of microscopy group) (Ljubljana, SLOVENIA), Christian RICOLLEAU (Professor) (Paris, FRANCE)
10:30 - 10:45 #5085 - MS00-OP180 Quantification of SiO2 nanoparticle sedimentation on A549 cells.
Quantification of SiO2 nanoparticle sedimentation on A549 cells.

In recent years many studies were published on toxicological effects of nanoparticles (NPs) on human tissue. Many dose-response studies rely on in vitro assays, in which cultured cells are exposed to a suspension of cell-culture medium and NPs. Limbach et al. [1] stated, that the sedimentation process of NPs in this setup is more complex than in the case of microparticles (MPs). While sedimentation of NPs is primarily driven by diffusion, it is mainly gravitational force, which influences MPs. To compare and quantify the sedimentation of particles of both scales, sedimentation studies and simulations with SiO2 particles were performed in this work. Scanning electron microscopy (SEM) was applied to measure the direct cellular dose, i.e., the sedimented areal densities of particles (AD) on the cells, as a more appropriate definition of dose according to Teeguarden et al. [2]. Simulations are based on ISDD, a computational model by Hinderliter et al. [3].

The samples were prepared as follows: A549 lung cancer cells were seeded onto indium-tin-oxide coated glass substrates in culture plates containing Dulbecco's Modified Eagle's Medium supplemented with fetal calf serum (FCS). After adhering overnight, the cells were incubated with SiO2 particles (from 70 nm up to 500 nm diameter) in cell culture medium for 1 and 4 hours. Using ISDD and assuming, that sedimented particles are homogeneously distributed, targeted ADs were calculated according to predefined incubation concentrations. After fixation with paraformaldehyde the cells were dehydrated with graded ethanol series, dried by critical point drying and investigated in a FEI Quanta 650 SEM. Particles were imaged with secondary electrons (SE) in intercellular regions between cells. Backscattered electron (BSE) images were taken to detect particles on cells. To increase the BSE contrast, a retarding bias was applied to the sample stage.

Fig. 1a shows a representative SE image of an intercellular region with 200 nm SiO2 particles after 1 h incubation. The particles appear homogeneously distributed. Fig. 1b depicts a cell surface of the same specimen with a substantially smaller AD (0.15 NP/µm² compared to 0.84 NP/µm² in Fig. 1a) and an inhomogeneous NP distribution. A smaller cellular AD is also observed for all other particles sizes and incubation times. This is shown in Fig. 2a, where measured cellular and intercellular ADs for each specimen are compared. In Fig. 2b, the simulated AD is plotted against the measured intercellular AD for all samples. Calculated and measured ADs agree well within the error bars, which can be considered as a verification of the ISDD model. However, the cellular ADs are significantly lower indicating that another effect must be taken into account. Cellular uptake can be ruled out as an explanation, because focused-ion-beam sectioning of whole cells did not show a high particle density in cells.

Fluorescence microscopy (FM) investigations with rhodamine labelled SiO2 NPs (Ø = 70 nm) and A549 cells qualitatively confirm the differences between cellular and intercellular ADs observed by SEM (Fig. 3a). The red fluorescence is much stronger in intercellular regions than on cells, however the signals stem from small agglomerates. Dynamic light scattering shows, that this agglomeration is caused by FCS coating of NPs applied before incubation. To study the impact of FCS in more detail, further in vitro experiments with and without FCS precoatings of NPs and/or substrates were performed. Preliminary results indicate, that protein coatings induce an attractive interaction between NPs with their protein corona and surface proteins. Since FM is unable to resolve single NPs, SEM is best suited for confirmation.

SEM is convenient to quantitatively determine ADs of NPs and reveals distinct differences between cellular and intercellular ADs. Our results highlight a major problem of bulk investigation methods relying on lysates, because intercellular and cellular regions cannot be distinguished.

References

[1] L.K. Limbach, et al., Environ. Sci. Technol., 39 (2005), pp. 9370–9376.

[2] J.G. Teeguarden, et al., Toxicol. Sci., 95 (2007), pp. 300–312.

[3] P.M. Hinderliter, et al., Part. Fibre. Toxicol., 7 (2010), p. 36.

We acknowledge the support of the BIF graduate school funded by the Helmholtz Association.

Thomas KOWOLL (Karlsruhe, GERMANY), Susanne FRITSCH-DECKER, Regina FERTIG, Erich MUELLER, Carsten WEISS, Dagmar GERTHSEN
10:45 - 11:00 #5767 - MS00-OP181 Evaluation of Electron Microscopy Techniques for the Purpose of Classification of Nanomaterials.
Evaluation of Electron Microscopy Techniques for the Purpose of Classification of Nanomaterials.

One current and much-debated topic in the characterization of nanomaterials (NM) is the implementation of the recently introduced recommendation on a definition of a nanomaterial by the European Commission [1]. According to this definition [1], a material is a NM when for 50% or more of the particles in the number size distribution, one or more external dimensions is in the size range 1 nm-100 nm. The European NanoDefine project [2] was set up to develop and validate a robust, readily implementable and cost-effective measurement approach to obtain a quantitative particle size distribution and to distinguish between NMs and non-NMs according to the definition [1].

All currently available sizing techniques able to address nanoparticles were systematically evaluated. It was demonstrated that particle sizing techniques like: analytical centrifugation, particle tracking analysis, single-particle inductively coupled plasma mass-spectrometry, differential electrical mobility analysis, dynamic light scattering, small angle X-ray scattering, ultrasonic attenuation spectrometry, but also gas adsorption analysis based on the BET-method can be applied for a screening classification. However, the quality of the results depends on the individual material to be classified. For well-dispersed, nearly spherical (nano)particles most of the sizing techniques can be applied in a quick and reliable way. In contrast, the classification of most real-world materials is a challenging task, mainly due to non-spherical particle shape, large polydispersity or strong agglomeration/aggregation of the particles. In the present study it was shown that these issues can be resolved in most cases by electron microscopy as a confirmatory classification technique [3-6].

Electron microscopy techniques such as TEM, STEM, SEM or TSEM (transmission in SEM) are capable of assessing the size of individual nanoparticles accurately (see Figures 1 and 2). Nevertheless the challenging aspect is sample preparation from powder or liquid form on the substrate, so that a homogeneous distribution of well-separated (deagglomerated) particles is attained. The systematic study in this work shows examples where the extraction of the critical, smallest particle dimension - as the decisive particle parameter for the classification as a NM - is possible by analysing the sample after its simple, dry preparation. The consequences of additional typical issues like loss of information due to screening of smaller particles by larger ones or the (in)ability to access the constituent particles in aggregates [5] are discussed.

By means of practical examples the inherent statistical evaluation of the particle size is highlighted together with all its pitfalls such as setting of a suitable threshold for delimitation of the particle boundaries in the electron micrograph or consideration of systematic (bias) deviations from the true particle size because of evaluation via surface sensitive secondary electron detectors, e.g. In-Lens.

[1] European Commission, Commission Recommendation of 18 October 2011 on the definition of nanomaterial (2011/696/EU). Official J. Europ. Union 54 (2011) p. 38.

[2] www.NanoDefine.eu

[3] F Babick et al, submitted.

[4] K Yamamoto, Microsc. Microanal. 21 (Suppl 3) (2015) p. 2399.

[5] P-J de Temmerman et al, Powder Technol. 261 (2014) p. 191.

[6] P Müller et al, Microsc. Microanal. 21 (Suppl 3) (2015), p. 2403.

[7] The research leading to these results has received funding from the European Union’s Seventh

Framework Programme (FP7/2007-2013) under grant agreement n° 604347-2. (www.nanodefine.eu).

Johannes MIELKE, Frank BABICK, Toni UUSIMÄKI, Philipp MÜLLER (Ludwigshafen am Rhein, GERMANY), Eveline VERLEYSEN, Vasile-Dan HODOROABA
11:00 - 11:15 #5959 - MS00-OP184 Playing around with shape and composition of nanoparticles for various applications.
Playing around with shape and composition of nanoparticles for various applications.

Nowadays, nanoparticles with sizes between 2 to 50 nm become more and more popular, because they can be applied in various fields such as materials science, chemistry, catalysis, medicine or biology. In particular nanoparticles with fancy shapes gain a lot of attention, also because of the low-cost synthesis methods. This study is focused on several types of catalytic nanoparticles (NPs), silver nanoparticles synthesized using green chemistry methods and in particular on their morphological and chemical analysis using HR(TEM). Three-dimensional (3D) catalysts [1] being promising for application in fuel cells were studied. These nanoparticles have a dodecahedron shape with Pt skin at the edges and a Ni core, Figure 1(a). When etching away the core, the remaining empty Pt frame offers a much larger active surface compared to spherical nanoparticles. The morphology of these 3D PtNi particles strongly depends on the synthesis parameters allowing fabricating dodecahedrons, Figure 1(a), core-shell or even dendritic structures, Figure 1(b).

SnO2 nanoparticles are excellent supports for noble metal NPs, as their combination exhibits good catalytic activity towards ethanol oxidation reaction [2]. Various synthesis routes including polyol and microwave assisted methods allowed producing different SnO2 NPs, Figure 2(a). The break of the C=C bond in the ethanol molecule occurs at the interface between PtRh and SnO2 particles, Figure 2(b). Therefore their physical contact is imperative for the effectiveness of the catalyst. Structural aspects and chemical analysis by TEM characterization techniques of the PtRh/SnO2/C catalysts were analyzed.

Green synthesis method using camomile extract was applied to synthesize silver nanoparticles in order to tune their antibacterial properties merging the synergistic effect of camomile and Ag [3]. Scanning transmission electron microscopy (STEM) revealed that camomile extract (CE) consisted of porous globular nanometer sized structures, which were a perfect support for Ag nanoparticles, Figure 3(a). The Ag nanoparticles synthesized with the camomile extract (AgNPs/CE) of 7 nm average size, were uniformly distributed on the CE support, Figure 3(b). The EDX chemical analysis showed that camomile terpenoids, Figure 3(c) act as a capping and reducing agent being adsorbed on the surface of AgNPs/CE, Figure 3(d), enabling their reduction from Ag+ and preventing them from agglomeration. Antibacterial tests using four bacteria strains, showed that the AgNPs/CE performed five times better compared to CE and AgNPs/G samples, reducing totally all the bacteria in 2 hours, Figure 4.

 

References

  1. C. Chen, Y. Kang, Z.Huo, Z. Zhu, W. Huang, H.L. Xin, J.D. Snyder, D. Li, J.A. Herron, M. Mavrikakis, M. Chi, K.L. More, Y. Li, N.M. Markovic, G.A. Somorjai, P. Yang, V.R. Stamenkovic, Science 343 (2014) 1339-1343.
  1. A. Kowal, M. Li, M. Shao, K. Sasaki, M.B. Vukmirovic, J. Zhang, N.S. Marinkovic, P. Liu, A.I. Frenkel and R.R. Adzic, Nature Materials 9 (2009) 325-330.
  2. M. Parlinska-Wojtan, M. Kus-Liskiewicz, J. Depciuch, and O. Sadik submitted to Bioprocess and Biosystems Engineering.

Acknowledgements

We thank the Center for Innovation and Transfer of Natural Sciences and Engineering Knowledge of the University of Rzeszow and the Division of Materials Processing Technology, Management and Computer Techniques in Materials Science, Institute of Engineering Materials and Biomaterials, of the Silesian University of Technology in Poland for using the TEM instruments. Financial support from the Polish National Science Centre (NCN), grant UMO-2014/13/B/ST5/04497 is acknowledged.

Magdalena PARLINSKA-WOJTAN (Krakow, POLAND), Grzegorz GRUZEL, Elzbieta ROGA, Joanna DEPCIUCH, Andrzej KOWAL
11:15 - 11:30 #5988 - MS00-OP185 Imaging interactions of iron oxide nanoparticles with organic ligands and the biological environment.
Imaging interactions of iron oxide nanoparticles with organic ligands and the biological environment.

Applications of nanoparticles for chemistry, biochemistry and biomedicine require a detailed understanding of the nature of the nanoparticles and of their surface interactions with both functionalizing organic groups and biological environments. Here we present ongoing studies of the nature and surface interactions of iron oxide nanoparticles (IONPs) intended for magnetic resonance imaging detection and hyperthermia treatment.

The IONPs are synthesized by a novel low temperature aqueous route devoid of surfactant chemistry or capping agents. In addition to routine characterization of the IONPs by TEM, XRD, FTIR, XPS, DLS, and magnetic characterization, we use aberration-corrected high resolution TEM to monitor their structural quality. Because of an observed sensitivity to the electron beam at a 200 kV high tension (HT), we use a HT of 80 kV, with effects of chromatic aberration reduced by implementing a monochromatic “rainbow” illumination in the incident beam [1]. This imaging allows observations of minute changes in surface structural quality. Initially formed particles are rounded, showing signs of slight structural disorder in the first atomic layers at the surface (Fig. 1). In contrast, aged particles have atomically sharp structural ordering at surfaces which are markedly more faceted (Fig. 2). No beam damage effects are observed other than the hopping of atoms on surface ledge sites.

For bioengineering, it is critical to understand how molecules and proteins interact with these IONPs. Beginning with the former, the particles are functionalized with folic acid, the molecule most often used for "nonspecific" targeting. Using the same monochromated, Cs-corrected imaging conditions as above, in Figure 3 we demonstrate the ability to image ligand attachment and surface coverage, similarly to Lee et al. [2]. The folic acid molecules show phase contrast with definition down to about 2 Å; this phase contrast will be compared to simulations for better interpretability. Ligands are observed to form two or three strand thick shells around the surfaces of the IONPs. Using fast frame acquisition, very specific effects of electron irradiation are recorded: the molecules wriggle under the beam, often with one end appearing to detach and reattach to the surface of the IONP. This behavior is explained in terms of relative bonding potentials for either the amino acid or the carboxylic ends of the molecules absorbed to the IONP surfaces.

Once the IONPs are injected into a biological environment it is known that any such functionalization is rapidly replaced by a surface coverage of proteins – the protein corona – which then dictates how the IONPs interact with this environment [3]. To understand this biomedically-important process, for the first time, we obtain in vivo conditions which create a protein corona whose nature mimics that observed for in vitro studies. As well as again using phase contrast imaging of dry samples, we also study the morphology of this corona by simple negative-staining to create contrast between the protein and the carbon support film. Contrary to the commonly held view that the proteins make a uniform encapsulating layer around nanoparticles, their coverage is distinctly patchy (Fig. 4); consistent with the attaching proteins having widely varying sizes. As a next step HAADF STEM tomography will be used to map the 3D morphology of this non-uniform coronal coverage.

With research ongoing on these IONPs for "magnetotheranostics", the other main perspective of this work is to pursue imaging of such interactions in representative liquid environments, particularly using high resolution cryo-TEM imaging to monitor better the nature of organic functionalization in an aqueous environment.

 

References and Acknowledgements

[1] P.C. Tiemeijer et al., Ultramicrosc. 114 (2012) 72–81.

[2] Z. Lee et al., Nano Lett. 9 (2009) 3365–3369.

[3] M.P. Monopoli et al., Nature Nano. 7 (2012) 779–786.

The authors acknowledge funding from Nano-Tera.ch, project Magnetotheranostics number 530 627.

Duncan ALEXANDER (Lausanne, SWITZERLAND), Débora BONVIN, Ulrich ASCHAUER, Heinrich HOFMANN, Marijana MIONIC EBERSOLD
11:30 - 11:45 #6730 - MS00-OP195 Evidence for the dissolution of molybdenum during tribocorrosion of CoCrMo alloy in the presence of serum proteins.
Evidence for the dissolution of molybdenum during tribocorrosion of CoCrMo alloy in the presence of serum proteins.

CoCrMo Metal on Metal (MoM) hip implants were designed to be durable, targeting a better quality of life for young, active patients. Current evidence suggests that such implants can release wear particles and metal ions due to a bio-tribocorrosion process [1,2]. This involves both corrosion of the implant surface itself, which is stimulated by the wear process removing a chromium-rich passivating film, and also mechanical wear of the surface to produce nanoparticulate wear debris that can be spread by lymphatic circulation and subsequently corrode. Both processes can therefore give rise to the release of metal ions which may have an inflammatory or toxic effect on cells and tissues, notably because certain metal ions can complex with proteins and disable the primary function.

This study focuses on the synthesis and characterization of CoCrMo nanoparticles, which mimic metal-on-metal (MoM) wear debris from hip implants. We have used a hitherto unexplored approach of mechanochemical milling to produce a large amount of CoCrMo mimetic wear debris over short time scales [3]. Using TEM, the nanoparticles produced were found to be similar in size, shape and composition to wear debris from CoCrMo hip implants produced both in-vivo and using hip simulators. An efficient separation of nanoparticles from solution prior to free ion concentration analysis by inductively coupled plasma mass spectrometry (ICP-MS) was developed using centrifugation combined with ultrafiltration (2 kDa ultrafilters; Figure 1). This revised preparation method has allowed the identification of the dissolution of specific alloy elements both in hip simulator lubricant systems and in different biological media and pH environments during both dynamic tribocorrosion studies (using milling) and also during static corrosion (Figure 2). The results indicate a much lower dissolution of cobalt than previously reported.  We also identified a significant fraction of Mo ions in solution that could be linked to the preferential binding of Mo by bovine serum albumin (BSA) proteins identified by FTIR and TEM [4; Figure 3]. Electrochemical corrosion tests in the presence of BSA confirm that the proteins play an important role in Mo dissolution from CoCrMo. We suggest that the interaction of Mo-rich surfaces with amide groups in serum proteins and the possible formation of metal carbonyl complexes are important because both can modify biological molecules, potentially altering the function.

In summary, the role of Mo as well as Co ions should be accounted for in the tribocorrosion of CoCrMo implant alloys, particularly in terms of inflammatory and toxicological responses.

 

[1] D. Cohen (2012) How safe are metal-on-metal hip implants?, Br. Med. J., 344, e1410–e1410.

[2] Y. Yan, A. Neville, D. Dowson, S. Williams, J. Fisher (2009) Effect of metallic nanoparticles on the biotribocorrosion behaviour of Metal-on-Metal hip prostheses, Wear, 267, 683–688.

[3] T.A. Simoes, A.E. Goode, A.E. Porter, M.P. Ryan, S.J. Milne, A.P. Brown, et al. (2014) Microstructural characterization of low and high carbon CoCrMo alloy nanoparticles produced by mechanical milling. J. Phys. Conf. Ser., 522, 012059.

[4]T.A. Simoes, A.P. Brown, S.J. Milne, R.M.D. Brydson (2015) Bovine Serum Albumin binding to CoCrMo nanoparticles and the influence on dissolution. J. Phys. Conf. Ser., 644, 012039.

Thiago SIMOES, Michael BRYANT, Andy BROWN, Steve MILNE, Mary RYAN, Angela GOODE, Alexandra PORTER, Anne NEVILLE, Rik BRYDSON (Leeds, UK)
11:45 - 12:00 #6517 - MS00-OP191 Chemical arrangement and surface effects in CoAu nanoparticles.
Chemical arrangement and surface effects in CoAu nanoparticles.

Nanoparticles associating a noble metal and a ferromagnetic metal are appealing from a magneto-plasmonics point of view, in addition to the problematics of magnetic anisotropy tailoring (interface anisotropy, phase transformation) and of nanoalloy original geometries. Because Co and Au are immiscible in the bulk phase, and since fcc cobalt and gold have highly different cell parameters, chemically separated structures (core-shell type) are expected for nanoparticles.

We have studied CoAu cluster assemblies, with a diameter between 3 and 10 nm, prepared by low energy cluster beam deposition (LECBD) where nanoparticles are formed in out-of-equilibrium conditions by laser vaporization, then deposited on a substrate under ultrahigh vacuum conditions and protected by a capping layer (amorphous carbon, to avoid oxidation). The nanoparticles’ structure and chemical arrangement (see figures) have been investigated by HRTEM, STEM-HAADF and STEM-EELS before and after annealing (2h around 500°C). A mapping of the low energy electronic excitations has also been performed by STEM-EELS, which is a challenge on such small nano-objects.

As prepared particles are found to be inhomogeneous (as deduced from EELS measurements), with interatomic distances always corresponding to pure gold, and they appear to be surrounded by a shell of lighter HAADF intensity which rapidly transforms upon electron beam exposure in STEM. After annealing, a phase separation is observed and CoAu nanoparticles adopt a core-shell structure where, as observed by HRTEM (with a clearly visible difference of inter-plane distances between Co and Au regions), STEM-HAADF and STEM-EELS, an off-centered cobalt core is surrounded by a gold shell (see figure). For both as-prepared and annealed nanoparticles, the Co is not oxidized thanks to the efficient protection of the thin carbon layer. Using low-loss STEM-EELS mapping, we are able to observe some surface contributions, which may reflect collective electronic resonances (surface plasmons), with significant differences between as-prepared and annealed CoAu nanoparticles (in particular, an intense peak is detected on the gold-rich regions, around 6 eV).

These results shed light on the atomic-scale behavior of the Co-Au nanoalloy, here in an amorphous carbon matrix, which will be compared to other dielectric matrices, and will help us to understand the original magnetic properties of these hybrid nanoparticles. 

Florent TOURNUS (VILLEURBANNE CEDEX), Ophelliam LOISELET, Kazuhisa SATO, Katia MARCH, Odile STEPHAN
12:00 - 12:15 #6318 - MS00-OP189 Revealing the symmetry breaking, growth and surface properties of gold nanorods.
Revealing the symmetry breaking, growth and surface properties of gold nanorods.

The unique properties of nanoparticle structures are directly determined by particle composition, size, shape and surface faceting; each offering a tuneable parameter with which one can tailor properties to a given application. Advances in wet-chemistry techniques now enable the synthesis of nanoparticles with shape anisotropy and novel surface faceting that exhibit exciting optical and catalytic properties.  Here, we develop and apply quantitative scanning transmission electron microscopy methods to observe key structural changes that enable anisotropic growth; characterise the nanorod faceting in three dimensions; and probe the relative surface “energies” of both high and low index facets.

Gold nanorods are an archetypal system with which to study anisotropic nanoparticle growth.  Yet despite intense research it remains unclear how or why a single crystal seed particle, with a cubic lattice, grows preferentially in two of six nominally symmetry-equivalent directions. Observations at various stages of gold nanorod growth reveal the onset of asymmetry occurs only in single crystal seed particles that have reached diameters between 4 and 6 nm.  In this size range only, small, asymmetric truncating surfaces with an open atomic structure become apparent, and in the presence of Ag+ ions are stabilised, becoming side facets in the embryonic nanorod structure [1,2].   These results provide the first direct observation of the structural changes that break the symmetry of the seed particle and provide key insights into the mechanism of anisotropic growth.

The various facets exhibited by the nanoparticle [3] and their relative surface energies are a crucial driver of shape control whilst also directly determining how the resulting particle will interact with its environment.  We apply a quantitative STEM technique [4] to count the number of atoms in each atomic column as identified in STEM images of a single crystal gold nanorod orientated in two different zone axes.  Using this method we are able to determine the morphology and facet crystallography of the nanorod, finding it is comprised of both high {0 1 1+√2} and low {110}, {100} index side-facets, each of comparable size and shape.  Furthermore, by applying this method at successive time intervals and comparing the images it is possible to quantify atomic movement on the surface and therefore determine the relative stability of different crystallographic facets and the overall stability of the nanoparticle shape [5].  These results provide important information on the effect of surfactants on the relative surface energies of high and low index facets, and shed new light on the growth kinetics of Au nanorods.

 

Acknowledgements: This work was supported by the Australian Research Council (ARC) grants DP120101573, DP160104679 and LE0454166

  1. M. J. Walsh, S. J. Barrow, W. Tong, A. M. Funston and J. Etheridge. ACS Nano, 2015. 9(1). 715-724.
  2. W.Tong, M. J. Walsh, P. Mulvaney, J. Etheridge and A.M. Funston.  In preparation 2016
  3. H. Katz-Boon, C. J. Rossouw, M. Weyland, A.M. Funston, P. Mulvaney and J. Etheridge. Nano Letters, 2010. 11(1): p. 273-278
  4. C. Dwyer, C. Maunders, C. L. Zheng, M. Weyland. P. C. Tiejmeijer and J. Etheridge.  Appl. Phys. Lett. 2012, 100, 191915
  5. H. Katz-Boon, M. J. Walsh, C. Dwyer, P. Mulvaney, A.M. Funston and J. Etheridge. Nano Letters. 2015, 15 (3), 1635
Joanne ETHERIDGE (Melbourne, AUSTRALIA), Michael J. WALSH, Hadas KATZ-BOON, Wenming TONG, Christian DWYER, Alison, M. FUNSTON
12:15 - 12:30 #6042 - MS00-OP186 Synthesis and characterization of bimetallic nanorods.
Synthesis and characterization of bimetallic nanorods.

Anisotropic metal nanoparticles (NPs), and especially nanorods (NRs) exhibit interesting optical properties, which arise from their localized surface plasmon. Unlike nanospheres, gold NRs have a longitudinal surface plasmon resonance (LSPR) in the visible or near-infrared range of the spectrum. By altering e.g. the shape or the dimensions (aspect ratio) of the NRs, the LSPR can be tuned, which makes them interesting materials for a broad range of light based applications, such as photocatalysis [1], data storage [2] and photothermal applications. The optical properties of gold nanoparticles can be extended even further by introducing a second metal. However, synthesizing bimetallic NRs with a good control over the metal composition and distribution while retaining the rod shape is challenging.

In this study we present bimetallic systems composed of gold-based NRs coated with a protective mesoporous silica layer. We show that it is possible to synthesize bimetallic core-shell nanorods within a mesoporous silica shell, by etching away part of the gold and overgrowing the remaining Au-core with a second metal while precisely controlling the core-size, metal-shell thickness and thus the metal-to-metal ratio [3]. Depending on the choice of metal, different growth behavior was observed. Overgrowth with Ag resulted in a smooth shell whereas the Pt and Pd metal shells had a rough morphology (Figure 1). The different types of bimetallic NRs were characterized in detail with advanced electron microscopy techniques such as Energy-dispersive X-ray spectroscopy (EDX), high-angle annular dark-field scanning transmission electron microscopy (HAADF-STEM) and electron tomography. Subsequently, we used these silica coated bimetallic core-shell structured rods as a starting material to make fully alloyed NRs, whereby the two metals were mixed via thermal treatment without loss of anisotropy [4]. The alloying process was followed in detail with in-situ HAADF-STEM heating and EDX measurements by making use of a special heating holder (Figure 2).

[1]           Suljo Linic et al., Nat. Mater. 14, (2015), 567-576. 
[2]           P. Zijlstra et al., Nature, 459, (2009), 410–413
[3]           T.S. Deng, J.E.S. van der Hoeven, A. Yalcin, H.W. Zandbergen, M.A. van Huis, A. van Blaaderen, Chem. Mater. 27, (2015), 7196-7203. 
[4]           W.A. Albrecht and J.E.S. van der Hoeven, T. Deng, P.E. de Jongh, A. van Blaaderen., submitted

Jessi VAN DER HOEVEN (Utrecht, THE NETHERLANDS), Wiebke ALBRECHT, Tiansong DENG, Petra DE JONGH, Alfons VAN BLAADEREN
12:30 - 12:45 #6930 - MS00-OP198 Ceria ro(a)d to cubes: a combined experiment and MD simulation study.
Ceria ro(a)d to cubes: a combined experiment and MD simulation study.

Since the first nanoparticle syntheses, research has focused on producing various nano-catalysts of different compositions and uses. Over the years, the importance of their shape, hence their exposed facets, grew in interest. The differences between facets, may sometimes remain unclear and controversial, but more and more effects in terms of reactivity and/or selectivity are shown in recent works. In this framework, rod and cubic-shaped cerium oxide nanoparticles were synthesized using a hydrothermal microwave-assisted process. Systematic TEM imaging and analysis reveal that rods are progressively replaced by cubes with increasing synthesis time (with fixed temperature and pH). Besides, a more careful look at these cubic-shaped particles highlights a slight anisotropy - depending on the zone axis they are imaged along - leading us to reconsider their nucleation and growth mechanisms.

These processes were first studied using radiation damages induced in ceria rods’ complex microstructures under focussed electron beam. Their periodical packing composed of single crystals aligned along a common low-index axis were firstly reduced and new nanometer-sized domains were formed with different (micro)structures. These zones tend to split from the initial rod in case of severe irradiation conditions and may, in-fine, lead to cuboidal particles via a solid-state transformation process.

As-synthesized cubes were then characterized by means of high resolution TEM coupled with molecular dynamics simulations. As expected, they exhibit single crystalline structure and {001} lateral facets are enclosed by {111} sharp corners and {011} average flat edges made of alternating {111} steps. More surprisingly, whatever the particle size might be, projected edges and corners width remains constant and corresponds to the diagonal of a four by four square formed by {002} planes. By keeping constant this truncation, surface plane ratios can be tuned only by adjusting ceria particles synthesis conditions.

Further on, when submitted to reducing conditions, both simulations and experiments tend to show a partially reversible flattening of the {011} edges combined with the observation of superlattice reflections and Moiré fringes. This indicates that oxygen anions have been removed from the fluorine CeO2 structure inducing an ordered oxygen vacancies lattice. Finally, oxygen vacancies movement into the ceria network will be discussed.

Uli CASTANET (PESSAC), Francesco CADDEO, Dean C. SAYLE, Jerome MAJIMEL

10:30-12:45
Added to your list of favorites
Deleted from your list of favorites

MS8-I
MS8 Geology and mineralogy, cultural heritage and archeology
SLOT I

MS8 Geology and mineralogy, cultural heritage and archeology
SLOT I

Chairmen: Trevor ALMEIDA (Glasgow, UK), Nicolas MENGUY (Professeur) (PARIS, FRANCE)
10:30 - 11:00 #8622 - MS08-S87 Magnetic microscopy of metallic meteorites: probing the magnetic state of the early solar system.
Magnetic microscopy of metallic meteorites: probing the magnetic state of the early solar system.

Meteorites are fragments of asteroids. They represent the oldest and most primitive materials in the solar system – rubble left over after the planets formed over four and half billion years ago. Information about the magnetic state of asteroids during the early solar system can, in principle, be recovered from meteorites. The paleomagnetic potential of meteoritic metal has been overlooked in the past, due to the fact that they are mostly comprised of kamacite (a soft magnet that makes a notoriously poor paleomagnetic recorder). Recent research, however, has uncovered regions buried within the Widmanstätten pattern that can capture reliable records of magnetic activity on asteroid bodies1. This discovery, combined with the advent of high-resolution magnetic imaging methods, has allowed us to decipher the magnetic signals encoded within meteoritic metal for the very first time2,3.

Our attention has focussed on the ‘cloudy zone’: a region just a few microns wide lying next to the kamacite lamellae that define the Widmanstätten pattern. The cloudy zone consists of a tightly packed array of nanoscale islands of tetrataenite, an ordered Fe0.5Ni0.5 phase that forms by the diffusive rearrangement of Fe and Ni atoms into alternating layers during slow cooling. The high intrinsic coercivity (> 2 Tesla) of tetrataenite makes it an excellent permanent magnet. Each island is ~100 nm or less in diameter, and is uniformly magnetised in one of six crystallographically defined directions. The proportions of islands magnetised in each direction are biased by interaction with the magnetic field of the asteroid. This bias is detected using X-ray photo-emission electron microscopy (XPEEM). Magnetic contrast is obtained using circularly polarised X-rays incident at a glancing angle to the surface, exploiting the resulting X-ray magnetic circular dichroism (XMCD) signal to reveal intricate nanoscale domains within the cloudy zone (Fig. 1). Combined with extensive image simulations, quantitative information about the magnetic state of the asteroid can be extracted from the data.

In this talk I will present an overview of nanopaleomagnetic XPEEM studies we have performed on metal from a range of meteorites that have experienced very different cooling rates on their parent bodies, from the most slowly cooled (<< 1 K/Myr) mesosiderites to the most rapidly cooled (>> 1000 K/Myr) IVA. The cooling rate is shown to have a dramatic effect on the length scale of the developing microstructures, with correspondingly dramatic effects on the magnetic properties. In the most extreme cases, six zones of distinct magnetic response are resolved over a 5 µm field-of-view, each zone defined by the properties of the underlying nanostructures. Combined with new insights into the chemical and crystallographic nature of the cloudy zone, as revealed by EDX tomography, atom probe tomography and scanning precession electron diffraction, we are discovering ever more about the underlying physics of these unique materials, and beginning to piece together the processes that helped shape some of the earliest objects in the solar system.

 

References:

1 J.F. Bryson, N.S. Church, T. Kasama, and R. Harrison, Earth Planet. Sci. Lett. 388, 237 (2014).

2 J.F. Bryson, J. Herrero-Albillos, F. Kronast, M. Ghidini, S.A. Redfern, G. Laan, and R. Harrison, Earth Planet. Sci. Lett. 396, 125 (2014).

3 C.I.O. Nichols, J.F.J. Bryson, J. Herrero-Albillos, F. Kronast, F. Nimmo, and R.J. Harrison, Earth Planet. Sci. Lett. 441, 103 (2016).

 

Acknowledgements:

R.J.H. would like to acknowledge funding under ERC Advance grant 320750- Nanopaleomagnetism. P.A.M. would also like to acknowledge funding under ERC Advance grant 291522 - 3DIMAGE.

Richard HARRISON (Cambridgeshire, UK), James BRYSON, Claire NICHOLS, Julia HERRERO-ALBILLOS, Florian KRONAST, Josh EINSLE, Paul MIDGLEY
Invited
11:00 - 11:15 #5814 - MS08-OP308 Mg-calcite formation in a freshwater environment (Lake Balaton): nucleation, growth, structure and composition.
Mg-calcite formation in a freshwater environment (Lake Balaton): nucleation, growth, structure and composition.

Lake Balaton can be regarded as a large scientific laboratory in which many interesting aspects of carbonate mineral formation can be studied, with relevance to the general understanding of nanoscale processes that govern crystal nucleation and growth in a natural aqueous system. The lake is extremely shallow for its size (on average 3.5 m deep and 70 km long), its calcareous water Mg-rich (with a Mg/Ca molar ratio ranging from (~1 to 4), typically displaying a chemical gradient along the W–E long axis of the lake, as a result of the main inlet and outlet being located at opposite ends. The bottom sediment is a soft grey mud, with 30 to 80% of it consisting of Mg-calcite, a mineral that precipitates from the water (Fig. 1a). We used SEM, TEM and STEM techniques to characterize this Mg-calcite, in order to obtain a better understanding of its formation and role in the ecosystem.

 

Freshwater calcite is known to nucleate on biological material, primarily on picoplankton (cyanobacterial) cells. In contrast, the Mg-calcite in Lake Balaton is closely associated with few-unit-cell-thick stacks of clay (smectite) layers (Figs. 1b and c), suggesting that it either nucleated on smectite or adsorbed to the clay flakes. Our laboratory experiments confirm that adding smectite to filtered lakewater dramatically induces Mg-calcite precipitation; thus, the nm-scale clay fragments likely serve as nucleation sites. Since the lake sediments are stirred up by even gentle winds, most of the time smectite particles are readily available for Mg-calcite nucleation. A special case occurs when the lake is frozen and even the clay particles can settle: only bacterial cells are available as nucleation sites (Fig. 2a), and encrustation of cells results in tube-shaped, porous Mg-calcite particles (Fig. 2b).

 

Mg-calcite that formed under „normal” conditions (i.e., nucleated on smectite) typically occurs in the shape of elongated, several μm-large, aggregate-looking particles (Fig. 2c). Even though they appear to be composed of many smaller crystals, SAED patterns suggest that the particles are perfect single crystals (Fig. 2d). The single crystalline nature of the particles probably results from a dissolution/reprecipitation process that preserves the original shapes of particles.

 

The Mg content of the calcite varies from 2 to 20 mol%, depending on the water budget (dilution) and geographical location in the lake. Mg-calcite that nucleated on bacterial cells is highly enriched in Mg; according to SAED patterns, Mg and Ca ions do not order in the structure. Our results provide new information on freshwater calcite nucleation and on the biologically assisted precipitation of high-magnesian calcite.

 

 

Acknowledgements: This research was supported by NKFIH grant no. K116732. Access to electron microscopes at Forschungszentrum Jülich was provided by the EU 7th Framework Programme ESTEEM2. 

Mihály PÓSFAI (Veszprém, HUNGARY), Ilona NYIRŐ-KÓSA, Ágnes ROSTÁSI, Éva BERECZK-TOMPA, Ildikó CORA, Maja KOBLAR, András KOVÁCS
11:15 - 11:30 #6002 - MS08-OP309 Interface migration mechanism on Corundum/Spinel/Periclase: atomic study via aberration-corrected STEM.
Interface migration mechanism on Corundum/Spinel/Periclase: atomic study via aberration-corrected STEM.

 In nature it is common that a new mineral grows between two minerals due to the inter-diffusion of elements. Understanding its growth mechanism is critical for reconstructing conditions and rates of mineral formation. The growth of the new phase is controlled by the coupling of interface reaction and long-range diffusion. To understand the interface reaction, it is essential to figure out the atomic structure of the interfaces.

 In this research, spinel (MgAl2O4, Spl) has been grown between periclase (MgO, Per) and corundum (Al2O3, Crn) via pulsed laser deposition [1] and uniaxial stress methods [2] for studying the early and late growth stages, respectively. Electron Backscatter Diffraction (EBSD) has been used to study the interfacial orientation relationship on both Per/Spl and Spl/Crn reaction interfaces. The EBSD mapping (Fig.1) shows that the Spl layer splits into two different sections: a thinner part in topotaxial orientation relationship with Per, and a thicker part topotaxial with Crn. Then Focused Ion Beam (FIB) was used to lift out the interface areas with representative orientation relationships, and the atomic structure studied by aberration-corrected Scanning Transmission Electron Microscopy (STEM).

 The atomic resolution images of the Spl/Crn interface show that the interface is located where the (001) lattice plane of Crn coincides the (111) lattice plane of Spl, which are both occupied by Al atoms exclusively. In another side, the Per/Spl interface shows a periodic configuration (Fig.2a), consisting of curved segments (convex towards Per) [3]. The image in Fig.2b reveals regularly spaced misfit dislocations at the positions of “cusps” (see the 2D model in Fig.2c), occurring every ~4.5 nm. A similar configuration is observed at another interface area equivalent with a 90° rotation of the structure in Fig.2b. These results unveil the 3D configuration of the interface, which has a grid of convex protrusions of spinel into periclase with misfit dislocation at each minimum (Fig2d). The structure reveals the mechanism of the interface migration: the climb of the misfit dislocations is the rate-limiting factor and therefore leads to this scalloped geometry. Furthermore, the extra atoms required for dislocation climb leave behind vacancies that eventually form pores at the interface, which provides additional resistance to interface motion and leads to doming of the interface on the scale of individual grains. These results also show that a fundamental understanding of the interface reaction and migration on the atomic scale is the key for understanding the interface migration on the larger scale.

 Reference:

[1] GÖTZE, L. C. et al., Phys. Chem. Minerals, 41, 681-693 (2014)

[2] Jeřábek, P. et al., American Journal of Science, 314, 940-965 (2014)

[3] C. Li et al., submitted (2016)

[4] This research was funded by the EU’s Horizon 2020 Marie Curie grants No. 656378–Interfacial Reactions (CL) and the Austrian Science Fund (FWF): I1704-N19 in the framework of FOR741-DACH (GH). 

Chen LI (Vienna, AUSTRIA), Thomas GRIFFITHS, Timothy J. PENNYCOOK, Clemens MANGLER, Lutz C. GÖTZE , Petr JEŘÁBEK , Jannik MEYER, Gerlinde HABLER, Abart RAINER
11:30 - 11:45 #6178 - MS08-OP310 3D analytical investigation of melting at lower mantle conditions in the laser-heated diamond anvil cell.
3D analytical investigation of melting at lower mantle conditions in the laser-heated diamond anvil cell.

Diamond anvil cell (DAC) is a unique tool to study materials under static high pressures up to several hundreds of GPa comparable to the pressures in the earth and planets interior. By using laser heating the temperature of the material inside the cell can be raised to several thousand degrees. This allows us to reach to the pressure and temperature conditions of deep mantle in laser heated diamond anvil cell (LHDAC). On the other hand small heated volume of the sample adjacent to the high thermally conductive diamonds results in large temperature and pressure gradients which affect the phase transformation and chemical distribution in LHDAC.

To fully understand the phase assemblages and equilibrium inside the LHDAC, it is essential to use three dimensional analytical characterization methods. As a proxy to deep mantle composition, San Carlos olivine has been chosen as a starting material for this study. To observe the effect of pressure and heating time, five samples are prepared. Three samples were melted at ~3000 K and at 45 GPa for durations of 1, 3 and 6 minutes. Other two samples were melted for 3 minutes at 30 GPa and 71 GPa. Each sample was then sliced by focused ion beam (FIB) with slice thickness of 50-100 nm. A secondary electron image and an energy dispersive x-ray (EDX) map were acquired from each slice by scanning electron microscope (SEM) in a dual beam FIB instrument. Half of the heated area in each sample was used for 3D FIB tomography and the other half is used to extract a 100 nm thick thin section for subsequent analysis by analytical transmission electron microscope (TEM). TEM is used to obtain accurate EDX maps from the phases. Also, the structure of crystalline phases has been characterized by electron diffraction technique.

3D reconstruction of SEM EDX maps (figure 1) shows that the heated area is roughly spherical and it consists of three main regions in all samples which correspond to ferropericlase (Mg­­, Fe)O (Fp), perovskite-structured bridgmanite (Mg,Fe)SiO3 (Brg) and iron-rich core. The bulk of the heated area is surrounded by ferropericlase shell. Then, we find a thick region of bridgmanite phase just inside the Fp shell and in the center lies an iron-rich core. In addition, in 45 GPa sample heated for 3 minutes we start to see another (Mg, Fe)O phase (Mw) around the core which is more iron-rich than the Fp shell. In the 45 GPa sample heated for 6 minutes this iron-rich oxide (Mw) entirely surrounds the iron-rich core. TEM analysis shows a third and even more iron-rich (Mg, Fe)O phase forming a thin layer (~70nm) between the Mw and the core. The core is getting richer in iron by increasing the pressure or heating time and its structure varies among the samples. For instance, in 45 GPa sample heated for 1 minute the core has eutectoid structure with iron nanoparticles distributed in it (figure 2) while in the 45 GPa sample heated for 6 minute we have a granular structure with the higher content of iron in the center of grains (figure 3). Moreover, we can see narrow Fp veins connecting the Fp shell to the iron-rich core in all of the samples, particularly in 71 GPa sample these veins are numerous and thick. In fact, they occupy a substantial part of Brg region in this sample.

Farhang NABIEI (Lausanne, SWITZERLAND), Marco CANTONI, James BADRO, Susannah DORFMAN, Richard GAAL, Hélène PIET, Philippe GILLET
11:45 - 12:00 #6575 - MS08-OP311 Alteration mechanisms of limestone used in built cultural heritage : use of isotoping labelling to determine the water penetration and reaction sitesIn France 52% of the historical monuments are made out of limestone; the preservation of this material is.
Alteration mechanisms of limestone used in built cultural heritage : use of isotoping labelling to determine the water penetration and reaction sitesIn France 52% of the historical monuments are made out of limestone; the preservation of this material is.

In France 52% of the historical monuments are made out of limestone; the preservation of this material is therefore an economic, scientific and cultural challenge. In urban area, limestones used in the façade of the buildings are exposed to a polluted environment and their degradation is already well documented. Several alteration processes are expected to occur such as phases precipitation and dissolution. In such environment, the most common alteration secondary phase is gypsum (CaSO4, 2H2O) formed from the reaction between calcareous stones, environmental water and sulfuric acid from the atmosphere.

Whatever the alteration process, water as rainfall (wet deposition) or as vapor state (dry deposition) is the alteration agent, so that it is the main parameter to focus on .Thus in order to better understand the stone/water interaction and to propose solutions for preserving the built cultural heritage, we developed an original methodology based on water isotopic tracers (D and 18O). Deuterium was used to localize water penetration front in the material, while 18O enabled to determine secondary phase reaction sites, mainly composed of gypsum.

Pristine samples from quarry and samples from Parisian monuments were selected to compare different alteration stages. Firstly, their main chemical and physical properties linked to the alteration were studied based on a multiscale characterization.(Saheb et al., 2015). Then, samples were altered in laboratory by realistic and controlled conditions of dry deposition during 2 months using isotopically labeled water. The reaction zones were analyzed by nano-SIMS. This experimentation enabled determining that water entirely penetrated in samples from quarry and from monuments, what highlights that the alteration layer does not seem to have a protection effect. In surface or deeper inside the sample, 18O enrichment highlights preferential reaction sites, localized in micro-cracks inside the gypsum zones and along grains of calcite (Figures 1 and 2).

This innovative methodology is a first step to understand the alteration mechanism formation on limestone used in the façade of the buildings. Understanding the mechanisms and especially the role of the alteration layer will contribute to improve the knowledge of stone chemical alteration processes to develop appropriate conservation strategies for the buildings.

References:

M. Saheb, J.D. Mertz, E. Colas, O. Rozenbaum, A. Chabas, A. Michelin, A. Verney-Carron, J.P. Sizun, Multiscale characterization of limestone used on monuments of cultural heritage, MRS Proceedings (2014), p. 1656. 

Adam DRICI (Creteil), Mandana SAHEB, Jean-Didier MERTZ, Aurélie VERNEY-CARRON, Loryelle SESSEGOLO, Laurent REMUSAT, Adriana GONZALEZ-CANO
12:00 - 12:15 #5749 - MS08-OP307 Correlative nonlinear optical microscopy and infrared nanoscopy reveals collagen degradation in altered parchments.
Correlative nonlinear optical microscopy and infrared nanoscopy reveals collagen degradation in altered parchments.

Non-invasive investigation techniques are strongly needed to avoid sampling during the examination of heritage artefacts. This study aims at developing specific and non-destructive mapping of the degradation of parchment, which is mainly composed of dermal fibrillar collagen. The main issue is to characterize gelatinization, its ultimate and irreversible alteration corresponding to collagen denaturation to gelatin, both from the morphological and chemical point of view. To that end, we implement correlative imaging of parchment using nonlinear optical (NLO) microscopy and nanoscale infrared spectroscopy (nanoIR).

NLO microscopy, also called multiphoton microscopy, advantageously provides non-invasive three-dimensional (3D) multimodal imaging of scattering samples with micrometer-scale resolution [1]. Among the collected signals, SHG signals are specific for dense non-centrosymmetric materials.  These signals therefore provide a unique structural probe of fibrillar collagen at the micrometer scale since the small signals from the collagen triple helices at molecular scale are amplified by constructive interferences at macromolecular scale due to the tight alignment of the collagen molecules to form collagen fibrils [2]. Nanoscale infrared spectroscopy is carried out thanks to an Atomic Force Microscope (AFM) coupled with an IR pulsed tunable laser [3]. This IR nanoscopy allows acquiring chemical mapping and local IR spectra to characterize and image samples at nanoscale.

Using this multiscale approach, key information about collagen and gelatin signatures is obtained in parchments and assessed by characterizing the denaturation of pure collagen reference samples. A new absorbing band is observed near the amide I band in the IR spectra, colocalized with the onset of fluorescence signals in NLO images. Meanwhile, a strong decrease is observed in Second Harmonic signals (see figure 1).

NLO microscopy therefore appears as a powerful tool to reveal collagen degradation in a non-invasive way. It should provide a relevant method to assess or monitor the condition of collagen-based materials in museum and archival collections and opens avenues for a broad range of applications regarding this widespread biological material.

[1] G. Latour, J.-P. Echard, M. Didier, and M.-C. Schanne-Klein, Opt. Express 20, 24623–24635 (2012).

[2] S. Bancelin, C. Aimé, I. Gusachenko, L. Kowalczuk, G. Latour, T. Coradin, and M.-C. Schanne-Klein, Nat. Commun. 5 (2014).

[3] A. Dazzi, C.B. Prater, Q. Hu, D.B. Chase, J.F. Rabolt, C. Marcott, Appl. Spectrosc. 66(12), 1365 (2012).

Gael LATOUR, Laurianne ROBINET, Alexandre DAZZI, François PORTIER, Ariane DENISET-BESSEAU, Marie-Claire SCHANNE-KLEIN (PALAISEAU CEDEX)
12:15 - 12:45 #8278 - MS08-S88 Watching works of art under the synchrotron lights to reveal their secrets.
Watching works of art under the synchrotron lights to reveal their secrets.

As early as he discovered X-rays in 1895, Roentgen envisaged their possible application for the study of works of art, and of paintings in particular. More than one century latter, X-rays are routinely used for the analysis of artworks. Beyond X-ray radiography (which application to paintings was patented already in 1914 (Bridgman, 1964)), many other X-ray based techniques can be used, such as X-ray fluorescence and X-ray diffraction, giving access to the chemical composition of ancient and artistic objects. Efforts are made to extend the use of such laboratory techniques in two main directions: on the one hand with the development of portable instruments, allowing on-site (museums, archaeological sites, historical buildings) analyses; on the other hand, exploiting the additional capabilities offered by synchrotron sources.

The brightness, collimation, polarization, emission spectrum and partial coherence of synchrotron beams offers: i) the possibility to focus the beam to few microns, down to few tens of nanometers, ii) improved detection limits, iii) access to the element speciation thanks to spectroscopy techniques, iv) reduced acquisition times, allowing the acquisition of many points, and in particular 2D and 3D scanning maps, v) additional contrast modes such as phase contrast imaging. Considering the technical constraints associated to the intrinsic properties of artistic materials (precious objects and samples of very limited size when available; highly heterogeneous at multi-scales; composed of a variety of materials (in/organic, un/crystallized, low/high Z elements)), these combined analytical strengths are highly beneficial (Bertrand, et al., 2012).

Here, we will present a set of recent examples of analysis of artistic and ancient materials, performed at the ESRF. Experiments usually aim at understanding how works of art were created (choice and synthesis of ingredients, in particular of pigments, firing temperature in the case of ceramics, etc). As an example, Figure1 shows a combined micro X-ray fluorescence, micro X-ray diffraction, micro X-ray absorption spectroscopy analysis of blue pigments in Chinese Qinghua porcelains. Fragments were sampled from sherds and prepared as thin sections (Top left). µXRF maps show particular concentration of Co and Ca in the pigment regions, and of iron on the surface (top right). Full-field XANES at the Co K-edge reveals the presence of two main Co species, in the pigment and in the glaze (bottom left: (a) transmission image recorded at 7670 eV, (b) the edge jump map, (c)  cluster maps and average XANES spectra obtained by PCA, (d) Speciation maps obtained from the least squares linear combination fitting (standards CoAl2O4 and Co in glaze)). Micro X-ray diffraction offers further determination of Co pigment lattice parameters (bottom right) (Wang, et al.).

Other experiments focus on degradation issues, and aim at identifying degradation products, and internal and external factors responsible for such instabilities. Figure 2 shows a typical example with the study of cinnabar darkening in Madonna with Child, St. Sebastian, St. John the Baptist and two donors, Boltraffio, Louvre (© C2RMF) (Cotte & Susini, 2009).

 

Acknowledgments

All the users whose work will be reported are highly acknowledged; all ESRF staff involved in the development and maintenance of instruments are thanked as well.

Reference

Bertrand, L., Cotte, M., Stampanoni, M., Thoury, M., Marone, F. &  Schöder, S. 2012. Development and trends in synchrotron studies of ancient and historical materials. Physics Reports, 519(2):51-96.

Bridgman, C. F. 1964. The amazing patent on the Radiography of Paintings. Studies in Conservation, 9(4):135-39.

Cotte, M. &  Susini, J. 2009. Watching Ancient Paintings through Synchrotron-Based X-Ray Microscopes. Mrs Bulletin, 34(6):403-05.

Wang, T., Zhu, T. Q., Feng, Z. Y., Fayard, B., Pouyet, E., Cotte, M., De Nolf, W., Salomé, M. &  Sciau, P. Synchrotron radiation-based multi-analytical approach for studying underglaze color: The microstructure of Chinese Qinghua blue decors (Ming dynasty). Analytica Chimica Acta.

Marine COTTE (GRENOBLE CEDEX 9)

10:30-12:45
Added to your list of favorites
Deleted from your list of favorites

LS3-I
LS3: Cell functional exploration
SLOT I

LS3: Cell functional exploration
SLOT I

Chairmen: Judith KLUMPERMAN (Utrecht, THE NETHERLANDS), Franck RIQUET (Lille, FRANCE)
10:30 - 11:00 #8648 - LS03-S10 mScarlet, a novel high quantum yield (71%) monomeric red fluorescent protein with enhanced properties for FRET- and super resolution microscopy.
mScarlet, a novel high quantum yield (71%) monomeric red fluorescent protein with enhanced properties for FRET- and super resolution microscopy.

mScarlet, a novel red fluorescent protein was generated from a synthetic template based on a consensus amino acid sequence derived from naturally occurring red fluorescent proteins and purple chromoproteins and on consensus monomerization mutations. The encoded synthetic red fluorescent protein was optimized by molecular evolution through site directed and random mutagenesis. Improved variants were selected by quantitative multimode screening for increased fluorescence lifetime, increased photo stability, increased quantum yield and for increased chromophore maturation.

Very bright variants were obtained with high fluorescence lifetimes up to 3.8 ns, quantum yields >70 % and complete maturation. The monomeric status of the variants was confirmed by OSER analysis and with a-tubulin fusions. The brightness of mScarlet is >2- fold increased as compared to bright red fluorescent proteins such as mCherry, mRuby2 and tagRFP-T as was analyzed with quantitative (single plasmid with viral 2A sequence) coexpression with mTurquoise2 in mammalian cells.

mScarlet can be used as a bright red fluorescent fusion tag for staining various subcellular structures in live cells. Because of their efficient maturation and high quantum yield, mScarlet vastly outperforms existing monomeric red fluorescent proteins in ratiometric FRET-microscopy applications due to seriously enhanced sensitized emission and lack of photochromism (figure 1).

During evolution mScarlet variants with substantially altered spectroscopic properties were generated including fluorescence lifetime variants, photo labile variants and strongly spectrally shifted variants. Besides looking at generally optimized properties such as increased lifetime, maturation and brightness, some of these variants were rescreened for blinking properties. An mScarlet variant (7Q2BMs-K) was identified that produced high spontaneous blinking upon illumination at 561 nm. This variant was fused to life-act and co-expressed with mVenus-Lifeact. With 488 and 561 nm excitation, dual life cell single molecule localization microscopy produced perfectly colocalized yellow and red actin structures in living cells (figure 2), demonstrating the usefulness of the novel red fluorescent proteins for life cell super-resolution microscopy.

Lindsay HAARBOSCH, Daphne BINDELS, Marten POSTMA, Mark HINK, Antoine ROYANT, Theodorus GADELLA (Amsterdam, THE NETHERLANDS)
Invited
11:00 - 11:30 #8437 - LS03-S11 Functionalized Carbon Nano-onions as Imaging Probes for Cancer Cells.
Functionalized Carbon Nano-onions as Imaging Probes for Cancer Cells.

Multimodal imaging probes based on carbon nano-onions (CNOs) have emerged as a platform for bioimaging because of their cell-penetration properties and minimal systemic toxicity. [1-3] We have developed a synthetic multi-functionalisation strategy for the introduction of different functionalities (receptor targeting unit and imaging unit) onto the surface of the CNOs. The modified CNOs display high brightness and photostability in aqueous solutions and their selective and rapid uptake in two different cancer cell lines without significant cytotoxicity is demonstrated. The localization of the functionalized CNOs in late-endosomes cell compartments is revealed by a correlative approach with confocal and transmission electron microscopy. [4] Understanding the biological response of functionalized CNOs with the capability to target cancer cells and localize the nanoparticles in the cellular environment, will pave the way for the development of a new generation of imaging probes for future biomedical studies.

[1] Yang, M. et al. Small 2013, 9, 4194. [2] Bartelmess, J. et al. RSC Adv.2015, 5, 50253–50258. [3] Marchesano, V. et al. Nanomaterials 2015, 5, 1331. [4] Frasconi, M. et al. Chem. Eur J. 2015 , 21, 1971.

Silvia GIORDANI (Genova, ITALY), Marco FRASCONI, Roberto MAROTTA
Invited
11:30 - 12:00 #7883 - LS03-S12 Nanotomy and CLEM techniques shed new light on biomedical processes.
Nanotomy and CLEM techniques shed new light on biomedical processes.

A spectrum of dyes and probes now enable to localize molecules of interest within living cells by fluorescence microscopy. With electron microscopy (EM), cellular ultrastructure has been revealed. Bridging these two modalities, correlated light microscopy and EM (CLEM) opens new avenues [1].

The first focus will be on recent developed labeling strategies for molecules that allow CLEM (Fig.1). These include particles and substrates to highlight endogenous proteins that are targeted using affinity, but also genetically-encoded probes [2], and traditional stains for light microscopy that aid in EM-analysis of samples. Probes that can only be detected in a single modality, and require image overlay, as well as combinatorial probes that can be visualized both at LM and EM levels will be discussed.

In addition, new approaches for large scale EM (“nanotomy”),  either TEM-based [3] or S(T)EM-based [4,5], to visualize macromolecules and organelles in the context of organized cell systems and tissues will be highlighted. Matching the areas of acquisition in CLEM and EM will not only increase understanding of the molecules in the context, but also is a straight forward manner to combine the LM and EM image data. While these new developments aid to better understand the contribution of molecules, organelles and cells in tissue-function, the amount of data is huge and quantification and sharing data need new solutions. We will show how millions of particles can be recognized and quantified using Fiji. Moreover, our open-access data-sharing (Fig.2) via nanotomy.org will be highlighted to easily access the ultrastructure in a variety of research projects, including blistering diseases that affect skin and mucosa in men and Type 1 diabetes.

Covering a variety of probes and approaches for image overlay will help to enable (new) users to broadly implement CLEM and/or nanotomy to better understand how molecules (mal)function in biology [6].

 

References:

 

[1] De Boer et al. (2015) CLEM: Ultrastructure lights up! Nature Methods 12:503

[2] Kuipers et al. (2015) FLIPPER for CLEM. Cell & Tissue Research 360:61

[3] Ravelli et al. (2013) Destruction (…) in T1D rats at macromolecular resolution. Sci. Reports 3:1804

[4] Sokol et al. (2015) Nanotomy of blistering diseases. J. Investigative Dermatology 135:1763

[5] Kuipers et al. (2015) SEM-based immunolabelling and nanotomy. Exp. Cell Res. 337:202

[6] Sponsored by ZonMW91111.006; NWO175-010-2009-023 ; STW Microscopy Valley 12718

Ben GIEPMANS (AV Groningen, THE NETHERLANDS)
Invited
12:00 - 12:15 #4565 - LS03-OP018 Monitoring Cell Death in Real-Time/Time-Lapse Studies.
Monitoring Cell Death in Real-Time/Time-Lapse Studies.

     Personalised medicine, chronic exposure drug toxicity and environmental contaminants amongst others have created the demand for in vitro assays which are more physiologically relevant. One aspect of this may be to run assays for many days and monitor them in real-time or time-lapse mode. Therein, it can be useful to assess the general cell survival or perhaps a specific cell death pathway in the context of the treatment or insult under observation.

     Depending on the sample under analysis, to date, this can be done by either marking the cells that remain functionally competent or measure release of ATP at a bulk level. It would be preferable if, conversely, only dead/damaged/apoptotic cells were marked in a binary manner, with a convenient and spectrally separated emission signature and with specificity for a predictable intracellular target such as gDNA, and cell-by-cell.

     To explore this need a novel far-red DNA binding viability probe, DRAQ7, has been developed. It has been shown to have undetectable toxicity in long-term / real-time cell based assays as validated in recent publications (Akagi et al., 2013, Marciniak et al., 2013, Wang et al., 2015, Liang et al., 2015, and internal data) including ultra-sensitive bioassays for DNA intercalators, ex-plant tissue culture and nano-particle toxicity, time-lapse apoptosis assays and importantly in the presence of toxicants / anti-cancer compounds. Thus, cells can be exposed to it at any stage of an assay to permit realtime monitoring of loss of membrane integrity (in apoptosis, death). Being DNA specific it allows monitoring cell-by-cell while its spectral properties mean it can be incorporated into multi-colour flow cytometry experiments, or with Hoechst 33342 or CyTRAK Orange (for simple cell health assays) or with mitochondrial membrane potential probes such as TMRM.

     DRAQ7 is truly cell impermeant yet retains the DNA targeting and far-red fluorescence of the parent DRAQ chromophore. Accordingly, it likewise can be used on HCS imaging platforms, fluorescence microscopes, cytometers and sorters. Its spectral properties (long wavelength) are particularly suited to penetrative imaging of multicellular structures and thick ex-plant tissue sections.

     The presentation will focus on the deployment and application of DRAQ7 to demanding real-time and 3D micro-tissue assays including patient-derived samples that enable strategies for personalised medicine, underpinned by core performance data that define its unique properties as a cross-platform imaging probe for cell viability.

Roy EDWARD (Shepshed, UK)
12:15 - 12:30 #5901 - LS03-OP020 High-Resolution Phosphorescence Lifetime Imaging of Oxygen in 3D Tissue Models.
High-Resolution Phosphorescence Lifetime Imaging of Oxygen in 3D Tissue Models.

Molecular oxygen (O2) plays a multitude of important roles in cell and tissue function and (patho)physiology. Real-time quantitative imaging of O2 by phosphorescence quenching method enables detailed mechanistic studies of cell and tissue physiology, responses to hypoxia, drug treatment and other stimuli. We applied high-resolution Phosphorescence Lifetime Imaging Microscopy (PLIM) to study several different 3D tissue models: multi-cellular spheroids, excised animal tissue and cultured organoids. To achieve efficient, stable passive staining of the cells and tissues, we have designed a panel of cell-penetrating phosphorescent nanosensors with variable surface charge and spectral characteristics, and a small molecule O2 probe, Pt-Glc, which provides fast and in-depth staining of most cell models. Using Pt-Glc and PLIM we studied cultures of PC12 (rat pheochromocytoma) and HCT116 (human colon cancer) cell spheroids, intestinal organoids and ex vivo brain, colon, bladder and vessel tissue samples. We showed that cell aggregates of >50 um size are significantly deoxygenated under ambient O2, but remain viable and respond to treatment with compounds affecting metabolism. For the neurospheres from embryonic rat brain, standard protocol was developed for O2 imaging by PLIM, multiplexed with immunofluorescence of cell type, proliferation and death markers. In giant umbrella cells of mouse bladder epithelium we observed marked intracellular gradients of O2 of up to 40-50 mM across the cell or 0.6 mM/mm, which may play physiological roles in tissue function. We found decreased respiration in the colonocytes in dextran sulfate sodium (DSS)-induced colitis mouse colon tissue. These results demonstrate the utility of cell-penetrating O2 probes and PLIM method for life science research. 

This work was supported by Science Foundation Ireland (SFI) grants 12/RC/2276 and 13/SIRG/2144.

Dmitri PAPKOVSKY (Cork, IRELAND), Alexander ZHDANOV, Ruslan DMITRIEV
12:30 - 12:45 #5802 - LS03-OP019 Multi-modal in vivo visualization of single cell dynamics by 1P, 2P, light-sheet, and on-chip technologies.
Multi-modal in vivo visualization of single cell dynamics by 1P, 2P, light-sheet, and on-chip technologies.

We made fluorescence multi-scale imager including 1P, 2P, light-sheet, microscope, and on-chip sensors for visualization of XYZT single cell dynamics in vivo. We integrated five microscope for systematic evaluation of living animals.

First visualization system is super-resolution imaging based on non-linear optics, with X- (resonance), Y-(galvano), and Z-(piezo) axis scanning. Real-time, multi-color XYZT multi-photon imaging enabled us to visualize single blood cell behavior in vessels and stroma. Spatial and time resolution was improved by pattern illumination. We analyzed thrombus formation, and inflammatory responses under microscope. We developed thrombus formation animal models, and elucidated cellular mechanisms of cardiovascular diseases. We directly manipulated cells by photo-chemical reactions, and two photon lasers to induce and observe thrombotic reactions. Using this system, we elucidated contribution of endothelial injuries to thrombus formation.

Second, high-resolutions and broad-imaging field was simultaneously enabled using 8K CMOS technologies, and 1P spinning disk confocal. We visualized whole organs and single cell in one image, and revealed complicated cell-cell interactions networks in single view. 8K, 60fps, and multi-color imaging visualized single platelet dynamics and tissue structural changes in single image.

Third, we performed light-sheet imaging for living mice, and enabled high-speed observation of living animals.

Fourth, macro imaging system for awake mice, pigs and humans was also developed, and free behavior monitoring revealed the tight association between metabolism and vascular reactions during daily life. Bioluminescent and fluorescent imaging from body surface using CMOS camera, image intensifier, and macro-lens enabled us to visualize cellular dynamics without anesthesia.

Last, wearable, implantable, and minimized devices for non-invasive recording were also developed using lens-less and on-chip technologies. We utilized SCMOS, micro-lens array, and LED illumination technologies.

 

In sum, we developed multi-scale imaging system which can evaluate cellular networks dysregulations in diseased conditions.

Satoshi NISHIMURA (TOCHIGI, JAPAN)

13:45
13:45-15:45
Added to your list of favorites
Deleted from your list of favorites

MS1-III
MS1: Structural materials, defects and phase transformations
SLOT III

MS1: Structural materials, defects and phase transformations
SLOT III

Chairmen: Patricia DONNADIEU (ST MARTIN D'HERES CEDEX, FRANCE), Randi HOLMESTAD (Trondheim, NORWAY), Simon RINGER (Sydney, AUSTRALIA)
13:45 - 14:15 #8770 - MS01-S68 ACOM-TEM and its application for the investigation of deformation pathways in nanocrystalline Pd and AuPd.
ACOM-TEM and its application for the investigation of deformation pathways in nanocrystalline Pd and AuPd.

Most of our current understanding of the deformation mechanisms active in nanocrystalline (nc) metals stems from in-situ deformation experiments on bulk materials using x-ray diffraction (XRD). However, XRD cannot directly resolve the local deformation processes, e.g. grain growth or twinning. For a local analysis, these processes are traditionally investigated using BF/DF-TEM. However, varying contrast due to local orientation changes, bending and defects during in-situ BF-TEM straining experiments make an accurate interpretation for nanometer sized grains difficult. On the other hand, Automated Crystal Orientation Mapping (ACOM-TEM) allows for the identification of the crystallographic orientation of all crystallites with sizes down to around 10 nm, well below the limit of electron back scatter diffraction (EBSD)1. Using template matching to reveal the crystal orientation, the ASTAR (Nanomegas) ACOM-TEM analysis offers an angular resolution that is nearly as good as EBSD1,2.

Recently, ACOM-TEM imaging in STEM modus was combined with in-situ straining inside a TEM3–5. This combination was the key to new data evaluation based on orientation maps. By tracking individual crystallites through a straining series the change of their orientation can be evaluated in order to distinguish between an overall crystallite rotation and sample bending. In addition, twinning/detwinning and grain growth can be directly followed and the automatic data evaluation leads to user independent quantitative statistical information such as grain size distribution3.

Recent investigations revealed some challenges using ACOM-TEM for in-situ experiments if there are overlapping crystallites. Overlapping crystallites lead to superimposed diffraction patterns that confuse the matching procedure. Tilting nc Pd in-situ and tracking the changes using ACOM-TEM, it became apparent that some orientations are more dominant than others during the matching procedure. Further, we present an ambiguity filter that reduces the number of pixels with a '180° ambiguity problem' (Fig. 1). The challenges discussed here for orientation mapping of nc materials do not only appear with ACOM-TEM, but are mostly an inherent problem of any TEM projection technique. However, using ACOM-TEM these limitations become apparent and can be properly analyzed, e.g. by mapping the rotation of many crystallites for a given area of interest. This enables to detect sample bending or tilting in a (in-situ) series of orientation maps, which cannot be measured by BF/DF-TEM.

The ACOM-TEM measurement and evaluation routine was applied to magnetron sputtered PdxAu1-x thin film samples of about 50 nm. Grain growth and grain fragmentation as well as twinning and detwinning have been observed to take place simultaneously at different locations. In addition, large angle grain rotations with ~39° and ~60° occur that can be related to twin formation, twin migration and twin-twin interaction as a result of partial dislocation activity (Fig. 2). Furthermore, plastic deformation in nanocrystalline thin films was found to be partially reversible upon rupture of the film. In conclusion, conventional deformation mechanisms are still active in nanocrystalline metals, however, with different weighting than in conventional materials with coarser grains.

We would like to acknowledge Christian Brandl, Edgar Rauch, Florian Bachmann, Ralf Hielscher, Ankush Kashiwar. This work was supported by the DFG under grant FOR714 and Karlsruhe Nano Micro Facility (KNMF).

References

(1)            Rauch, E. F.; Portillo, J.; Nicolopoulos, S.; Bultreys, D.; Rouvimov, S.; Moeck, P. Zeitschrift für Krist. 2010, 225 (2-3), 103–109.
(2)            Zaefferer, S. Cryst. Res. Technol. 2011, 46 (6), 607–628.
(3)            Kobler, A.; Kashiwar, A.; Hahn, H.; Kübel, C. Ultramicroscopy 2013, 128, 68–81.
(4)            Kobler, A.; Kübel, C. Imaging Microsc. 2014, No. 1.
(5)            Mompiou, F.; Legros, M. Scr. Mater. 2015, 99, 5–8.

Aaron KOBLER (Eggenstein-Leopoldshafen, GERMANY), Christian KÜBEL, Horst HAHN
Invited
14:15 - 14:30 #4875 - MS01-OP199 Measuring Charge Distribution in Nanoscale Magnesium Aluminate Spinel by Electron Energy-Loss Spectroscopy and Electron Holography.
Measuring Charge Distribution in Nanoscale Magnesium Aluminate Spinel by Electron Energy-Loss Spectroscopy and Electron Holography.

Charge distribution resulting in the formation of a space charge zone (SCZ) in ionic materials has a critical role on functional properties [1].  Even though significant advances in theoretical models have been accomplished, experimental evidence in nanoscale granular materials is indirect.

Here, we investigated the distribution of cations and defects on the formation of a SCZ in a nanoscale granular model system of non-stoichiometric MgO∙nAl2O3 (MAS, n= 0.95 and 1.07). The SCZ was investigated experimentally by electron energy-loss spectroscopy (EELS) and off-axis electron holography (OAEH).

EEL spectra were collected along directions perpendicular to grain boundaries (GB’s), from which the magnesium-to-aluminum relative cation concentrations were calculated, as presented in Fig.1. We found that regardless of annealing processes, the vicinity of GB’s of the Mg rich spinel has excess Mg+2 cations while the vicinity of GB’s of the Al rich spinel has excess of Al+3 cations. Additionally, the cation distribution shows strong dependency on the grain size. For non-stoichiometric MAS, cation concentration is proportional to the defect concentration, because deviation from stoichiometry results in adjacent defects that compensate for the electric charge [2, 3, 4]. In both materials, the cation distribution is inhomogeneous for grains smaller than 40 nm. For larger grains, the defect concentration approaches the bulk value at the center of the grain. Furthermore, excess of Mg (Al) cations at the vicinity of the GB decreased with increase of grain size. Maier et al. [1] calculated that for grain size at the scale of the Debye length (estimated at 9nm for non-stoichiometric MAS studied here [7]), the GC is no longer electrically neutral, instead influenced by accumulation or depletion of charge at the boundaries.

Due to the lack of accurate values for defect formation energy [5, 6], we applied OAEH to measure directly the electrostatic charge distribution in nano-sized MAS. We show that charge distribution and the buildup of electrostatic potential between GB and core are linked to the spatial distribution of defects rather than the overall composition of MAS (Fig. 2). At the vicinity of GB’s, excess Mg+2 or Al+3 cations accumulate depending on the composition, the magnitude of which increases with decreasing grain size. Indeed, the potential distributions show the relation between the excess cation species, grain size and the Debye length, in agreement with theoretical models [1].

References:

[1] J. Maier, Prog. Solid State Chem. 23, 171-263 (1995).

[2] M. Rubat du Merac et al., J Am Ceram Soc. 96 (2013) 3341-3365.

[3] S.T. Murphy et al, Philosophical Magazine. 90 (2010) 1297-1305.

[4] Y. Chiang, W.D. Kingery, J Am Ceram Soc. 73 (1990) 1153-1158.

[5] K. Lehovec, J. Chem. Phys. 21, 1123-1128 (1953).

[6] K. Kliewer & J. Koehler,. Phys. Rev. 140, A1226 (1965).

[7] M. Halabi, V. Ezertzky, A. Kohn and S. Hayun, submitted.

Mahdi HALABI (Beer-Sheva, ISRAEL), Amit KOHN, Shmuel HAYUN
14:30 - 14:45 #6841 - MS01-OP214 Improved Quantitative Compositional Analysis of γ’ and γ’’ in Additively Manufactured Alloy 718 Using STEM X-ray Energy Dispersive Spectrometry.
Improved Quantitative Compositional Analysis of γ’ and γ’’ in Additively Manufactured Alloy 718 Using STEM X-ray Energy Dispersive Spectrometry.

Selective laser melting (SLM) is an additive manufacturing technique where successive laser beam passes are used to melt metal powder which forms a solid layer on solidification with high densification, little material waste, and large design freedom [1]. The application of SLM to repair high temperature components that often need reconditioning requires an understanding of the microstructural and compositional developments of the chosen material throughout the SLM process. Alloy 718 is a Ni-Cr-Fe-Nb-Ti-Al alloy used in applications where high strength is needed while maintaining corrosion and creep resistance, making this alloy a prime candidate for SLM structural and compositional characterization. Precipitation hardening is one of the primary strengthening mechanisms of Alloy 718, where intermetallic phases of L12-ordered Ni3(Al, Ti) (γ’) or D200-ordered Ni3(Nb, Ti) (γ’’) may form coherent precipitate particles in the face-centered cubic matrix (γ) [2]. Additional phases that may be present in the microstructure of Alloy 718 include D0a-ordered Ni3Nb (δ), MC, M6C, M23C6, and (Ni, Cr, Fe)2(Nb, Mo, Ti) Laves [3, 4]. The complex microstructures in this alloy system are further complicated by the multiple heating and cooling cycles present in the SLM process, thus requiring characterization on the nanoscale in order to understand the microstructural development during processing.                           

Analytical electron microscopy allows the identification of the particular microstructural components on the micro and nano scales.  Alloy 718 is of particular interest in that the γ’ precipitates can nucleate on the (001) surface of γ’’ precipitates in the as-SLM condition. The structure and chemical composition of these precipitates was investigated through X-ray energy dispersive spectrometry (XEDS) using an aberration-corrected FEI Titan G2 ChemiSTEM equipped with the Super X EDX X-ray detector configuration. Figure 1 shows a γ’’ precipitate with γ’ precipitates on the two elongated sides of the γ’’ in both scanning transmission electron microscopy (STEM) bright-field (BF) and high-angle annular dark-field (HAADF) imaging modes. The understanding the formation of γ’/γ’’ requires both structural information about the interface and chemical analysis across the interface of the two precipitates. Figure 2 displays 4 XEDS spectrum images of Ni, Nb, Ti, and Al showing the location of these elements throughout the precipitates present in the γ matrix. Quantitative XEDS analysis was performed on an as printed Alloy 718 specimen where the γ matrix composition was found to be 50.5 wt % Ni, 1.4 wt % Nb, 0.3 wt % Al, and 0.09 wt % Ti with γ’ and γ’’ having compositions of 66.6 wt % Ni, 7.13 wt % Nb,  3.18 wt % Ti, and 2.4 wt % Al and 65.0 wt % Ni, 25.4 wt % Nb, 0.37 wt % Al and 3.6 wt % Ti, respectively.             

References:

[1] Gu, D.D., Meiners, W., Wissenbach, K., Poprawe, R., Int. Mater. Rev., 2012, 57, 133-164.

[2] Azadian, S., Wei, L.Y., and Warren, R., Mater. Charact., 2004, 53, 7.

[3] Burke, M.G. and Miller, M.K., J. de Physique, 1989, 50, C8 395-400.

[4] Rama, J.G.D., Reddya, A.V., Raob, K.P., Reddyc, G.M., Sundar, J.K.S., J. Mater. Process. Tech., 2005, 167, 73.

C. Austin WADE (Manchester, UK), Giacomo BERTALI, Thijs WITHAAR, David FOORD, Bert FREITAG, Grace BURKE
14:45 - 15:00 #6925 - MS01-OP216 Kinetic Behavior of Fe-Ni-C Martensitic Steels during Aging at Room Temperature.
Kinetic Behavior of Fe-Ni-C Martensitic Steels during Aging at Room Temperature.

The kinetic behavior of a Fe–24 Ni–0.4 C (weight percent) martensitic steel during aging at room temperature has been investigated by transmission electron microscopy. Electron diffraction, coupled with imaging techniques in a transmission electron microscope, including high resolution studies, have been used at various stages during the aging process, in order to characterize the microstructural features of the analyzed samples.

Rapid cooling (quenching) in liquid nitrogen of the austenitic state of the above-mentioned material leads to the formation of a martensitic phase which, at room temperature (RT), begins to transform, through a process called spinodal decomposition. As a result of this process, a modulated structure is formed, in which carbon-rich regions (precipitates) occur in a periodic manner throughout the matrix, leading to the presence of diffuse streaks or satellite spots around each fundamental (matrix) reflection on the electron diffraction patterns, see Figure1. This process is thus accompanied by a reduction of the tetragonality of the martensitic phase (bct), which evolves towards the formation of a cubic structure (bcc), see Figure2, corresponding to alpha-iron (ferrite). After long ageing times, Fe3C is observed. The presence and structure of intermediate carbides is also studied.  

The microstructure of the carbon-rich phase is also analyzed by electron diffraction. The evolution of the tetragonality with time, translated as the c/a ratio, which is also function of the carbon composition, is studied. Moreover, the variation of the distance between the carbon-rich precipitates (in fact, the periodicity of the modulated structure), along with the evolution of their width and length, are also observed.

 

Thanks are due to the Clym ( Centre Lyonnais de Microscopie) for the access to the TEM 2010F.

Sergiu CURELEA, Sophie CAZOTTES (VILLEURBANNE CEDEX), Thierry EPICIER, Frédéric DANOIX, Héléna ZAPOLSKY, Mikola LAVRSKYI, Philippe MAUGIS, Sara CHENTOUF, Mohamed GOUNE
15:00 - 15:15 #6327 - MS01-OP210 Influence of grain boundary character on void formation in nanotwinned copper.
Influence of grain boundary character on void formation in nanotwinned copper.

Most materials used in nuclear reactors are prone to radiation damage and their mechanical properties degrade with time, limiting their service life. Defects produced by high energy particle radiation can be highly mobile at high temperatures with their mobility being influenced by the local stress fields associated pre-existing defects and grain boundaries.  The relaxation of radiation defects can produce clusters and induce defect diffusion to interfaces and other pre-existing defects, where they can be absorbed and alter the microstructure that can detrimental to the material’s mechanical properties.  Radiation mechanisms that promote biased point defect diffusion can induce significant microstructural change and degrade properties. For instance, interstitials, being more mobile than the vacancies, are quickly absorbed by nearby dislocations, inducing creep by dislocation climb and dislocation multiplication that results in work hardening and embrittlement. Likewise, a small excess of remnant vacancies can agglomerate, leading to the formation of voids that cause swelling, increased residual stresses, the formation of microcracks, and the eventual failure of the material. The long-term stability of a microstructure under irradiation depends on its neutrality towards biased point defect dynamics. Materials with a high density of sites that can act as sinks for the point defects produced by high-energy particles would thus be an enabling technology for reliable and clean nuclear energy. 

 

We have explored radiation damage mechanisms in one such material, nanotwinned copper, which has a microstructure comprising a high density of coherent twin boundaries distributed in a “latter-like” morphology within columnar grains having high misorientation angle intercolumnar boundaries. By performing in-situ particle irradiation in a MeV TEM and characterizing the evolution in the microstructure using the Nanomegas ASTAR system, we have studied the irradiation induced damage and correlated the radiation induced grain boundary migration and void preferential void formation to the local grain boundary character and network.  In-situ observations were made at two different temperatures, room temperature and temperature above 573 K that stimulate void nucleation and growth. Figure 1 shows an example montage of the irradiation induced void formation and grain boundary migration with increased time and dose.  The radiation induced grain boundary migration (RIGM) of several high-angle random and low–angle grain boundaries within the irradiated zone was observed at both temperatures.  However, low-energy CSL type boundaries were not observed to migrate even at elevated temperatures as well as grain boundaries connected to junctions that contain one or more Σ3 boundaries. Thus, coherent twins are thought to stabilize the microstructure from radiation induced coarsening.

 

At temperature above 573 K during high-energy electron irradiations, copious amounts of voids were observed to nucleate and growth. They were predominantly observed to nucleate and grow in the regions with a high density of the coherent twin boundaries (CTBs).  Very few voids were observed around HARGBs and LAGBs as shown in Figure 2, and only a limited of number of small voids were observed in the vicinity of incoherent twin boundaries. Non-coordinated boundaries having excess free volume are efficient sinks for both vacancies and interstitials, and thus do not stimulate biased point defect diffusion that promote void formation, explaining the observed lack of voids in the vicinity of HARGBs. On the other hand, the observed high density of large voids in regions with finely spaced CTBs may arise from the biased diffusion of point defects. Interstitials being more mobile that are able to diffuse through CTBs can rapidly migrate and annihilate at free surfaces or other sinks such HARGBs. This leaves an excess of vacancies in the vicinity of the CTBs which can coalesce and cause the nucleation of voids.  Given these observations, a nanocrystalline material having a high fraction of CTBs, though being resistant to radiation induced coarsening, may not possess the ideal topology for radiation resilience and may be prone to increased void formation and swelling.

Thomas LAGRANGE (Lausanne, SWITZERLAND)
15:15 - 15:30 #5362 - MS01-568 HRTEM study of structural defects and related deformation mechanisms induced by nanocompression of silicon.
MS01-568 HRTEM study of structural defects and related deformation mechanisms induced by nanocompression of silicon.

Over the last years, progress in nanomaterials design and manufacturing has revolutionized technology and opened up prospects for many scientific researches. The investigations of material properties (optical, electronic, mechanical...) at small scales have revealed amazing behaviors, different from those currently observed in bulk samples. For instance, silicon, which is known to behave at room temperature as a brittle material, shows an unexpected ductile behavior when the sample size decreases below a few hundreds of nanometers [1]. The mechanisms leading to this phenomenon remain, however, poorly understood. In this context, this research project aims at investigating in more details the deformation behavior of silicon nanopillars by combining experimental techniques (SEM, FIB, HRTEM) and molecular dynamics simulations. In this work, various nanopillars with different orientations and diameters (from 100 nm to 1 µm), were patterned by Reactive-Ion Etching and FIB micromachining. These pillars were then compressed with a slow-strain-rate (10-4 s-1) at room temperature using a nanoindenter equipped with a flat punch and operated in displacement-control mode. Post mortem observations of deformed nanopillars performed by SEM and TEM reveal the activation of different slip systems. The comparison between experimental and simulated HRTEM images notably evidences the simultaneous propagation of partial and perfect dislocations in {111} planes. In addition, unexpected plastic events have also been observed in {113} planes. On the basis of the microscopic observations, various possible deformation mechanisms involved during the nano-compression of the pillars are proposed.

 

[1] F. Östlund et al., Brittle-to-Ductile Transition in Uniaxial Compression of Silicon Pillars at Room Temperature, Adv. Func. Mat., 19, p1(2009).

 

This work is performed within the framework of the ANR-funded research project « BrIttle-to-Ductile Transition in Silicon at Low dimensions » (ANR-12-BS04-0003-01, SIMI4 program).

Amina MERABET (Marseille), Michaël TEXIER, Christophe TROMAS, Marc VERDIER, Anne TALNEAU, Olivier THOMAS, Julien GODET
15:30 - 15:45 #5906 - MS01-OP211 Strain relaxation defects in Ge crystals grown on Si pillars.
Strain relaxation defects in Ge crystals grown on Si pillars.

Due to the differences in lattice parameters and thermal expansion coefficients, Ge films grown on Si substrates suffer from a high density of threading dislocations, cracks and wafer bowing. One method to avoid these problems is to grow Ge crystals on (001)-Si pillars by low energy plasma enhanced chemical vapor deposition (LEPECVD). With this strategy, threading dislocations of 60° type with Burgers vector (b) b=1/2 can be avoided in the bulk as they end at the sidewalls of the Ge crystals. However, the misfit strain is not totally relaxed; it is accommodated by misfit dislocations (MDs) of 60° and 90° types, while planar defects have not been reported in these Ge crystals.

This work presents a detailed analysis of defects formed to relax the misfit strain between the Ge crystals and Si pillars using high angle annular dark-field scanning transmission electron microscopy (HAADF-STEM).

In Ge/Si interface, the misfit strain is accommodated by 60° and 90° MDs. The 90° MDs are Lomer MDs lying on (001) planes which are formed by interaction of two 60° MDs. These pairs of MDs form steps at the interface leading to an atomic roughness.

Most interestingly, besides the MDs, there is a high density of planar defects which has not been reported before. These 2D defects are formed at the Ge/Si interface and extend between 3 to 40 nm along the {111} planes. We observed coherent twin boundaries (CTB), incoherent twin boundaries (ITBs) and stacking faults (SF) bounded by partial dislocations (PDs). Figure 1a shows CTBs of the ∑3{111}-type (red arrows), the distance between them is 2.95 nm. The inset of Fig. 1a is its Fourier transform (FFT), the twin plane correspond to GA(1-1-1)=GB(1-11) (GA and GB are grains A and B) and the corresponding planes for GA and GB are (1-11) and (1-1-1) respectively. The CTBs ends inside the crystal with an ∑3{112}-ITB (Fig 2b) having 6, 7, 5 rings along the boundary. Figure 1c shows an image of the CTB at the interface, the yellow arrows are the PDs accommodating the mismatch in the CTB. They are 30° Shockley PDs since the Ge grows in compressive films. The CTB are formed nearby the steps formed by the pairs of MDs.

Figure 1d shows a misfit partial dislocation (MPD) at the end of a SF.  The Burgers circuit around the dislocation gives b=1/6[1-12] which corresponds to a Shockley PD. The stacking sequence of the SF changes from AaBbCcAaBbCc… to AaBbCcBbAaBbCc (Fig. 2e). The SF is of extrinsic type since there is an extra plane in the stacking sequence. For the compressive Ge, the Shockley MPD with an extrinsic SF corresponds to 90° at the interface which in addition has  a perfect 60°  dislocation close to the PD. Figure 2f shows two SFs in (1-11) and (1-1-1) planes, they annihilate each other forming a stair rod dislocation (SRD). The stacking sequence of the SFs changes from AaBbCcAaBbCc… to AaBbCcBbCcAaBb… The SFs are of intrinsic type since there is a missing plane in the stacking sequence. The configuration of the MPDs at the interface in Ge with intrinsic SFs is 30° MPD at the interface and 90° MPD in the Ge. The two 90° PDs in the Ge interact with each other and form the SRD with b=1/3[-110].

SFs are also found in the Ge crystal (Fig. 1g). The Burgers circuit around the PDs gives b=1/6[1-12] which corresponds to Shockley PDs. These intrinsic SFs (Fig. 1h) are formed by the dissociation of perfect 60° dislocations.

We conclude that the misfit strain between Si and Ge is accommodated by 60° and 90° MDs, the 60° MDs splitting to form MPDs and CTB. The MPDs form extrinsic or intrinsic SFs. Intrinsic SFs are interacting with each other forming SRD. Overall, it can be summarized that the defect chemistry in Ge pillars is more complex than reported in earlier studies. This might have significant impact on the electronic properties of these binary semiconductor heterostructures.

Yadira ARROYO ROJAS DASILVA (Dübendorf, SWITZERLAND), Marta D. ROSSELL, Rolf ERNI, Fabio ISA, Giovanni ISELLA, Hans VON KÄNEL, Pierangelo GRÖNING

13:45-15:45
Added to your list of favorites
Deleted from your list of favorites

IM8-IV
IM8: Spectromicroscopies and analytical microscopy
SLOT IV

IM8: Spectromicroscopies and analytical microscopy
SLOT IV

Chairmen: Gerald KOTHLEITNER (Graz, AUSTRIA), Anders MEIBOM (Lausanne, SWITZERLAND), Bénédicte WAROT-FONROSE (Toulouse, FRANCE)
13:45 - 14:15 #8675 - IM08-S56 From core and valence excitations to orbital mapping: a theorist's perspective.
From core and valence excitations to orbital mapping: a theorist's perspective.

Ab initio spectroscopy is a powerful combination of quantum-based theories and computer simulations, covering a wide range of theoretical and computational methods that incorporate many-body effects and interactions showing up in the excited state. This framework, combining density-functional theory (DFT) with many-body perturbation theory, not only allows for analyzing data obtained by experimental probes, but also for shining light onto the underlying physics. I will demonstrate this with a series of selected materials and excitation processes:

Oxygen K-edge spectra from the wide-gap semiconductor Ga2O3 will reveal how signals from atoms located in a particular environment can be selectively enhanced or quenched by adjusting the crystal orientation [1]. These results suggests ELNES, combined with ab initio many-body theory, to be a very powerful technique to characterize complex systems, with sensitivity to individual atomic species and their local environment.

With the example of self-assembled phases of functionalized azo-benzene it will be shown, how excitonic effects in core and valence excitations can be tuned by molecular packing [2], and how this may affect the switching functionality of the molecules.

Finally, I will discuss how transmission electron microscopy can be used for mapping atomic orbitals, exploring its capabilities by a first principles approach [3]. For defected graphene, exhibiting either an isolated vacancy or a substitutional nitrogen atom, different kinds of images are to be expected, depending on the orbital character (see Figure 1).

 

[1] C. Cocchi, H. Zschiesche, D. Nabok, A. Mogilatenko, M. Albrecht, Z. Galazka, H. Kirmse, C. Draxl, and C. T. Koch, submitted to Phys. Rev. B (2016).

[2] C. Cocchi and C. Draxl, Phys. Rev. B 92, 205105 (2015).

[3] L. Pardini, S. Löffler, G. Biddau, R. Hambach, U. Kaiser, C. Draxl, and P. Schattschneider, Phys. Rev. Lett. (2016); in print.

Claudia DRAXL (Berlin, GERMANY)
Invited
14:15 - 14:30 #5709 - IM08-OP140 EELS simulations in III-Nitride ternary alloys by DFT.
EELS simulations in III-Nitride ternary alloys by DFT.

III-V nitride ternary alloys, composed of two different third-column metals, e.g. Al, Ga, In... , and nitrogen, as in the case of AlxGa1-xN, are semiconductor materials that for the last years have been playing a crucial role in the development of novel applications. They are of foremost importance for the optoelectronic industry, for instance for the recent development of blue laser applications. Often in these devices, the desirable reduction of the typical integrated circuit dimensions is translated in increasing challenges to the growth and characterization techniques employed. Among the later, analytical transmission electron microscope (TEM) is an invaluable for its ability to obtain structural and chemical information about the structures and materials at the nanometer scale. For instance, electron energy-loss spectroscopy (EELS), a technique that is available in most modern TEM machines, allows the measurement of important valence properties by probing the low-loss region of the spectrum, containing signals from inter-band transitions and plasmon excitation.

 

We present a theoretical study of low-loss EELS using super-cell models for different concentrations of the metals, x, that allow to systematically study the whole compositional range, 0<x<1, with Δx = 0.125 resolution [1]. This study is carried out for the three foremost III-nitride semiconductor ternary alloys, AlxGa1-xN, InxAl1-xN and InxGa1-xN. In order to do this, automated DFT simulations have been carried out using Wien2k software and home-made scripts. Additionally to the typical DFT simulation scheme, we have corrected our calculations using the modified Becke-Johnson (mBJ) exchange-correlation potential. This correction represents a critical improvement over the former calculation, using generalized gradient approximation (GGA), which predicted wrong band-gap values.

 

For each concentration, x, of the ternary nitride compounds, AxB1-xN, with elements A and B combinations of Al, Ga and In, we obtain from our DFT simulations the complex dielectric function (CDF), ε(E) = ε1+i·ε2, where E is the energy-loss. Energy-loss spectra are proportional to the imaginary part of the inverse CDF, Im[-1/ε], also called the energy-loss function (ELF, see Figs. 1 and 2). Figure 1 contains the ELF-series obtained for the three studied ternary nitrides. In these series, the composition-related behavior of the most intense peak in EELS, the plasmon, is depicted.  It is generally accepted that the observed features in ternary nitride EELS, like band gap and plasmon onset energy, are related to the features observed in the pure binaries through a Vegard law of the form,

 

Ei(AxB1-xN) = x·Ei(AN) +(1-x)·Ei(BN) + x·(1-x)·b .

 

Where Ei is the observed energy for a feature i, and b is called the bowing parameter. We have used this formula to analyze both band gap and plasmon energy, Egap and Ep, respectively. In this sense, Egap is directly measured in the calculated band structures and density of states. Conversely, Ep is retrieved from a model-based fit of the ELF series. For this purpose, we chose the Drude model of quasi-free electron gas, which is a typical approach in experimental EELS. Figure 3 contains the results from these analyses, where the Egap and Ep appear in red and green color, respectively. The results show a somewhat inconsistent behavior of these two parameters in terms of slope and bowing of the derived Vegard laws (solid lines). This problem is especially poignant in In-rich compounds, in which the band gap offset is greater and also interband transitions are more important.

 

Because of this inconsistency, we have developed an alternative method to locate the plasmon energy, following the zero of the real part of the CDF (Fig. 2); Ecut, such that ε1(Ecut) = 0. A more reasonable agreement with the theoretical band gap as well as the experimentally measured Vegard law is obtained from this parameter. Finally, the role played by interband transitions in the calculations of In-rich compounds is also addressed.

[1]: A. Eljarrat, X. Sastre, S. Estradé and F. Peiró Microscopy and Microanalysis 2016 (Accepted) doi: 10.1017/S1431927616000106.

Alberto ELJARRAT ASCUNCE (Barcelona, SPAIN), Xavier SASTRE, Sònia ESTRADÉ, Francesca PEIRÓ
14:30 - 14:45 #5913 - IM08-OP146 Dislocation Modelling: Calculating EELS Spectra for Edge Dislocation in Bismuth Ferrite.
Dislocation Modelling: Calculating EELS Spectra for Edge Dislocation in Bismuth Ferrite.

Recent advances in electron energy-loss spectroscopy (EELS) triggered by the implementation of aberration correctors and novel spectrometers have enabled atomic resolution and single atom sensitivity. The energy-loss near-edge structure (ELNES) in core-loss EELS provides insight into the electronic structure of individual atomic species containing information about their bonding characteristics such as, e.g., oxidation state, charge transfer and site coordination. Yet the electronic structure information is buried in the spectral fine structure which can be regarded as a “fingerprint” of the atom’s bonding characteristics. In order to overcome this shortcoming of a purely experimental approach, the ELNES of core-loss EELS ionization edges can be obtained from first-principles electronic structure calculations.

 

BiFeO3 (BFO) is a multiferroic perovskite that exhibits antiferromagnetism coupled with ferroelectric order. Besides, because of their astounding electromechanical properties, BFO thin films are promising candidates for the replacement of lead-based ceramics in microelectromechanical system devices. It is well known that performance of ferroelectric devices is reduced by the presence of crystal defects such as edge dislocations. This type of crystal defects within this material1 are due to the lattice mismatch between the BFO film and the SrRuO3 substrate as  strain compensation mechanism. However, the electronic structure close to the dislocation core is not yet well understood. In this work, we investigate the influence of edge dislocations on the material’s local electronic properties, using a combined experimental and theoretical strategy based on HAADF-STEM, EELS and atomistic simulations.

 

In this study the edge dislocation within BFO was modelled based on Peierls-Nabarro Model (P-N model)2 theory. The initial guess was obtained from the P-N model and further optimized using BVVS classical potentials for BFO3 as incorporated in the LAMMPS package4. Fig 1 shows the final geometry of the edge dislocation model obtained after optimizing it with this potential. FEFF95, a real space multi-scattering  Core level spectroscopies code will be used on the final optimized model to obtain the O K-edge and Fe L-edge spectra at the dislocation core.  Fig 2 shows the experimentally recorded O K-edge EELS on a line of atomic columns (points 1 to 9) crossing the dislocation core. The comparison of the experimental spectrum with the calculated EELS allows shedding atomistic insight on the EELS peak structure. In particular, it is possible to explain to which extent a defect in a bulk material (in this case, an edge dislocation) locally affects its electronic properties, thus enhancing the power of electron energy-loss spectroscopy in the high-resolution description of complex materials.

 

 

(1)           Lubk, A.; Rossell, M. D.; Seidel, J.; Chu, Y. H.; Ramesh, R.; Hÿtch, M. J.; Snoeck, E. Electromechanical Coupling among Edge Dislocations, Domain Walls, and Nanodomains in BiFeO3 Revealed by Unit-Cell-Wise Strain and Polarization Maps. Nano Lett. 2013, 13 (4), 1410–1415.

(2)           Yao, Y.; Wang, T.; Wang, C. Peierls-Nabarro Model of Interfacial Misfit Dislocation: An Analytic Solution. Phys. Rev. B 1999, 59 (12), 8232–8236.

(3)           Liu, S.; Grinberg, I.; Rappe, A. M. Development of a Bond-Valence Based Interatomic Potential for BiFeO 3 for Accurate Molecular Dynamics Simulations. J. Phys. Condens. Matter 2013, 25 (10), 102202.

(4)           Plimpton, S. Fast Parallel Algorithms for Short-Range Molecular Dynamics. J. Comput. Phys. 1995, 117 (1), 1–19.

(5)           Rehr, J. J.; Kas, J. J.; Vila, F. D.; Prange, M. P.; Jorissen, K. Parameter-Free Calculations of X-Ray Spectra with FEFF9. Phys. Chem. Chem. Phys. 2010, 12 (21), 5503–5513.

Piyush AGRAWAL (Dübendorf, SWITZERLAND), Marta D. ROSSELL, Cécile HÉBERT, Daniele PASSERONE, Rolf ERNI
14:45 - 15:00 #6448 - IM08-OP157 Theoretical and numerical investigation of the interaction between phase-shaped electron probes and plasmonic modes.
Theoretical and numerical investigation of the interaction between phase-shaped electron probes and plasmonic modes.

Optical vortices - i.e. electromagnetic fields with phase singularity - are well known objects and have been used for a decade in a wide range of applications in e.g. optics or astrophysics. Recently, it has been demonstrated that such vortices can be created in an electron microscope by tailoring the phase of the beam [1] and these so-called electron-vortex-beams have already proven their efficiency detecting magnetic state in a material or to probe chirality in a crystal [2].

Simultaneously, EEL spectroscopy in the low-loss region has attracted a tremendous interest due to its efficiency in resolving plasmonic resonance at the nanometer scale [3] and the underlying formalism is now firmly established [4]. However, because of the invariance of the electron probe along the propagation axis, low-loss EELS remained unable to detect plasmonic optical activity. However, electron-vortex-beams constitute a perfect candidate to overcome this limitation and measure the dichroic behavior of plasmons in an electron microscope - as recently pointed out through simulations by Asenjo-Garcia and García de Abajo [5].

In the present work, we developed a semiclassical formalism describing the interaction between an electron probe with an arbitrary phase profile and a plasmonic mode. Following [6], we used a quasi-static and classical description of the plasmon resonances while the electron probe is described in a fully quantum way. We showed that the equation ruling this interaction takes the elegant form of a transition matrix - between two electron states mediated by the eigenpotentials of the plasmon modes. Starting from this formalism, we built an analytical model describing the interaction between point charges and a vortex beam, which gave us a good insight into the physics of plasmonic dichroism. Important experimental inputs, such as convergence and collection angles, were considered. We also implemented a Matlab script within MNPBEM [7] in order to compute our equation and investigate the dichroic behavior of arbitrary plasmonic nano-structures (see Figure 1). 

In the conference, we will present the theoretical formalism and a wide variety of numerical studies of interactions between different nano-structures (e.g. helix, rod) and phase shaped electron probes (e.g. vortex beams, HG-like beams...), with a special emphasis on the experimental feasibility of the proposed geometries.

Acknowledgments:

GG, and JV acknowledge funding from ERC Starting Grant No. 278510-VORTEX.JV and MK also acknowledge under a contract for an Integrated Infrastructure Initiative, reference No. 312483- ESTEEM2

References

[1] Verbeeck and al, Nature 467, 301-304, 2010.

[2] Juchtmans and al, Phys. Rev. B 91, 094112, 2015.

[3] Colliex and al, Ultramicroscopy 162, A1-A24, 2016.

[4] García de Abajo, Kociak, Phys. Rev. Lett. 100, 106804, 2008.

[5] Asenjo-Garcia, García de Abajo, Phys. Rev. Lett. 113, 066102, 2014.

[6] Boudarham, Kociak, Phys. Rev. B 85, 245447, 2012.

[7] Hohenester, Trügler, Comput. Phys. Commun 183, 370, 2012.

Hugo LOURENÇO MARTINS (Palaiseau), Giulio GUZZINATI, Jo VERBEECK, Mathieu KOCIAK
15:00 - 15:15 #5877 - IM08-OP144 Determination of elemental ratio in an atomic column by STEM-EELS.
Determination of elemental ratio in an atomic column by STEM-EELS.

The elemental signals do not necessarily localize at atomic-column positions because the spatial resolution of an EELS signal is constrained by the delocalization of inelastic scattering and electron channeling process. These complexities make it difficult to perform quantitative analysis with atomic resolution. When we estimate the exact value of elemental ratio with atomic scale, full quantum mechanical simulations combined with experimental result are necessary. On the other hand, if there is a criterion about accuracy of experimental result about elemental ratio for an atomic column without simulation, it would be very useful. 

In this study, atomic-resolution quantification of the elemental ratio of Fe to Mn at octahedral and tetrahedral sites in brownmillerite Ca2Fe1.07Mn0.93O5 (Fig. 1) is demonstrated using STEM-EELS. It is known that Fe and Mn ions are nearly all ordered but not fully ordered, i.e., a small number of Fe and Mn ions reside in octahedral and tetrahedral sites, respectively. It was found that a considerable oversampling of the spectral imaging data yields a spatially resolved area that very nearly reflects atomic resolution (~1.2 Å in radius) for Fe and Mn L2,3-edge (Fig. 2). And the average relative compositions of Fe to Mn within the region were 17.7 to 82.3 ± 13.1 in octahedral sites and 80.7 to 19.3 ± 9.9 in tetrahedral sites. The actual atomic ratio was estimated by calculating the mixing of signals from nearest-neighbor columns using simple simulation based on multislice technique. It was concluded that the ratio of Fe to Mn was 14 to 86 at octahedral sites. It agrees well with the previous neutron diffraction experiment (14.4 to 85.6) which can correctly decide such information for bulk sample [1]. On the other hand, the experimental value and the estimation value from tetrahedral site have relatively large error compare with the result of neutron diffraction experiment (92.2 to 7.8). This means that an experimental oversampling SI data of Fe and Mn L2,3-edge from octahedral site in perovskite-like structure is probably interpreted with an uncertainty of approximately 10% without simulation.

[1] Hosaka, Y.; Ichikawa, N.; Saito, T.; Haruta, M.; Kimoto, K.; Kurata, H.; Shimakawa, Y. Bull. Chem. Soc. Jpn. 2015, 88, 657-661

Acknowledgements

This work was supported by JSPS KAKENHI Grant Numbers 26706015, 19GS0207 and 22740227. It is also supported by a grant for the Joint Project of Chemical Synthesis Core Research Institutions from the Ministry of Education, Culture, Sports, Science and Technology of Japan and by the Japan Science and Technology Agency, CREST.

Mitsutaka HARUTA (Kyoto, JAPAN), Yoshiteru HOSAKA, Noriya ICHIKAWA, Takashi SAITO, Yuichi SHIMAKAWA, Hiroki KURATA
15:15 - 15:30 #6374 - IM08-OP156 How multiple scattering simulations help for EELS compositional analysis of hard metals and ceramics.
How multiple scattering simulations help for EELS compositional analysis of hard metals and ceramics.

In the manufacturing process of hard metals and ceramics used as tooling materials, there is strong interest to determine the phases formed within the sintered bulk material or the thin coating of this sintered substrate. A compositional analysis of these, frequently sub-stoichiometric, phases based on electron energy-loss spectroscopy (EELS) is challenging and requires - besides the knowledge of other parameters - accurate ionization cross-sections. Most commonly, EELS ionization cross-sections are derived from analytical models (hydrogenic approximation) or from Hartree Slater oszillator strengths, and hence lack EELS fine-structure details (ELNES). Since fine-structures, however, can be indicative of chemical phases, a proper calculation of cross-sections around the ELNES regime can be beneficial for a more reliable analysis.

For this, an alternative route was followed, trying to obtain detailed differential cross-sections from ab initio multiple scattering calculations, as implemented in the FEFF9 code [1]. With this tool one can calculate EEL spectra (and energy differential cross-sections) based on Green’s functions theory when fed with crystal structure data.

 

One point of interest first of all is how integrated cross-sections for K-shell ionization edges between the hydrogenic model, calculated with the SIGMAK3 program [2] and FEFF9 [1] simulations compare with each other (Conditions for calculations were: E0= 200 eV; β= 5 mrad, integration window (Δ)= 200 eV starting at the edge’s threshold) (Fig.1). This is also interesting in the light of the fact that, over the past decades, hydrogenic calculations have been adjusted to better match experimental data for instance by adjusting the inner-shell screening constant (s) for non hydrogen-like atoms [e.g. 3, 4] (Fig.1).

Secondly, structural data, needed as input for FEFF, for known phases can be taken from data bases (e.g. FIZ Karlsruhe). Alternate, for materials with unknown structure or undefined phases the atomic positions can be determined by electron diffraction tomography experiments directly from single nano crystalline domains, as shown by Kolb et al. [e.g. 5,6].

In certain cases selected area electron diffraction (SAED) along with diffraction pattern (DP) simulations can help for phase determination of potential EELS reference samples as the example of TixSiC1-x shows (Fig. 2): three different structures of TixSiC1-x are described in literature [7]. The different stoichiometries can be distinguished by the length of the c-axis as there are a different numbers of Ti-layers between the Si-layers (Fig. 2.a). SAED does not only reveal epitactic growth of the carbide coating on the corundum substrate (Fig. 2(b, d, e)) but also, when compared to a diffraction simulation (JEMS) the stoichiometry of the coating can be clearly determined as Ti3SiC2 as the distances and intensities of the reflexions match the simulation (Fig. 2(c, f)).

The paper will discuss new possibilities in quantification and findings with this approach.

 

Acknowledgments

This work was carried out with the financial support by Sandvik Coromant and Sandvik Mining.

 

References:

[1] J.J. Rehr et al., Phys. Chem. Chem. Phys., 12, pp. 5503-5513 (2010).

[2] R.F. Egerton, EELS in the Electron Microscope, 3rd edition, Springer (2011).

[3] E. Clementi and D.L. Raimondi, J. Chem. Phys., 38, pp. 2686-2689 (1963).

[4] R.F. Egerton, Ultramicroscopy, 63, pp. 11-13 (1996).

[5] U. Kolb et al., Ultramicroscopy 107: 507-513 (2007).

[6] U. Kolb et al., Ultramicroscopy 108: 763-772 (2008).

[7] H. Högberg et al., Surf. Coat. Technol. 193: 6.10 (2005).

Lukas KONRAD (Graz, AUSTRIA), Martina LATTEMANN, John REHR, Ute KOLB, Zhao HAISHUANG, Gerald KOTHLEITNER
15:30 - 15:45 #6336 - IM08-OP152 New Opportunities in multi-frame STEM Spectroscopy & Fractional Beam-current EELS.
New Opportunities in multi-frame STEM Spectroscopy & Fractional Beam-current EELS.

    Electron energy-loss spectroscopy (EELS) and energy-dispersive x-ray spectroscopy (EDX) are two of the most common means of chemical analysis in the scanning transmission electron microscope (STEM). While the instrumentation hardware has progressed markedly in recent years, the way that microscopists operate these spectrometers has changed very little. In general small areas are scanned just using coarse pixilation, slow scan-speeds and high beam-currents. Finally, once acquired these chemical signals are usually expressed in only relative or arbitrary units. This stands in stark contrast with modern best-practice in STEM imaging, where wider fields of view are surveyed utilising multi-frame acquisitions, faster scan speeds, finer pixel sampling and lower electron doses. More recently it has also become common place to express STEM image data in units of fractional-beam-current facilitating direct comparison with simulation.

    We will present EELS and EDX results where best-practice techniques from STEM imaging have been repurposed to improve chemical map quality. Multi-frame spectrum-image data were recorded with simultaneous EELS and EDX spectra, before using non-rigid data registration [1] in the spatial domain, Figure 1. This not only reduces scan-distortion but also unlocks tools such as; digital super-resolution, strain mapping performed directly on chemical maps, and the digital accumulation of weak signals such as those from monochromated EELS. We also present an ‘equal fixed-dose’ fractionation study where sample damage was reduced drastically using a fast-scanning multi frame approach compared with its single scan equivalent.

    EELS spectrum data fidelity was improved by energy-drift tracking and correction [2,3]. After this, the limiting performance in multi-frame spectra becomes the spectral-noise and was found to depend strongly on the quality of the available dark-reference, Figure 2. To mitigate this artefact, separate dark-references and full-beam gain-references were recorded along with wholly un-processed experimental spectra. Further improvement in the noise performance was released by stepping the EELS spectra in energy between each scan before un-stepping them in post-processing [4]. Dark-correction was performed offline before dividing by the gain-reference to finally express the EELS spectra in terms of fractional beam current [5].

    It is expected that leveraging all these individually small improvements collectively will deliver more precise chemical maps while minimising sample damage and experimental time overheads.

Acknowledgments

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 EPSRC grant code, EP/K040375/1, for the ‘South of England Analytical Electron Microscope’. SuperSTEM is the UK's facility for aberration corrected STEM funded by EPSRC.

References

[1]         L. Jones, H. Yang, T. J. Pennycook, M. S. J. Marshall, S. Van Aert, N. D. Browning, M. R. Castell, and P. D. Nellist, Adv. Struct. Chem. Imaging 1, 8 (2015).

[2]         K. Kimoto, K. Ishizuka, T. Asaka, T. Nagai, and Y. Matsui, Micron 36, 465 (2005).

[3]         Y. Sasano and S. Muto, J. Electron Microsc. (Tokyo). 57, 149 (2008).

[4]         M. Bosman and V. J. Keast, Ultramicroscopy 108, 837 (2008).

[5]          Y. Zhu and C. Dwyer, Microsc. Microanal. 20, 1070 (2014).

Lewys JONES (Oxford, UK), Aakash VARAMBHIA, Demie KEPAPTSOGLOU, Quentin RAMASSE, Robert FREER, Feridoon AZOUGH, Sergio LOZANO-PEREZ, Richard BEANLAND, Peter NELLIST

13:45-15:45
Added to your list of favorites
Deleted from your list of favorites

IM3-IV
IM3: Innovative SEM, imaging and analytical instruments
SLOT IV

IM3: Innovative SEM, imaging and analytical instruments
SLOT IV

Chairmen: Emmanuel BEAUREPAIRE (Palaiseau, FRANCE), Christian COLLIEX (Orsay, FRANCE), Jörg ENDERLEIN (Göttingen, GERMANY), Andreas ENGEL (Delft, THE NETHERLANDS), Ernst H.K. STELZER (Professor) (Frankfurt am Main, GERMANY)
13:45 - 14:15 #7823 - IM03-S43 Integrated microscopy: Matching scales and capabilities in light and electron microscopy.
Integrated microscopy: Matching scales and capabilities in light and electron microscopy.

Superresolution techniques have pushed the resolution of fluorescence microscopy (FM) towards that of electron microscopy (EM)[1]. Meanwhile, developments in scanning EM (SEM) are revolutionizing EM, moving lateral image dimensions to typical FM fields-of-view[2] and extending imaging capability into the third dimension[3] and the live-cell regime[4]. By correlating data from both techniques[5], molecules can be localized within the context of cells and tissue and with reference to their live dynamics, but throughput and quantification are hindered by elaborate, expert procedures involving separate microscopes. In this presentation, I will show an integrated approach, with high-numerical aperture FM in a SEM, such that the electron beam can be positioned anywhere within the fluorescence field of view[6, 7]. Using electron-beam excited cathodoluminescence from the transparent sample substrate, we achieve automated FM-EM image registration (Fig. 1) with an accuracy that can be pushed to 5 nm, i.e. equalling biomolecular length scales. Besides integrated correlation microscopy, I will show our progress towards novel applications bridging fluorescence and electron microscopy, such as fluorescence-guided live-cell EM[8], and electron-beam identification and localization of labels, molecules, and cells.

[1]   B. Huang, M. Bates, and X. Zhuang, Annual review of biochemistry, 78, 993-1016 (2009).

[2]   R. B. G. Ravelli et al, Scientific Reports 3 (2013)

[3]   C. J. Peddie, and L.M. Collinson, Micron 61, 9-19(2014)

[4]   N. de Jonge, and F. M. Ross, Nature Nanotechnology 6 (11), 695-704 (2011)

[5]   P. de Boer, J.P. Hoogenboom, and B.N.G. Giepmans, Nature Methods 12(6.), 503–513 (2015).

[6]   A.C. Zonnevylle et al., Journal of Microscopy 252, 58-70 (2013).

[7]   N. Liv et al, PLoS ONE 8 (2), e55707 (2013)

[8]   N. Liv et al, ACS Nano 10, 265-273 (2016)

Jacob HOOGENBOOM (Delft, THE NETHERLANDS)
Invited
14:15 - 14:30 #5983 - IM03-OP090 Beyond elemental analysis in the electron microscope: accessing isotopes with in-situ TEM-SIMS correlative analysis.
Beyond elemental analysis in the electron microscope: accessing isotopes with in-situ TEM-SIMS correlative analysis.

Despite having superior spatial resolution, the analytical capabilities of Transmission Electron Microscopy (TEM), either Energy Dispersive X-ray Spectroscopy (EDS) or Electron Energy-Loss Spectroscopy (EELS) [1], are fundamentally limited in terms of isotopic analysis. Moreover, the analysis of light elements still remains out of reach. However, these analyses count among the advantages of Secondary Ion Mass Spectrometry (SIMS), which also provides high chemical sensitivity and high dynamical range, but offers poor lateral resolution [2]. Therefore, a combination of both techniques is the logical step to complement the strengths of SIMS with a high-resolution imaging technique such as TEM. As ex-situ approaches are prone to sample modification artefacts, we have developed an in-situ combination of TEM and SIMS called the Parallel Ion-Electron Spectrometry (PIES).

The developed prototype instrument [3] is based on a Tecnai F20 which has had its octagon and pole pieces modified to accommodate a FEI Magnum Ga+ focused Ion Beam (FIB) primary column and secondary ion extraction optics. This extraction optics, with extraction efficiency up to 90%, is coupled with a double focusing magnetic-sector mass spectrometer developed in house. Finally, another needed addition involved the sample holder, which was designed to be biased to high voltages (4.5 kV) and to act as the first electrode in the extraction system.

To highlight the new methodology developed in the instrument, lithium carbonate (Li2CO3) was chosen. Lithium exists in two stable isotopes in nature (6Li and 7Li) with very different abundances, 7.5% and 92.5%, respectively. A powder sample of Li2CO3 enriched in 6Li up to 95% was mixed with a natural-abundance sample. The goal of the investigation was the distinction of particles according to their isotopic abundance.

The correlative approaches for the information extraction can be regarded as two-directional. The first option, for samples with areas of interest depending on structure, TEM should be used first for identification of those areas, while SIMS will provide the analysis afterwards. The second option suits samples marked with isotopic (or elemental) labels. SIMS will then be used for localisation of the area of interest, to be further imaged at high resolution with TEM. Using this second approach, mass spectra were acquired for each of the individual materials before mixing them. Mass filtered images were acquired form the mixture. The subsequent imaging with TEM of the same area was used to assign each of the particles to one of the original materials according to their isotopic composition.

As the correlative approach leads to a multimodal imaging of the same regions of interest, and, in order to successfully exploit the results, there is need of correction of imaging artefacts as well as data fusion for obtaining high spatial resolution isotopic images. We will show how these new kinds of datasets can be obtained with this new instrument.

References

[1] R. F. Egerton, “Electron Energy-Loss Spectroscopy in the Electron Microscope” (3rd edition), (Springer, New York, 2011)

[2] A. Benninghoven, H. W. Werner, F. G. Rüdenauer, “Secondary ion mass spectrometry: Basic concepts, instrumental aspects, applications and trends” (Wiley-Interscience, New York, Chichester, Brisbane, Toronto, Singapore, 1987)

[3] T. Wirtz, et al, Nanotechnology 26, 434001 (2015)

Lluis YEDRA (Belvaux, LUXEMBOURG), Santhana ESWARA, David DOWSETT, Tom WIRTZ
14:30 - 14:45 #6871 - IM03-OP100 A JEOL-based cooling holder with a low specimen drift allowing sub 1Å STEM imaging.
A JEOL-based cooling holder with a low specimen drift allowing sub 1Å STEM imaging.

With the development of aberration corrected microscopy has come the necessity of more stable sample stages, one area that needs further refinement is cryogenic sample system, both for biological imaging and for sensitive materials samples that would benefit from cryo-stability.  The basic principles of the holder are given in Fig. 1. It shows the main components of the cooling holder, the main coldness losses and the measures taken to create a thermal equilibrium, such that the drift is minimized. This allows for drifts of less than 2 nm/min and sub 1 Å resolution. In the cross-section of the holder, schematically given in Fig. 1, the outside tube contains a heater close to thermal isolator near the tip, such that the temperature of the outside tube can be regulated to the temperature of the goniometer (or any other temperature). This is done to minimize the drift related to thermal expansion and to prevent any cooling (from the tip of the holder) of parts of the goniometer. 

The thermal expansion stability is important since the cooling of the goniometer parts can result in unpredictable (slow) changes in the specimen position, for instance in the electron beam direction resulting in a focus change. In case of the temperature difference between the holder and the goniometer, the heat transfer between the outside rod and the goniometer is uncertain because it depends on the mechanical contact (that can vary with tilt) and the effect of the thermal gradient on each of the relevant components in the goniometer.   Likewise the “coldness” influx depends on quite a number of parameters, like the size of the beard, how deep it is in the liquid nitrogen, whether the dewar has a “good” cover to isolate and to prevent turbulence. For that reason we have an option of a heater close to the cryo-beard. Fig. 2 shows the experimental setup of the JEOL ARM, the holder and the dewar.

Our initial tests with the prototype sample holder on a sample of small gold particles on a thin C film show a resolution stability in HRTEM showing sub 1Å information transfer (Fig. 3).  This stability is sufficient to operate the holder in STEM HAADF imaging conditions (Fig. 4) which was acquired at 2k resolution with 17.6 us pixel dwell time.  Further work needs to be done to increase the operational duty cycle, however the basic configuration of the holder seems to provide a workable solution for high-resolution cryo-imaging.

David BELL (Cambridge, USA), Henny ZANDBERGEN
14:45 - 15:00 #6551 - IM03-OP096 Power of FIB-SEM Tomography for Biological Samples.
Power of FIB-SEM Tomography for Biological Samples.

Three-dimensional (3-D) spatial distribution of organelles within cells at nanometer resolution is essential to better understand cellular processes and functions. Currently, FIB-SEM tomography provides a promising technology to generate volume data with nanometer resolution (Bushby et al., 2011; Heymann et al., 2009; Knott et al., 2011; Villinger et al., 2012; Wei et al., 2012). A FIB-SEM microscope is a scanning electron microscope (SEM) combined with a focused ion beam (FIB) where both beams coincide at their focal points. This combination enables to locally section bulk embedded-resin samples by ion milling, creating a fresh surface for subsequent imaging with the electron beam. This process is repeated automatically to generate 3-D information of relatively large volumes with a field of view of several micrometers.

Any stained and embedded resin samples prepared for transmission microscope can be used for FIB-SEM tomography offering the possibility to visualise directly in 3D a wide range of biological samples and typically volumes with a pixel size of 10 nm can be achieve (Figure1). However, to obtain a better resolution, considerations have to be given concerning ultrastructure preservation and samples preparation (Kizilyaprak et al., 2015). High pressure freezing (HPF) combined with freeze-substitution (FS) and resin embedding constitutes a method of choice to find the best compromise between ultrastructural preservation and high contrast of cellular components (Figure2).

In conclusion, we propose biological sample preparation protocols that can serve as starting point to visualize in 3-D wide range of biological samples at nanometer resolution including HPF/FS samples and correlative microscopy approach using FIB-SEM Tomography.

References:

Bushby, A.J., P'Ng K, M., Young, R.D., Pinali, C., Knupp, C., Quantock, A.J., 2011. Nat Protoc 6, 845-858.

Heymann, J.A., Shi, D., Kim, S., Bliss, D., Milne, J.L., Subramaniam, S., 2009. Journal of structural biology 166, 1-7.

Kizilyaprak, C., Longo, G., Daraspe, J., Humbel, B.M., 2015. Journal of structural biology 189(2), 135-146.

Knott, G., Rosset, S., Cantoni, M., 2011. J Vis Exp, e2588.

Villinger, C., Gregorius, H., Kranz, C., Hohn, K., Munzberg, C., von Wichert, G., Mizaikoff, B., Wanner, G., Walther, P., 2012. Histochem Cell Biol 138, 549-556.

Wei, D., Jacobs, S., Modla, S., Zhang, S., Young, C.L., Cirino, R., Caplan, J., Czymmek, K., 2012. Biotechniques 53, 41-48.

Caroline KIZILYAPRAK (Lausanne, SWITZERLAND), Florence NIEDERGANG, Damien DE BELLIS, Willy BLANCHARD, Jean DARASPE, Niko GELDNER, Bruno HUMBEL
15:00 - 15:15 #6736 - IM03-OP097 Helios G4: Combination of ultrathin damage-free TEM sample preparation and high-resolution STEM imaging in a single instrument.
Helios G4: Combination of ultrathin damage-free TEM sample preparation and high-resolution STEM imaging in a single instrument.

Development in semiconductor industry as well as in materials research has lead to a further decrease in observed features sizes. STEM in SEM imaging has been a well based and widely used technique to address this type of work. As a consequence of reducing feature size, there have been parallel requirement increases for high quality TEM sample preparation as well as for improved optical and detection performance of SEM/FIB systems. Here we introduce the very new generation of the FEI Helios product family, which addresses both of upper demands (preparation and resolution). Helios G4 introduces a brand new workflow, which combines ultrathin TEM sample preparation and high resolution STEM imaging in a single instrument.

 

Newly introduced Helios system enables the preparation of all common sampling methods: top-down, plan view, inverted and tomography. It integrates three sample manipulation devices, the well-established piezo stage for sample bulk processing, nanomanipulator for sample lift-out and double-tilt STEM Rod (TEM-like manipulator) for lamella thinning and high-resolution STEM imaging. All  manipulators are controlled via the microscope controller software, which allows automated switching between each manipulator and enables automation of the complete workflow. STEM Rod design with opening from one side enables access to sample by Focus Ion Beam (FIB), gas injection system, nanomanipulator and also provides optimized signal collection by analytical detector such as X-Ray EDS.

 

Focus Ion Beam is a known and commonly used technology for thin sample preparation. Helios G4 family employs the newly developed Phoenix FIB, which brings an advancement in low kV performance, allowing for precise beam placement at low acceleration voltage down to 500 V without compromising milling rates at high voltages. As a result, operator can create site-specific TEM samples with thickness below 10 nm with damage layer less than 1 nm.

 

Sample lift-out and transfer from bulk sample to the liftout grid is realized using FEI EasyLiftTM nanomanipulator. The grid is loaded in the STEM Rod with double-tilt functionality for SEM to access both sides of the lamella for precise end-pointing. STEM Rod is used for final thinning and cleaning and for subsequent high-resolution STEM imaging of prepared sample. The whole process is done in a system chamber without breaking vacuum. STEM Rod is also designed for sample transfer out of the specimen chamber for further analysis. The option is also to load previously prepared sample into the system and run the STEM imaging.

 

Helios G4, namely FX configuration, introduces the most advanced STEM imaging capability in existing SEM/FIB product portfolio. It is based on proven Elstar SEM column equipped with Schottky field emission gun with improved monochromator, delivering better resolution at higher beam currents. Elstar column works in a setup with configurable objective lens geometry and can be operated in conventional SEM/FIB mode, STEM/FIB end-pointing mode and high-resolution STEM imaging mode. Design of new STEM imaging and detection system brings significant leap forward in  resolution and also improvement in STEM contrast.

 

Helios G4 FX capabilities can be demonstrated on imaging of lattice planes on several materials as Carbon nanotubes (CNT, 0.34 nm at 20-30 kV) or Tungsten Disulfide (0.27 nm at 30kV). Users can benefit from simultaneous collection of various SEM/STEM signal types and acquire information about sample surface, structure and composition in a single scan.

Jan SKALICKY (Brno, CZECH REPUBLIC), Tomas VYSTAVEL, Lubomir TUMA, Richard YOUNG
15:15 - 15:30 #4542 - IM03-OP080 ATOM-PROBE TOMOGRAPHY AND NANOSCIENCES.
ATOM-PROBE TOMOGRAPHY AND NANOSCIENCES.

The design of Atom probe tomography (APT) at Oxford and Rouen universities 25 years ago has been an outstanding breakthrough in the microscopy world. APT is the only analytical microscope able to provide 3D images of a material at the atomic scale [1]. Because of its ultimate spatial resolution (0.1 nm in depth, a few tenths of a nm at the sample surface), combined with its quantitativity in composition measurements, APT has played a major role for the investigation of the segregation of impurities to crystal defects or of the early stages of phase separation in solids. APT was the first to show Cottrell atmospheres (tiny clouds of impurity atoms around dislocations in crystals) at the atomic-scale in the dimensions of space [2].

 

A new breakthrough has been achieved ten years ago with the implementation of ultrafast pulsed laser (duration < 1ps) to atom probe tomography [3]. This new generation of instrument, designed in our lab and abroad, has made it possible the analysis of semi-conductors (figure 1) and oxides that are key materials in micro-electronics and nanosciences [4]. This major innovation with the use of FIB ion milling (focused ion beam)  to prepare samples in the region of interest opened a new insight in many fields of nanoscience related to nano-wires [5], nanostructured magnetic thin films [6], heavily-doped ultra-shallow junctions in microelectronics [7]. One of the biggest challenge being the atomic-scale reconstruction of MOSFET transistors [8,9]. Unique capabilities of atom probe tomography in nanoscience and salient findings will be highlighted on the basis of some selected illustrations. 3D APT reconstructions related to phase separation in GeMn self-organised thin films (figure 2) will be confronted to atomistic Kinetic Monte Carlo simulations conducted on rigid lattice [6].

 

[1] D. Blavette, A. Bostel, J.M. Sarrau, B. Deconihout  and A. Menand, 1993, Nature 363, 432

[2] D. Blavette, E. Cadel, A. Fraczkiewicz, A. Menand, 1999, Science 17, 2317

[3] B. Gault, F. Vurpillot, A. Vella, M. Gilbert, A. Menand, D. Blavette, B., 2006, Rev. Sci. Instr. 77, 043705

[4] S. Duguay, T. Philippe, F. Cristiano, D. Blavette, Applied Physics Letter (2010) 97, 242104

[5]  W. Chen et al. JAP, 111, 094909-094916

[6] I. Mouton, R, Larde, E. Talbot, C. Pareige, D. Blavette, JAP 115, 053515  (2014)

[7] Yang Qiu, Fuccio Cristiano, Karim Huet, Fulvio Mazzamuto, Giuseppe Fisicaro, Antonino La Magna, Maurice Quillec, Nikolay Cherkashin, Huiyuan Wang, Sébastien Duguay, and Didier Blavette, Nanoletters (2014) DOI: 10.1021/nl4042438

[8] R Estivill, M Juhel, M Gregoire, A Grenier, V Delaye, D Blavette , Scripta Materialia  (2015)113, 231-235

[9] A. Grenier, R. Serra, G. Audoit, Jp Barnes, S. Duguay, D. Blavette, N. Rolland, F. Vurpillot, P. Morin, P. Gouraud, Applied Physics Letters 106, 213102 (2015); doi: 10.1063/1.4921352

Didier BLAVETTE (ST ETIENNE ROUVRAY CEDEX), Isabelle MOUTON, Sébastien DUGUAY
15:30 - 15:45 #6933 - IM03-OP101 Dynamic-Transmission Electron Microscopy at the Relativistic Electron Gun for Atomic Exploration (REGAE) for live cell imaging.
Dynamic-Transmission Electron Microscopy at the Relativistic Electron Gun for Atomic Exploration (REGAE) for live cell imaging.

With the relativistic electron gun for atomic exploration (REGAE) we seek to observe structural dynamics both in real space imaging and diffraction [1]. REGAE is based on a RF gun accelerator and operates in the range from 2 to 5 MeV. RF gun technology allows high brightness for electron pulses at high energy. The machine is equipped with an RF cavity system, allowing for picosecond bunches with energy spread compensation to prevent limitations by chromatic aberrations and space charge. A custom made lens system for 3 MeV pulse energy system is set in place.

The presentation will include a discussion on space charge induced aberrations in dynamic HVEM, give estimates about the anticipated resolution and discuss the prospects for dynamic transmission electron microscopy of organic samples in environmental cells. Recent results making use of nano-fluidic cell technology developed in house will be presented as an outlook: DNA – nanoparticle multimers have been studied regarding their dynamic in liquid and stability under electron irradiation during TEM imaging in a conventional 200 keV microscope [2]. First experiments towards imaging dynamics in living cancer cells under environmental conditions in the TEM will be discussed [3].

 

 References

[1]     S. Manz et al. Faraday Discuss. 2015, 177, 467-491.

[2]     S. Keskin et al., J. Phys. Chem. Lett. 2015, 6, 4487−4492.

[3]     S. Besztejan et al., submitted, 2016.

Stephanie MANZ (Hamburg, GERMANY), Sercan KESKIN, Stephanie BESZTEJAN, Benno ZEITLER, Sana AZIM, Guenther KASSIER, Robert BUECKER, Deybith VENEGAS-ROJAS, Svenja RIEKEBERG, Dongfang ZHANG, Albert CASANDRUC, Rolf A. LOCH, Yinpeng ZHONG, Hossein DELSIM-HASHEMI, Sascha W. EPP, Klaus FLOETTMANN, Hoc Khiem TRIEU, Andrea RENTMEISTER, R. J. Dwayne MILLER

13:45-15:45
Added to your list of favorites
Deleted from your list of favorites

MS5-III
MS5: Energy-related materials
SLOT III - Solar cells, thermoelectrics, storage

MS5: Energy-related materials
SLOT III - Solar cells, thermoelectrics, storage

Chairmen: Wolfgang JÄGER (Kiel, GERMANY), Joachim MAYER (Aachen, GERMANY), Philippe MOREAU (Associate Professor) (Nantes, FRANCE)
13:45 - 14:15 High Resolution and 3-D STEM of Energy Materials. Ilke ARSLAN (Richland, USA)
Invited
14:15 - 14:30 #6370 - MS05-OP270 In situ TEM analysis of structural changes in metal-halide perovskite solar cells under electrical bias.
In situ TEM analysis of structural changes in metal-halide perovskite solar cells under electrical bias.

Organic-inorganic metal-halide perovskite solar cells are emerging as a promising photovoltaic technology to harvest solar energy, with latest efficiencies now surpassing 22%1 - an impressive increase from the first reported value of 3% in 2009.2 In addition to low manufacturing costs, the optical properties of such cells can be tailored to form efficient tandems when combined with high-efficiency silicon solar cells.3 A typical perovskite cell structure as investigated here is based on a methylammonium lead trihalide absorber (MAPbI3) that is placed between hole- (Spiro-OMeTAD) and electron-selective contacts (a fullerene-based material).

 

While new record efficiencies are frequently reported, the commercial application of this solar cell technology remains hindered by issues related to thermal and operational stability. Different mechanisms that are still debated modify cell properties with time, temperature, illumination and general operating conditions.4 In order to correlate applied voltage (V) and resulting current (I) to changes in active layer chemistry and structure on the nanometre scale, we performed both ex situ and in situ transmission electron microscopy (TEM) experiments, involving (scanning) TEM (STEM) imaging, selected-area electron diffraction, energy-dispersive X-ray spectroscopy and electron energy-loss spectroscopy. Samples were prepared by focused ion beam (FIB) milling, with exposure to air during transfer to the TEM minimised to <5 minutes to reduce any degradation of MAPbI3.

 

First, the effects of exposure to air and electron beam irradiation were assessed in relation to FIB final thinning parameters. Once adequate sample preparation and observation conditions were identified, changes in morphology during cell characterisation were assessed ex situ by comparing lamellae extracted from as-manufactured and tested cells and then in situ by contacting FIB-prepared samples to a microelectromechanical systems (MEMS) chip mounted in a TEM specimen holder5 (Fig. 1a). Cell manufacturing parameters led to iodine diffusion into the hole collector, with the width of this diffused layer remaining constant during I-V characterisation. Similarly to ex situ experiments, the MAPbI3/Spiro interface was observed to delaminate during in situ electrical measurements, resulting in the presence of a ~5 nm Pb-rich layer on the hole-transparent-layer side (Figs. 1b-c). In addition, PbI2 nanoparticles were observed to nucleate within the MAPbI3 layer at the hole-collector interface and at the positions of structural defects (Figs. 1b-d).

 

Overall, the active MAPbI3 layer was observed to be sensitive to sample preparation, exposure to air, observation conditions and I-V stimulus, resulting in the need for great care to deconvolute each effect. Different mechanisms that may all contribute to the decrease in efficiency of the cell were identified both ex situ and in situ, including ionic migration, PbI2 formation and local delamination of interfaces.

 

Acknowledgments

Financial support is gratefully acknowledged from the Swiss National Science Foundation (SNSF) Sinergia project DisCO and the European Union Seventh Framework Program under Grant Agreement 312483 – ESTEEM2 (Integrated Infrastructure Initiative – I3). The authors also wish to thank Max Kruth, Doris Merteens and Vadim Migunov.

 

1         NREL research cell efficiency records, 2016, http://www.nrel.gov/ncpv/.

2         A. Kojima et al., J. Am. Chem. Soc., 2009, 131, 6050–6051.

3         J. Werner et al., J. Phys. Chem. Lett., 2016, 7, 161–166.

4         S. D. Stranks et al., Nat. Nanotechnol., 2015, 10, 391–402.

5         M. Duchamp et al., Microsc. Microanal., 2014, 20, 1638–1645.

Quentin JEANGROS (Neuchâtel, SWITZERLAND), Martial DUCHAMP, Jérémie WERNER, Rafal DUNIN-BORKOWSKI, Björn NIESEN, Christophe BALLIF, Aïcha HESSLER-WYSER
14:30 - 14:45 #5058 - MS05-OP261 Stabilization of the cubic perovskite BSCF phase by Y-doping.
Stabilization of the cubic perovskite BSCF phase by Y-doping.

Among complex oxides Ba0.5Sr0.5Co0.8Fe0.2O3-d (BSCF) exhibits excellent oxygen permeability due to its high oxygen non-stoichiometry. However, the favoured cubic perovskite phase tends to partly decompose in the desired operation temperature between 700 and 900 °C. Various secondary phases with high Co-content and low O-non-stoichiometry are formed close to grain boundaries [1]. The decomposition can be correlated with thermally activated O-vacancies. By lowering the temperature (from e.g. 1000 °C), the O-vacancy concentration decreases forcing the multivalent transition metal atoms (mainly Co) to increase their valence state to maintain charge neutrality. This leads to the collapse of the cubic phase because the tolerance factor shifts into the hexagonal regime [2]. The formation of secondary phases leads to a substantial reduction in the O-permeation [2].

To overcome this issue, doping was suggested which proved to be beneficial for several different cations (e.g. Zr, Y) [3,4]. Among these candidates, Y was chosen due to its monovalent character and large ionic radius which are both beneficial for stabilizing the cubic phase. Transmission/scanning electron microscopy (TEM/SEM) investigations were carried out on 1% (BSCF1Y) to 10% Y (BSCF10Y) B-site doped samples to analyse the phase constitution. O-permeation measurements were performed to assess the long-time permeability. Due to the intermediate ionic radius of Y the lattice position was checked using atom location by channeling enhanced microanalysis (ALCHEMI) [5].

BSCF tends to form a variety of secondary phases. Especially at high temperatures Co tends to diffuse out of the cubic lattice. This happens during the sintering process at high temperatures of 1050…1150 °C when CoO grains form close to grain boundaries. At lower temperatures partial phase decomposition to Ban+1ConO3n+3(Co8O8) (n ≥ 2, BCO), Co3O4 and Ba0.5+xSr0.5-xCoO3 (hexagonal) phases was demonstrated. Valence-state analysis of Co employing the Co-L2,3 white-line distance method [7] revealed that secondary phases with higher Co-valence state tend to be more stable at lower temperatures. Permeation measurements of BSCF10Y show that the degradation of O-permeation is reduced compared to undoped BSCF. This can be understood by the suppression of secondary phase formation in BSCF10Y at temperatures ≥800 °C (cf. Fig. 1 and Fig. 2). Secondary phases like BCO (cf. Fig. 3) and CoxOy are completely supressed by ≥ 3 at% Y-doping. The surfaces of permeation pellets show only minor amounts of BaSO4 which can be attributed to sulphur impurities in the feed gas. At lower temperatures (~700 °C) small volume fractions of the hexagonal phase are formed even in BSCF10Y (cf. Fig. 4). However, the volume fraction was negligible compared to the amount of secondary phases in undoped BSCF revealing a stabilizing effect on the cubic BSCF phase even at 700 °C.

ALCHEMI experiments showed unintended partial Y-occupation of up to 55 % on the A-site [8]. Therefore, Y doping is expected to generate B-site vacancies. These can contribute to the increased stability of the cubic phase because BSCF seems to tend towards B-site deficiency. To verify this hypothesis BSCF with 5% B-site deficiency was intentionally prepared which indeed showed less Co-outdiffusion and less pronounced secondary phase formation.

 

References

[1] P. Müller et al., Chem. Mater. 25, 564–573 (2013).

[2] S. Švarcová et al., SSI 178, 1787–1791 (2008).

[3] S. Yakovlev et al., Appl. Phys. Lett. 96, 254101 (2010).

[4] P. Haworth et al., Sep. Purif. Technol. 81, 88–93 (2011).

[5] J.C.H. Spence & J. Taftø, J. Microsc. 130, 147–154 (1983).

[6] M. Arnold et al., J. Membrane. Sci. 293, 44–52 (2007).

[7] P. Müller et al., Microsc. Microanal. 19, 1595–1605 (2013).

[8] M. Meffert et al., Microsc. Microanal. 22, 113-121 (2016).

[9] Financial support from the German Science Foundation (DFG) is gratefully acknowledged.

Matthias MEFFERT (Karlsruhe, GERMANY), Lana-Simone UNGER, Stefan BAUMANN, Christian NIEDRIG, Heike STÖRMER, Wolfgang MENESKLOU, Stefan WAGNER, Wilhelm Albert MEULENBERG, Ellen IVERS-TIFFÉE, Dagmar GERTHSEN
14:45 - 15:00 #6702 - MS05-OP274 Elemental mapping of perovskite solar cells using STEM and multivariate analysis.
Elemental mapping of perovskite solar cells using STEM and multivariate analysis.

Over the last few years, the interest in perovskite based solar cells has boomed, due to a surprisingly fast increase in terms of their efficiency that has now reached values comparable with established photovoltaic technologies. Nevertheless, the understanding of the optoelectronic properties of such nanostructured materials is still an open problem and issues related to their stability and degradation pathways represent the current hot topic in this research area.

Organic-inorganic solar cells present a complex composition as well as a composite structure that are strongly related to device fabrication. In this work four processing methods of the organic-inorganic halide perovskite have been investigated, varying the deposition method (single step or double step) and the atmospheres in which the synthesis has been carried out1. We compared interface quality, morphology, chemical composition and efficiency of the resulting cells. A fluorine doped tin oxide (FTO) glass layer was coated first by a compact (hole blocking) TiO2 layer and then by a nanoporous TiO2 layer. The TiO2 scaffold was infiltrated and capped by a methyl-ammonium lead iodide. Spiro-MeOTAD, acting as hole transport layer was spin coated on the perovskite layer; Au contacts were deposited on top. The devices were analysed using several complementary characterisation techniques: Scanning Transmission Electron Miscroscopy (STEM) used in conjunction with EDX analysis, time of flight secondary ion mass spectrometry (ToF-SIMS) and X-rays photoelectron spectroscopy (XPS). In particular, the use of FIB specimen preparation, combined with analytical transmission electron microscopy, represents a powerful and versatile tool for the characterization of devices based on hybrid composites with nano- and micro-scale structural and chemical features. EDX maps were then treated using multivariate analysis in order to optimise signal-to-noise ratio and obtain high quality EDX maps. The application of this method plays a key role in the analysis of data acquired with low electron doses to minimise specimen damage, and is unique in allowing the identification of different materials present in the sample as compounds rather than individual elements.

 

Having fully characterised the devices, we investigated the different degradation processes that affect the perovskite based solar cell. Air exposure and temperature were proven to be  responsible for the drastic reduction in device performance. STEM-EDX analysis was thus used in conjunction with the in situ heating, bringing the cells to 250 °C2. The main result was the direct observation of elemental migration, particularly evident in iodine maps, and the simultaneous formation of metallic Pb precipitates, resulting in the depletion of the initial perovskite region.

Lastly, we studied the changes in chemical composition and morphology after 2 months of air exposure in dark3. In this case the principal variations in local elemental composition and in morphology concerned respectively the migration of lead and iodine into the HTL layer towards the Au electrode, resulting in a severe degradation of the photoactive layer and the physical formation of bubbles in the Spiro-OMeTAD.

 

[1] F Matteocci, Y Busby, JJ Pireaux, G Divitini, S Cacovich, C Ducati, A. Di Carlo, Interface and composition analysis on perovskite solar cells, ACS applied materials & interfaces 7 (47), 26176-26183 (2015)

[2] G. Divitini, S. Cacovich, F. Matteocci, L. Cinà, A. Di Carlo, C. Ducati, In situ observation of heat-induced degradation of perovskite solar cells, Nature Energy, 15012 (2016)

[3] S. Cacovich, G. Divitini, C. Ireland, F. Matteocci, A. Di Carlo, C. Ducati, Study of ageing processes in perovskite solar cells using Scanning Transmission Electron Microscopy and multivariate analysis, Under review

Stefania CACOVICH (Cambridge, UK), Giorgio DIVITINI, Fabio MATTEOCCI, Yan BUSBY, Jean-Jacques PIREAUX, Aldo DI CARLO, Caterina DUCATI
15:00 - 15:15 #5985 - MS05-OP267 Investigating Cu diffusion in CdTe solar cells via aberration-corrected STEM: Cu2-xTe precipitates at CdTe twins and the CdTe/CdS interface.
Investigating Cu diffusion in CdTe solar cells via aberration-corrected STEM: Cu2-xTe precipitates at CdTe twins and the CdTe/CdS interface.

 CdTe is one of the most promising materials for thin film solar cells due to its near-optimum bandgap, high efficiency and low cost of fabrication. In order to reduce the contact resistance, it is common to deposit a thin Cu layer at the back of the cell. However, Cu is found to diffuse easily into the p-type CdTe layer and even to the n-type CdS layer. This diffusion causes a problem because too much Cu diffusion has been shown to cause performance degradation [1-2]. Therefore it is important to figure out the pathways of Cu diffusion, and the configurations that Cu stabilizes inside the CdTe solar cells.

 In this research, we employed aberration-corrected Scanning Transmission Electron Microscopy (STEM) combined with Electron Energy Loss Spectroscopy (EELS) or Energy Dispersive X-Ray Spectroscopy (EDX) to probe Cu diffusion in CdTe solar cells. We first investigated the cells with a typical Cu-annealing process at 150° C for 45 min in dry air under ambient pressure, but did not find very much Cu diffusion. Therefore we extended the annealing at 150° C for 8 hours in vacuum using our sample-baking system. The vacuum annealed samples show clear Cu diffusion traces along some of the grain boundaries (GBs). However not all the GBs became Cu-enriched, indicating that Cu diffusion may have a preference for different GBs (Fig.1). Moreover, STEM-EDX reveals Cu-rich precipitates embedded inside high-density twin boundaries in CdTe (Fig.2). The Cd concentration decreases while Te does not change on the precipitates, and no other elements were found, indicating the precipitates are most likely to be CuxTey. The CuxTey precipitates exist all the way along the high density of twin boundaries from the contact layer to the CdS layer, demonstrating another pathway for Cu diffusion.

 Similar CuxTey precipitates were found at the CdTe/CdS interface. A precipitate located at the edge of the STEM specimen without overlapping CdTe or CdS was used to quantify the ratio of Cu to Te via both EELS and Z-contrast image intensity. Both results show that the ratio of Cu to Te is slightly lower than 2:1, therefore the precipitate can be described as Cu2-xTe. Moreover, the atomic structure of the precipitates show variations from the normal CuTe structure, for example formation of 2×2 supercells (Fig.3). More detailed results will be shown and the photovoltaic behavior of the Cu2-xTe precipitates will also be discussed.

 

References:

[1] K.D.Dobson, et al, Stability of CdTe/CdS thin-film solar cells, Solar Energy Materials & Solar Cells, 62, 295 (2000)

[2] S.H.Demtsu, et al, Cu-related recombinations in CdS/CdTe solar cells, thin solid films, 516, 2251 (2008)

[3] This research was supported by the US DOE, Office of Energy Efficiency and Renewable Energy (F-PACE, DE-FOA-0000492), (CL, NP, YFY, SJP), the Office of DOE-BES, Materials Science and Engineering Division (ARL). The EDX Research was sponsored in part by the UK EPSRC through the UK National Facility for Aberration-Corrected STEM (SuperSTEM) (TJP, SJH).  

Chen LI (Vienna, AUSTRIA), Timothy J. PENNYCOOK, Sarah HAIGH, Andrew LUPINI, Naba PAUDEL, Yanfa YAN, Stephen PENNYCOOK
15:15 - 15:30 #6829 - MS05-OP275 Investigation of the heterojunction morphology in donor/acceptor stacks by energy filtered transmission electron microscopy.
Investigation of the heterojunction morphology in donor/acceptor stacks by energy filtered transmission electron microscopy.

One challenge in the development of efficient organic photovoltaic devices (OPVs) is the optimization of the donor/acceptor(D/A) interface morphology, as this has tremendous impact on the electrical transport properties. In an optimized absorber layer morphology, excitons should be generated always in close proximity to the D/A interface to reach it within their lifetime, where theydissociate into free charge carriers [1-3]. In OPVs however the free charge carriers exhibit a small mobility which requires short pathways towards the electrodes to prevent recombination. Substrate temperature during deposition as well as surface chemistry are important process parameters [4-7]. In this respect alternating D/A layers are promising candidates for an optimized active layer morphology [8-10].

In this contribution we present results obtained from a stack assembly of three Zn-Phtalocyanine/C60 layer (3 nm/3 nm nominal thickness) pairs (for further details of the preparation see [11]). The complete OPV device is shown in Fig. 1. The morphology of the individual D/A layers was investigated after each deposition step of ZnPc or C60, respectively. For this, energy filtered electron transmission microscopy in a Zeiss LIBRA 200 FE operated at 80 kV was used. In order to be able to investigate the same area after each deposition step, a special finder TEM grid, coated with a graphene monolayer was used. Zero-loss filtered images were acquired with an energy filter slit width of 10 eV. In Fig. 2 we show images which where obtained after the first, second and third deposition, respectively (deposition at room temperature (RT)). The first image (Fig. 2a) shows that the ZnPc film is not closed, but islands have formed. The same can be observed for the successive C60 layer (Fig. 2b). However, the latter agglomeration is stronger, thus larger clusters are formed, indicating a higher surface mobility of C60. In the further course of the deposition steps (not shown) the contrast between image features of the first two layers decreases. This indicates that the successive deposition steps lead to closed films. Such a morphology is unfavorable as there are limited percolation paths from the different interfaces towards the respective electrodes. The morphology in the layer stack changes substantially when the substrate temperature is increased to 80 °C. Figure 3 shows the corresponding morphologies after the 1st, 2nd and 3rd deposition step, respectively. In contrast to the experiment at RT agglomeration is observed in all following layers. In order to investigate the inter-layer links the morphology in layer n was visualized by calculating the difference image dn = Imagen – Imagen-1. The result for the first six layers is also depicted in Fig. 3. We observe a high porosity and a well established crosslinking between the ZnPc on the one hand and the C60 layers on the other. These morphology results in an improved efficiency of η=2.1 % compared to η=0.5 % for the device deposited at RT.

References:

[1] B. Johnev, M. Vogel, K. Fostiropoulos, B. Mertesacker, M. Rusu, M.-C. Lux-Steiner, A. Weidinger, Thin Solid Films, vol. 488, no. 1, pp. 270 – 273 (2005).

[2] S. Senthilarasu, S. Velumani, R. Sathyamoorthy, A. Subbarayan, J.A. Ascencio, G. Canizal, P.J. Sebastian, J.A. Chavez, and R. Perez, Appl. Phys. A, 77:383 – 389 (2003).

[3] B. P. Rand, D. Cheyns, K. Vasseur, N. C. Giebink, S. Mothy, Y. Yi, V. Coropceanu, D. Beljonne, J. Cornil, J. L. Bredas, J. Genoe, Adv. Func. Mater., 22:2987–2995 (2012).

[4] K. Fostiropoulos, W. Schindler, Phys. Status Solidi B246, 2840 – 2843 (2009).

[5] A. F. Bartelt, C. Strothkämper, W. Schindler, K. Fostiropoulos, R. Eichberger, Appl. Phys. Lett. 99, 143304 (2011).

[6] S. Pfuetzner, J. Meiss, A. Petrich, M. Riebe, K. Leo, Appl. Phys. Lett. 94, 253303 (2009).

[7] P. Peumans, A. Yakimov, and S.R. Forrest, J. Appl. Phys., 93(7), 3693–3723 (2003).

[8] J. Xue, B. P. Rand, S. Uchida, and S.R. Forrest, J. App. Phys., 98(12):124903 (2005).

[9] B. P. Rand, J. Xue, S. Uchida, and S. R. Forrest, J. Appl. Phys., 98, 124902 (2005).

[10] Z. R. Hong, B. Maenning, R. Lessmann, M. Pfeiffer, and K. Leo, Appl. Phys. Lett. 90, 203505 (2007).

[11] G. Chouliaras, W. Schindler, M. Wollgarten, M. Ch. Lux-Steiner, K. Fostiropoulos, submitted.

Georgios CHOULIARAS, Markus WOLLGARTEN (Berlin, GERMANY), Wolfram SCHINDLER, Konstantinos FOSTIROPOULOS
15:30 - 15:45 #5888 - MS05-OP266 Design and characterization of mesopores in photocatalytically active oxynitride single crystals using structural and chemical TEM analysis.
Design and characterization of mesopores in photocatalytically active oxynitride single crystals using structural and chemical TEM analysis.

Mesoporous single crystals have been a matter of intense discussion in the last years in the context of solar energy harvesting, since it is expected that this material class can contribute significantly to the improved design of highly efficient solar energy conversion devices [1]. Especially for photocatalytic or photoelectrochemical hydrogen generation high surface area and good charge-transport properties are key features to enhanced device performance [2, 3]. Good conductivity is usually obtained in large defect free structures such as single crystalline materials, where consequently the surface area is small. High surface area, however, is obtained best by porous agglomerates of nanoparticles, where the conductivity is low because of multiple grain boundaries. The possibility to achieve performance improvement by combining both concepts has been demonstrated via the fabrication of large single crystals on the micrometer scale with a mesoporous structure using a template based approach [4]. In comparison to nanocrystalline materials, the improved charge carrier conductivity has been shown in this material class as well as its competitive surface area. Moreover, other template-free synthesis routes are known, where porous structures in solids are formed spontaneously. One example is the solid-gas phase reaction carried out for the synthesis of oxynitrides, i.e. thermal ammonolysis [5]. However, the control of pore size and density requires a detailed understanding of the reaction mechanisms during synthesis.  

Some of these oxynitride materials i.e. LaTiO2N (LTON) or LaTaON2 are photocatalytically active [6-8]. The main characterization techniques used to evaluate the pore quantity and quality have been powder techniques (for example physisorption), giving information about the open porosity, and qualitative scanning electron microscopy (SEM) and transmission electron microscopy (TEM), which suggested that open and closed pores are formed [6-8]. However, little is known about the size and shape distribution especially of the closed porosity or about the pore formation process.

In this contribution we will focus on microscopic pore characterization of LTON as a function of the synthesis method by combining several TEM techniques (Figure 1, 2) [9,10]. The pores themselves were explored mainly by electron tomography and by scanning TEM (STEM) with an high angle annular dark field (HAADF) detector, while the crystallinity was investigated using a combination of high resolution transmission electron microscopy (HREM), selected area diffraction (SAD) and nanobeam diffraction. The local chemical composition was studied by electron energy loss spectroscopy. With the improved understanding of the pore formation mechanism in LTON we enabled porosity tuning in large oxynitride single crystals leading to enhanced performance in photocatalytic and photoelectrochemical water-splitting.

[1]          C. Ducati, Nature 495 (2013) 180-181.

[2]          T. Hisatomi, J. Kubota and K. Domen, Chem. Soc. Rev., 43 (2014) 7520-7535.

[3]          A. Kay, I. Cesar and M. Grätzel, J. Am. Chem. Soc., 128 (2006) 15714-15721.

[4]          E.J.W. Crossland, N. Noel, V. Sivaram, T. Leijtens, J.A. Alexander-Webber, and H.J. Snaith, Nature,

              495 (2013) 215-219.

[5]          S.G. Ebbinghaus, H.-P. Abicht, R. Dronskowski, T. Müller, A. Reller and A. Weidenkaff,

              Prog. Solid State Chem., 37 (2009) 173-205.

[6]          A.E. Maegli, S. Pokrant, T. Hisatomi, M. Trottmann, K. Domen, and A. Weidenkaff, J. Phys. Chem. C,

              118 (2014) 16344-16351.

[7]          T. Takata, C. Pan and K. Domen, ChemElectroChem, 3 (2016) 31-37.

[8]          N.-Y. Park and Y.-I. Kim, J. Mater. Sci., 47 (2012) 5333-5340.

[9]          S. Pokrant, M.C. Cheynet, S. Irsen, A.E. Maegli and R. Erni, J. Phys. Chem. C, 118 (2014)

              20940-20947.

[10]        S. Pokrant, S. Dilger and S. Landmann, J. Mater. Sci., (2016) DOI: http://dx.doi.org/10.1557/jmr.2016.9

Acknowledgements: The authors thank the SNSF for the PrecoR grant 20PC21_155667.

Simone POKRANT (Dübendorf, SWITZERLAND), Stefan DILGER, Steve LANDSMANN

13:45-15:45
Added to your list of favorites
Deleted from your list of favorites

IM7-II
IM7: Phase Microscopies
SLOT II

IM7: Phase Microscopies
SLOT II

Chairmen: David COOPER (Engineer) (Grenobles, FRANCE), Christoph KOCH (Professor) (Berlin, GERMANY)
13:45 - 14:15 #8361 - IM07-S52 A closer look at high-resolution electron holography.
A closer look at high-resolution electron holography.

In the recent years high-resolution off-axis electron holography made huge advancements. Electron microscopes with increased numbers of electron optical biprisms and electron lenses allow more flexible ray paths. Due to the combination of double biprism holography and higher magnification at the lower biprism, holograms in the high resolution regime are now recorded with negligible artifacts [1]. Furthermore, smart averaging schemes for hologram series allow to prolong the effective exposure time for a hologram to time scales, where the experimental errors are not dominated by the shot noise anymore [2].

 

One of the advantages of holography is the posterior correction of residual aberrations. However, the aberrations still have to be sufficiently known. By quantitative comparison of reconstructed wave functions with calculations the imaging parameters can be retrieved with sufficient precision allowing the reconstruction of aberration corrected exit-wave functions. Thus nowadays holograms with (sub-)angstrom resolution can be obtained on a regular basis.

 

In off-axis holography only the side-band channel of the information recorded in the hologram is reconstructed. As hologram series are recorded and the reconstructed wave function can be propagated to arbitrary defocus values anyway, also focal series of conventional images (center-band) can be obtained in the same measurement.

 

In the figures the reconstructed exit waves of a GaAs wedge, obtained by inline and off-axis holography are shown [3, 4]. While differences in the low-frequency reconstruction are expected, also significant differences in thicker specimen parts are found. The latter can be easily recognized on the left side of the linescan profiles, where the amplitudes signals at the columns positions exhibit different behaviours. There are several possible reasons for this discrepancy. The numerical inversion of the imaging process within the inline method becomes worse conditioned since the thicker parts exhibit stronger non-linear imaging. Thus the reconstruction algorithm might converge in a wrong minimum of the overall error figure. Also, any changes of the object during the acquisition of the series will exhibit different behaviour in the measurement. Furthermore, both channels (side-band and center-band) refer to different parts of the density matrix due to the different quantum-mechanical nature of the underlying interference experiment.

 

[1] F. Genz et al., Ultramicroscopy 147 (2014) 33

[2] T. Niermann et al., Micron 63 (2014) 28

[3] T. Niermann et al., J. Phys. D: Appl. Phys. 49 (2016) 194002

[4] The FEI TrueImage software packages was used for inline reconstruction.

Tore NIERMANN (Berlin, GERMANY), Michael LEHMANN
Invited
14:15 - 14:30 #6940 - IM07-OP137 Long-range focal series reconstruction in the TEM.
Long-range focal series reconstruction in the TEM.

Focal series wave reconstruction in the Transmission Electron Microscope (TEM) is a well established holographic technique employed in both the medium and atomic resolution regime to study electric, magnetic and strain fields in solids as well as atomic configurations at crystal defects or grain boundaries. Focal series reconstruction does not require an undisturbed reference wave as off-axis holography and may be conducted under relaxed partial coherence provided that the latter is well-behaved and well-known in advance [1]. Moreover, focal series holography may be considered as an instance of the more general quantum state tomography (see Fig. 1) that is successfully employed to study mixed (i.e., incoherent) quantum states of matter (e.g., atoms) and light [2].
These advantages are opposed by ambiguities in the reconstructed wave function, e.g., due to inconsistent and incomplete focal series data. In reality every focal series is inconsistent, e.g., due to the presence of partial coherence, shot and detector noise, as well as geometric and chromatic aberrations depending on the defocus. Similarly, every focal series is incomplete because of a limited number of foci, typically limited to the near field regime, and the restriction to isotropic foci, where astigmatic foci are necessary to provide a dataset allowing for an unabiguous reconstruction of an underlying wave function [3]. For instance, the problematic reconstruction of low spatial frequencies can be traced back to missing focal series data in the far field.
Here, we elaborate on focal series reconstruction from the perspective of quantum state tomography and use the obtained results to increase the scope of the technique in terms of convergence and uniqueness in particular for low spatial frequencies. Moreover, we explain a number of previous results by exploiting the above analogy, and open pathways to further improvements.
We particularly report on the recording, preprocessing, calibration and reconstruction of a long range focal series ranging from the near to the far field in a TEM. We calibrate the focal series, including the effective defocus and magnification, by a careful calibration of the proportionality between squared current and reziprocal focal length in a magnetic lens. We derive non-linear focal sampling schemes from the phase space analogy. Subsequently, we adapt a modified Gerchberg-Saxton algorithm to the long range focal series by exploiting the link to randomized Kaczmarz (ART) algorithm used in tomography [4]. We use different numerical propagation regimes in the near and far field to take into account the scaling of the wave function and overcome convergence problems by replacing the Kaczmarz iteration with the Landweber (SIRT) iteration as proposed by Allen et al.. [5]. To overcome remaining ambiguities in the reconstruction (e.g, pertaining to a different starting guess in the Gerchberg-Saxton algorithm) resulting from inconsistencies in combination with the non-convex nature of the set of wave functions possessing the same modulus, we discuss several additional constraints such imposed by the topology of the starting guess [6].
To illustrate the above reconstruction principles, we perform a case study on a higher-order vortex beam with topological charge (winding number) 3 truncated by a square aperture (Fig. 2). The beam possesses a non-trivial topology by design, which is nicely suited to discuss the impact of (implicit) topology constraints, rotation alignment as well as other issues.

[1] Koch, C. T., Micron, 2014, 63, 69-75
[2] Schleich, W. P., Quantum Optics in Phase Space, Wiley VCH, 2001
[3] Lubk, A. & Röder, F., Phys. Rev. A, 2015, 92, 033844
[4] Natterer, F., Wübbeling, F., Mathematical Methods in Image Reconstruction,SIAM, 2001
[5] Allen, L. J.; McBride, W.; O'Leary, N. L.,Oxley, M. P., Ultramicroscopy, 2004, 100, 91-104
[6] Martin, A. & Allen, L., Optics Communications, 2007, 277, 288-294
[7] Financial support by the DIP programme of the DFG is greatly acknowledged.

Axel LUBK (Dresden, GERMANY), Karin VOGEL, Daniel WOLF, Falk RÖDER, Laura CLARK, Jo VERBEECK
14:30 - 14:45 #5921 - IM07-OP130 Superresolution and depth sensitivity in HRTEM through structured illumination.
Superresolution and depth sensitivity in HRTEM through structured illumination.

In the weak phase object approximation the wave function in the objective lens’ back focal plane is a convolution of the illuminating wave, ψin, with the projected potential, V, multiplied with the coherent transfer function, H, i.e. ( ψin * ( 1 + σV ) ) H, with σ the interaction constant and * the convolution operator. Therefore, frequencies from above the cutoff-frequency in H’s objective aperture enter the image formation in aliased form, as is illustrated in Fig. 1a. In this abstract, fully dynamical multislice simulations are carried out to investigate if these higher frequencies can be disentangled and superresolution can be achieved.

Simulations of [112] Si are set up with FDES [1], the supercell is 5.12 nm by 5.12 nm wide and is 5 nm thick on the left side and 2.5 nm on the right-hand side. The in-plane sampling distance and the slice thicknesses were set to 0.01 nm. Si in the [112] direction displays dumbbells that are 0.075 nm apart, corresponding to a spatial frequency of 13.3 nm-1 or 33.4 mrad at an acceleration voltage of 200 kV. The imaging is done with illumination that is structured to have random-but-known phases and that is bandwidth limited to 20 mrad, see Fig. 1b and Fig. 2a. The objective lens behind the sample is set to Scherzer conditions with an objective aperture of 20 mrad, a defocus of -18.8 nm and a spherical aberration of 94.0 µm, thus yielding a point resolution of 0.125 nm that is clearly insufficient for resolving the dumbbells. Eight images are recorded with the simulation shifted by 0.27 nm between consecutive recordings; the first image is shown in Fig. 2b, the others look similar.

A three-dimensional reconstruction was set up in IDES [2,3] with three slices 2.5 nm apart and a pixel size of 0.01 nm; amplitude and phase of the impinging wave and the settings of the objective lens were assumed known. Multiple scattering was accounted for by propagating the wave between slices with the Fresnel propagator. A Polak-Ribière conjugate gradient search was used to minimize an error function that was chosen as the sum of absolute differences between measurement and corresponding data simulated from the current object estimate. Contrary to [2,3] no sparse regularization has been applied. In Fig. 3a the middle slice of the reconstruction is shown, displaying a clear separation of the Si-dumbbells. Furthermore, the lower slice displayed in Fig. 3b shows that depth sensitivity with a resolution of at least 2.5 nm is achieved as the right-hand side, which by construction does not contain any atoms in its lower half, indeed does not display any. [4]

[1] W. Van den Broek, et al. “FDES, a GPU-based multislice algorithm with increased efficiency of the computation of the projected potential.” Ultramicroscopy 158 (2015), pp. 89–97.
[2] W. Van den Broek and C.T. Koch. “Method for retrieval of the three-dimensional object potential by inversion of dynamical electron scattering.” Phys. Rev. Lett. 109 (2012), p. 245502.
[3] W. Van den Broek and C.T. Koch. “General framework for quantitative three-dimensional reconstruction from arbitrary detection geometries in TEM.” Phys. Rev. B 87 (2013), p. 184108.
[4] The Carl Zeiss Foundation is gratefully acknowledged by all authors. C.T. Koch also acknowledges the DFG (KO 2911/7-1).

Wouter VAN DEN BROEK (Berlin, GERMANY), Christoph T. KOCH
14:45 - 15:00 #5023 - IM07-OP126 Towards understanding of charging effects of thin-film phase plates.
Towards understanding of charging effects of thin-film phase plates.

In the past few years, physical phase plates (PP) have become a viable tool to enhance the contrast of weak-phase objects in transmission electron microscopy (TEM). Here we focus on thin-film PPs where the mean inner potential is exploited to impose a phase shift on electrons propagating through the PP [1]. The application of thin-film PPs is hampered by deviations of the phase shift from its desired value which occur due to charging of the thin film. Our experimental approach to overcome charging of thin-film PPs was using the metallic glass alloy Pd77.5Cu6.0Si16.5 (PCS) with a high specific conductivity of 1.18×106 S/m [2] as PP material. However, Hilbert PPs fabricated from thin PCS films nevertheless show pronounced distortions of the Thon-ring system during illumination with 200 keV electrons. These observations initiated the development of a theoretical model to obtain an improved understanding of charging, which is presented in this work.

 

Charging is described by assuming a charge-dipole layer to be present at the PP. A possible source for such a dipole layer could be an insulating contamination layer on top of the PCS film in the illuminated PP region which could capture and fix low-energy secondary electrons generated by the primary electrons in the PCS film. Together with its positive mirror charge in the grounded electrically conducting PCS film this fixed charge would form a dipole layer. The dipole strength is assumed to be proportional to the current density distribution in the back focal plane which can be qualitatively obtained from a diffraction pattern. The proportionality factor for the dipole strength is a fit parameter denoted as phase mask amplitude in the following.

 

As a test of our model we compare power spectra obtained from an experimental image of an amorphous carbon (aC) thin-film test object with a simulation based on our model. The experimental reference image was obtained by using a PCS film-based Hilbert PP installed in the back focal plane of a Philips CM 200 FEG/ST transmission electron microscope. In the simulation we assume the Hilbert PP to be illuminated by the current density distribution given by the diffraction pattern and calculate the phase shift in the back focal plane. The resulting dipole strength distribution is then fed into an image simulation procedure which yields simulated power spectra as a function of the phase mask amplitude. The latter is optimized by a comparison with the experimental reference. The experimental reference and calculated power spectra are compared pixel by pixel. The sum of all pixel comparisons belonging to one pair of power spectra serves as measure of agreement, which is plotted for different phase mask amplitudes in Figure 1 for the region below the cut-on frequency. Best agreement of almost 75% is obtained. This is remarkable taking into account that experimental and simulated spectra (based on noisy input for the illumination data) contain noise.

 

Figure 2 shows a montage of a simulated and experimental power spectrum. Two regions, below and above the cut-on frequency, can be distinguished in the power spectra. The cut-on frequency is given by the distance between the PP edge and the zero-order beam and is marked by vertical white lines. Note the good agreement between experimental and simulated spectra, especially below the cut-on frequency.

 

Overall our method seems to be a promising approach to analyze and explain phase shift distortions due to charging in thin-film PP applications.

 

[1] R. Danev and K. Nagayama, J. Phys. Soc. Jpn. 73 (2004), p. 2718.

[2] B. Chelluri and R. Kirchheim, J. Non-cryst. Solids 54 (1983), p. 107.

[3] Financial support by the Deutsche Forschungsgemeinschaft (DFG).

Roland JANZEN (Karlsruhe, GERMANY), Jonas SCHUNDELMEIER, Simon HETTLER, Manuel DRIES, Dagmar GERTHSEN
15:00 - 15:15 #6017 - IM07-OP132 Quantitative comparison of phase contrast imaging in conventional TEM focal series and STEM ptychography.
Quantitative comparison of phase contrast imaging in conventional TEM focal series and STEM ptychography.

 In phase contrast imaging, three-dimensional, quantitative information about the specimen is encoded in the object wave function, which results from the scattering of the electron wave with the specimen potential. For weak scattering materials, such as a mono-atomic layer of graphene, the phase of the object wave contains all the structural information. However, this phase information is lost in the image recording process. In order to recover the phase, a variety of numerical reconstruction methods are available, including off-axis electron holography [1], focal series reconstruction (FSR) [2] and ptychography [3]. To understand quantitatively materials’ properties, matching of experimental phases to simulations is required. In practice, the quantitative information that is obtained from the experimental object wave is often in disagreement with simulations, even for the simple case of a mono-atomic layer of graphene. This disagreement is a phase mismatch, that resembles the contrast mismatch, or Stobbs factor, [4] found between images and simulations.

 In this contribution, we focus on a comparison between conventional focal series phase restoration in transmission electron microscopy (TEM) and ptychographic phase restoration in scanning transmission electron microscopy (STEM) for the simple case of graphene. These techniques provide two independent measurements of the phase of a monolayer, which we subsequently compare to explore the physical meaning of the restored object wave phase. Figure 1 shows the restored phase from a conventional focal series reconstruction (a) and from a ptychography reconstruction (b). The focal series was recorded using an aberration corrected JEOL 2200MCO equipped with an in-column Omega-type energy filter, while the ptychography data set was acquired in a probe corrected JEOL ARM200CF fitted with a direct electron pixelated detector from PNDetector. The detector has an array of 264x264 pixels and can achieve a speed of up to 20,000 fps through binning/windowing. The focal series and ptycographic reconstructions are based on Wiener filter [5] and Wigner distribution deconvolution [6] algorithms, respectively. Following both restorations, the range of recovered phases is compared for the two methods. The theoretical phase of the object function is also determined for both cases by performing multislice frozen phonon calculations and reproducing, step by step the experimental restoration procedures. The preliminary results of these calculations are shown in Figure 2 for the TEM case. Without introducing any unknown fitting parameters in the simulations, the mismatch between the calculated and experimental phases is found to be 1.7.

 The experimental comparison with the ptycographic simulations, and those for FSR recorded at elevated temperatures, will be further discussed to show the influence of thermal motion on the restored object wave, as well as the effects of using zero-loss filtered images to exclude inelastic scattering contributions [7].

 

References: 

[1] M. Lehmann et al., Ultramicroscopy 54 (1994) 335 – 344.

[2] W. M. J. Coene et al., Ultramicroscopy 64 (1996) 109 – 135.

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

[4] M. J. Hÿtch and W.M. Stobbs, Ultramicroscopy 53 (1994) 191 – 203.

[5] A. I. Kirkland et al. Ultramicroscopy 57 (1995) 355 – 374.

[6] J. M. Rodenburg and R. H. T. Bates, Philosophical Transactions of the Royal Society A: Mathematical, Physical and Engineering Sciences, 339 (1992) 521-553.

[7] A. Howie, Ultramicroscopy 98 (2004) 73 – 79.

[8] The authors acknowledge funding from the European Union Seventh Framework Programme under Grant Agreement 312483-ESTEEM2 (Integrated Infrastructure Initiative-I3) and from the EPSRC under grant number EP/M010708/1.

Emanuela LIBERTI (Oxford, UK), Hao YANG, Gerardo MARTINEZ, Peter NELLIST, Angus KIRKLAND
15:15 - 15:30 #6136 - IM07-OP133 Study on the robustness of electron ptychography for phase imaging in the STEM using fast pixelated detectors.
Study on the robustness of electron ptychography for phase imaging in the STEM using fast pixelated detectors.

The use of electron ptychography methods for phase imaging in the scanning transmission electron microscope (STEM) have gained renewed interest due to the recent developments of fast pixelated detectors. One of those being the pnCCD (S)TEM camera [1] developed by PNSensor and PNDetector, it allows the recording of a 4D-dataset that consists of the full convergent beam electron diffraction pattern for each probe position that is scanned over the sample. This dataset contains all the scattering information generated by the electron sample interaction in the STEM experiment. The technique has allowed improved resolution beyond conventional limits [2] and phase contrast imaging at atomic resolution of different materials [3, 4]. In combination with other STEM techniques, it enables imaging of heavy and light elements of radiation-sensitive materials at the atomic scale [4].

In this work, we present an analysis and discussion of two different phase reconstruction algorithms, such as the Single Side Band (SSB) [3, 5, 6] and the Wigner-Distribution Deconvolution (WDD) [7] methods, in terms of their robustness to dynamical effects. The main difference of both methods is that the phase difference retrieved by the SSB method is still affected by the probe aberrations, meanwhile the WDD method allows to correct for this [4, 7].  However, in both cases a multiplicative interaction between the specimen function and the electron wave is assumed, which complicates the interpretation of the reconstructed wave when dynamical effects start to play a major role. Figures 1 and 2 show experimental results of retrieving the modulus (a) and the phase difference (b) of a gold nanoparticle measured with the pnCCD (S)TEM camera and using the discussed ptychography methods. It can be observed that by correcting the probe aberrations [3] while using the WDD method, the quality of the reconstructed image is dramatically improved. Moreover, a ring-like shape of the reconstructed phase difference for some atomic columns is observed. This behavior is studied in detail by image simulations (c) and discussed in terms of how dynamical effects influence the reconstruction algorithm. Furthermore, we explore the comparison of this method with other phase imaging techniques.

[1] H. Ryll, M. Simson, R. Hartmann, P.Holl, M. Huth, S. Ihle, Y. Kondo, P. Kotula, A. Liebel, K. Müller-Caspary, A. Rosenauer, R. Sagawa, J. Schmidt, H. Soltau, L. Strüder, manuscript accepted at Journal of Instrumentation

[2] P.D. Nellist, B.C. McCallum and J.M. Rodenburg, Nature 374 (1995) 630-632.

[3] T. J. Pennycook, A. R. Lupini, H. Yang, M. F. Murfitt, L. Jones, P. D. Nellist, Ultramicroscopy 151 (2015) 160 – 167.

[4] H Yang, R.N. Rutte, L. Jones, M. Simson, R. Sagawa, H. Ryll, M. Huth, T.J. Pennycook, H. Soltau, Y. Kondo, B.D. Davis, P.D. Nellist,  manuscript submitted.

[5] J. M. Rodenburg, B. C. McCallum, P. D. Nellist, Ultramicroscopy 48 (1993) 304 – 314.

[6] H. Yang, T. J. Pennycook, P. D. Nellist, Ultramicroscopy 151 (2015) 232 – 239.

[7] J. M. Rodenburg and R. H. Bates, Phil. Trans. R. Soc. Lond. A. 339 (1992) 521 – 553.

 

Acknowledgement

The research leading to these results has received funding from the EPSRC (EP/M010708/1).

Gerardo T MARTINEZ (Oxford, UK), Hao YANG, Lewys JONES, Martin SIMSON, Martin HUTH, Heike SOLTAU, Lothar STRÜDER, Ryuusuke SAGAWA, Yukihito KONDO, Peter D NELLIST
15:30 - 15:45 #5952 - IM07-OP131 Zorro: multi-reference dose-fractionated image registration.
Zorro: multi-reference dose-fractionated image registration.

    The technique of single-particle analysis (SPA) in cryo-TEM has recently been revolutionized by the simultaneous introduction of direct electron detectors (DEDs) and maximum likelihood, reference-free reconstruction tools such as Relion [S. Scheres] and Bayesian refinement tools such as Frealign [N. Grigorieff]. CMOS DEDs have much better DQE than scintillator-coupled CCDs and fast read-out permits collection of dose-fractionated image stacks which can be effectively drift-corrected. SPA is essentially an ensemble-average in-line holography technique.  Particles (i.e. proteins) are embedded in a vitreous ice matrix, such that they have random projections. Many micrographs are recorded at various defocus values (typically 1.0 – 2.5 µm). With sufficient particle projections recorded (typically 50k-500k), it is possible to estimate orientations of the particles by maximum likelihood iterative refinement and reconstruct the 3-D volume via back-projection using the central projection theorem. 

    Proteins are highly radiation sensitive (critical dose ~20 e-2 = 0.03 C/cm2) therefore useful per-frame dose is in the range 0.5 – 2 e-2. Ice films are viscous and thought to be semi-insulating under electron illumination. High, charged-driven drift rates are observed (x10 compared to hard specimens on carbon). Radiation damage and non-rigid particle motion results in rapid decay of correlation information amongst frames. In addition, DED’s gain reference change quickly due to radiation damage, such that correlated noise is often greater than the correlated signal.

    We will present Zorro, a new drift registration package, to be open-sourced (github.com/C-CINA/zorro) in the near-future after publication. Zorro takes a similar multiple reference approach as Motioncorr [Y. Cheng], in that it overcomes noise by cross-correlating each frame to ~10 frames of similar dose, resulting in an over-determined set of image shifts.  In Zorro the error amongst image shifts is optimized by a global-minimizer, the Basin hopping algorithm.  The individual correlations are logistic-weighted in the minimizer, based on the statistical significance of the correlation peak compared to the background noise. To deconvolve the impact of correlated noise shared between frames, a masked intensity-normalized cross-correlation (MNXC) is used. The MNXC algorithm also deconvolves non-uniform ice thickness and illumination. 

    A test specimen of urease particles are shown in Fig. 1 with (top) single frames with dose 0.3 e-2 and (bottom) the aligned sum of 60 frames. The set of correlations can be represented by an upper triangular matrix (Fig. 2), where the diagonal represents the frame number and the horizontal the adjacent cross-correlation maxima. The correlations rapidly drop to the noise level after the adjacency exceeds  frames due to non-rigid particle motion and radiation damage. Registration success at low-resolution may be assessed by independently registering the even- and odd-numbered frames, and calculating the cross-correlation between the independent halves. The normalized, rotationally averaged correlation, known as the Fourier Ring Correlation (FRC), is shown in Fig.3 for a small (145 kDa) protein. FRC oscillations are due to defocus.

    Zorro is a general-purpose algorithm without technique-specific heuristics.  Here we show it applied to dose-fractionated HAADF-STEM.  Shown in Fig. 4, a (top) rapid-scan rate with 100 ns dwell can be combined with dose-fractionation to (bottom) dampen scan errors in the sum.

 

Acknowledgements: Julia Kowal recorded the image used for the FRC in Fig. 3.  Kenneth Goldie and Ariane Fecteau-Lefebvre are thanked for maintaining the TEM instruments.

Robert MCLEOD (Basel, SWITZERLAND), Benedikt HAAS, Henning STAHLBERG
15:45 - 16:00 #6399 - IM07-OP135 Quantitative phase imaging with using orientation-independent differential interference contrast (OI-DIC) microscopy.
Quantitative phase imaging with using orientation-independent differential interference contrast (OI-DIC) microscopy.

Conventional differential interference contrast (DIC) microscope shows the two-dimensional distribution of optical phase gradient encountered along the shear direction between two interfering beams. Therefore, contrast of DIC images varies proportionally to cosine of the angle made by azimuth of the phase gradient and the direction of wavefront shear. The image contrast also depends on the initial phase difference (bias) between the interfering beams. To overcome the limitations of DIC systems, we have developed a quantitative orientation-independent differential interference contrast (OI-DIC) microscope, which allows the bias to be modulated and shear directions to be switched rapidly without mechanically rotating the specimen or the prisms [1]. A set of raw DIC images with orthogonal shear directions and different biases is captured within a second. Specialized software computes the phase gradient vector map and then the quantitative phase image.

The new OI-DIC beam-shearing assembly is shown in Fig.1. It consists of two standard DIC prisms with a liquid crystal 90º polarization rotator in between. The shear plane of the first prism DIC1 is oriented at 0º, and the shear plane of the second prism DIC2 is oriented at 90º. Another liquid crystal cell works as a phase shifter, which modulates the bias. Its principal plane is oriented at 0º. We employed a twisted-nematic liquid crystal cell as 90º rotator and an untwisted nematic cell as phase shifter. The OI-DIC technique can use any high-NA objective lens at the full aperture and provides an optical path length (OPL) or phase map with the highest resolution. Unlike other phase mapping techniques, the OI-DIC does not require phase unwrapping and calibration. The OI-DIC can also be combined with other imaging modalities such as fluorescence and polarization.

An example of the computed OPL gradient map is shown in Fig. 2. The image displays a 4-µm thick glass rod that is embedded in immersion liquid with refractive index 1.47. The image brightness is linearly proportional to OPL gradient magnitude. White level corresponds to gradient magnitude 200 nm/nm. The hue depicts the gradient direction, as it is illustrated by the color wheel in the left bottom corner. We used microscope Olympus BX61 equipped with objective lens UPlanSApo 100x/1.40 Oil.

The obtained OPL gradient map was processed by Fourier integration to compute the OPL (phase) map, which is represented in Fig. 3. The image brightness is linearly proportional to OPL and phase. White corresponds to 500nm (OPL) and 5.75rad (phase) at wavelength λ=546 nm.

Fig.4 displays cross-sections of the OPL and phase maps of 4-µm thick glass rods in immersion liquids with the refractive indices 1.47 (red curve), 1.51 (orange curve), 1.54 (blue curve), 1.56 (violet curve), and 1.58 (green curve). Refractive index of the glass is 1.56. An extremum OPL is determined by formula:

OPL=(nr -nim)d,

where nr  and nim are refractive indices of rod and immersion, respectively, d is diameter of the rod. As one can see, the OPL maxima and minimum are practically equal to the theoretical values 360nm, 200nm, 80nm, 0nm, and -80nm.

The OI-DIC assemblies fit into existing slots of a regular research grade microscope. We confirmed that a microscope upgraded with the OI-DIC provides lateral resolution ~200 nm and axial resolution ~100 nm at wavelength 546 nm. The OPL noise level was ~0.5nm. According our best knowledge, the images with such high level of resolution cannot be produced by any other currently available interference and phase microscopy techniques.

Acknowledgements

This publication was made possible by Grant Number R01-GM101701 from the National Institute of General Medical Sciences, National Institutes of Health (USA). Its contents are solely the responsibility of the author and do not necessarily represent the official views of the National Institute of General Medical Sciences or the National Institutes of Health.

References:

[1] M. Shribak, “Quantitative orientation-independent DIC microscope with fast switching shear direction and bias modulation, ” The Journal of the Optical Society of America A, 30 (2013), p. 769-782.

Shribak MICHAEL (Woods Hole, USA)

13:45-15:45
Added to your list of favorites
Deleted from your list of favorites

IM6-II
IM6: Quantitative Diffraction
SLOT II

IM6: Quantitative Diffraction
SLOT II

Chairmen: Tatiana GORELIK (Mainz, GERMANY), Damien JACOB (Lille, FRANCE)
13:45 - 14:15 #8769 - IM06-S48 Three-dimensional nanostructure determination from large diffraction dataset.
Three-dimensional nanostructure determination from large diffraction dataset.

  Electron tomography relies on mass-thickness contrast. However, chemically homogeneous nanostructures can not be studied using this technique. Nanocrystalline materials in general are structurally featured by a large density of grain boundaries and large surface/volume ratio, which have attracted significant interest for their unique mechanical, chemical and electronic properties.1-3 Transmission electron diffraction (TED) is an appropriate technique for complex nanostructure analysis because it is highly sensitive to local structure and it can be obtained using a small electron beam. 

  Previously we have developed a TEM based scanning electron nanodiffraction (SEND) technique that uses the built-in TEM deflection coils to shift the beam.Here, we report a new technique called 3D-SEND by coupling the SEND with diffraction tomography. This technique aims at determining the 3D morphologies and orientations of grains in nanocrystalline materials. A special holder design is employed for a high-angle (up to 87°) sample rotation.5 The design employs a needle-shape specimen and a sample rotation driven by the goniometer itself. Diffraction pattern recording and beam scanning are automated using a DigitalMicrograph® script to control the TEM deflection coils and camera readout.6 A stack of diffraction patterns are acquired for each sample rotation step.

  The reconstruction starts with the identification of 2D grain morphologies, using dark-field images generated with diffraction spots. Similar contrasts are grouped using the normalized cross-correlation. Diffraction spots belong to one grain are sorted into one diffraction pattern. Electron diffraction pattern indexing is achieved by a combination of diffraction peak search and peak indexing using both length and angle information. The tomographic reconstruction of the grain is performed using the algebraic reconstruction technique. We apply prior conditions concerning the sample outline and the damping of scattering intensity. A smoothed isosurface is created to illustrate the grain’s 3D morphology. The grain orientation is determined from the indexing results.

  We demonstrate the performance of 3D-SEND on a TiN thin-film nanocrystalline sample prepared by unbalanced magnetron sputtering. The FIB cut and lift-out was performed perpendicular to the growth direction. The sample was milled to a tip with a diameter of 200 nm. The beam was set to 7nm in FWHM. The scanning covered a 2626 pixels area and each step is 11nm. The sample was tilted over a range of ±85° with a step of 5°. In total 23660 DPs are recorded. Figure 1 shows the reconstruction results of seven major grains and their orientations. The reconstruction result is shown in Fig. 1. 

  The spatial resolution of this technique is ultimately limited by the electron probe size under the column approximation. For a JEOL 2100 TEM, we can form a probe with a FWHM of 2.3 nm using a 10 um condenser aperture in the CBD mode with alpha = 1. Compared with the alternative technique 3D-OMiTEM,7 3D-SEND has a better diffraction pattern resolution, a wider sample rotation range and a much lower electron dose. Our approach also grants a combinative study with other techniques such as atom-probe tomography and in-situ deformation. In the future, 3D-SEND may potentially be improved for the five-parameter characterization of grain boundaries.

References:

1.   E. Abe et al, Acta Materialia, 50(2002), p. 3845-3857.

2.   D. Wang et al, ACS Nano, 3(2009), p. 907-914.

3.   C. C. Koch et al, MRS Bulletin, 24(1999), p. 54-58.

4.   K. H. Kim et al, Micron, 71(2015), p. 39-45.

5.   S. Mao et al, Acta Materialia, 82(2015), p. 328-335.

6.   J. Tao et al, Physical Review Letters, 103(2009), p. 097202.

7.   H. H. Liu et al, Science, 332(2011), p. 833-834.

*We thank Prof. Huang of National Tsinghua University for providing the TiN sample and DOE BES DEFG02-01ER45923 support. 

Yifei MENG, Jian Min ZUO (Urbana, IL, USA)
Invited
14:15 - 14:30 #6606 - IM06-OP122 Soft matter, and the many flavours of diffraction tomography.
Soft matter, and the many flavours of diffraction tomography.

With the arrival in 2009 of diffraction tomography for crystal structure solution [1] of sub-micron sized crystals, there has been a renaissance in the application of electron diffraction, with diffraction tomography providing a relatively easy path to quantitative analysis of crystal structures, compared to more conventional electron diffraction approaches. The collection of diffraction tomography data occurs by tilting the crystal around an arbitrary axis, which runs counter to the traditional notions of electron diffraction from low index orientated zones. Given the success of the method, it has produced a series of derivative techniques on this theme, of collecting tomographic diffraction from around an arbitrary axis. Each of these methods comes with its strengths and weaknesses.

An investigation into the relative merits of the various diffraction tomography techniques, ADT[1], RED[2], EDT[3], Rotation method[4], MicroED[5]. With an emphasis on the application of the techniques for soft matter and radiation sensitive materials, such as organic and protein crystals. Evaluation of the nuances in data collection strategies, hardware requirements, specimen requirements, and optional components, will all be considered, with details of how they affect the experimental setup, data acquisition, and processing.

We will discuss the need for validation of structural models and the added difficulties presented in doing so with soft matter materials, often leading to the requirement of validation by additional techniques.

[1] E. Mugnaioli, T. Gorelik and U. Kolb, Ultramicroscopy 109 (2009), 758-765.
[2] D. Zhang, P. Oleynikov,, S. Hovmoller & X. Zou, Z. Kristallogr. Cryst. Mater., 225 (2010), 94-102.
[3] M. Gemmi, P. Oleynikov, Z. Kristallogr. 228 (2013), 51–58.
[4] I. Nederlof, E. van Genderen, Y.-W. Li & J. P. Abrahams,  Acta Cryst.  D69 (2013), 1223–1230.
[5] D. Shi, B. Nannenga, M.G. Iadanza and T. Gonen (2013), eLife – 2:e01345: 1 - 17.

Andy STEWART (Limerick, IRELAND)
14:30 - 14:45 #6799 - IM06-OP125 First woven covalent organic framework solved using electron crystallography.
IM06-OP125 First woven covalent organic framework solved using electron crystallography.

Making fabric by weaving is known as one of the oldest and most enduring methods. Nevertheless such an important design concept still needs to be emulated in extended chemical structures. Linking molecules into weaving structures would be of a great help to create materials with exceptional mechanical properties and dynamics. For this purpose a woven covalent organic framework-505 (COF-505) has been synthesized using a designed strategy [1]. However, COF-505 is not well crystallized, which gives rise to a poorly resolved PXRD pattern. Therefore, approaches based on electron crystallography methods have been used. The structure of this COF has been solved by a combination of 3D electron diffraction tomography (3D EDT, [2]), high-resolution TEM imaging and structure modeling.

3D EDT dataset was collected from a single sub-micron crystal in a tilting range of –41.3° to +69.1°. The reconstructed 3D reciprocal lattice was identified as a C-centered orthorhombic Bravais lattice with the unit cell parameters of a = 18.9 Å, b = 21.3 Å, c = 30.8 Å, and V = 12399 Å3, which have been used to index reflections observed in both PXRD pattern and Fourier diffractograms of HRTEM images. The derived reflection conditions were summarized as hkl: h+k = 2n; hk0: h, k = 2n; h0l: h = 2n and 0kl: k = 2n, leading to five possible space groups (s.g.): Cm2a (39), Cc2a (41), Cmca (64), Cmma (67) and Ccca (68). Cm2a, Cmma and Ccca were excluded because their projected plane group symmetries along [1-10] do not coincide with those of the experimental HRTEM images (pgg). Furthermore, by performing Fourier analysis of the HRTEM images and imposing symmetry to the reflections, Cu(I) positions were determined from the reconstructed 3D potential map (Fig. 1). The structure of COF-505 was built in Materials Studio by putting Cu(PDB)2 units at copper positions and connecting them through biphenyl (reacted BZ) molecules. The chemical composition was determined by the elemental analysis, which indicated that the unit-cell framework is constructed by 8 Cu(PDB)2 and 16 biphenyl units (Fig. 2). However, symmetry operations of the s.g. Cmca require two PDB units connected to one copper onto a mirror plane perpendicular to the a–axis that is not the energetically favorable geometry. The final s.g. determined as Cc2a and was used to build and optimize a structure model. The PXRD pattern calculated from the model is consistent with the experimental pattern of activated COF-505.

 

Acknowledgements:

Grants from Swedish Research Council/VR (Y.M. and P.O.) and JEOL Ltd, Japan (P.O.); EXSELENT and 3DEM-Natur, Sweden (O.T.) and BK21Plus, Korea (O.T.).

 

References:

[1] Y. Liu, Y. Ma, Y. Zhao, X. Sun, F. Gandara, H. Furukawa, Z. Liu, H. Zhu, C. Zhu, K. Suenaga, P. Oleynikov, A. S. Alshammari, X. Zhang, O. Terasaki, O. M. Yaghi. Weaving of organic threads into a crystalline covalent organic framework. Science, 351 (2016) 365–369.

[2] M. Gemmi, P. Oleynikov. Scanning reciprocal space for solving unknown structures: energy filtered diffraction tomography and rotation diffraction tomography methods. Z. Krist. 228 (2013) 51–58.

Yuzhong LIU, Yanhang MA, Yingbo ZHAO, Peter OLEYNIKOV (Stockholm, SWEDEN), Osamu TERASAKI, Omar YAGHI
14:45 - 15:00 #6316 - IM06-OP119 Structure solution of the complex γ-La6W2O15.
Structure solution of the complex γ-La6W2O15.

Oxides in the Ln2O3-MO3 (M = Mo and W) system are of significant technological interest for their laser applications [1], ionic conduction [2], catalytic [3] and ferroelectric [4] properties. The La2O3-WO3 phase diagram has been studied by a number of groups [5-7], but little detailed crystallographic information was reported due to the lack of good single crystals. Some of the reported compositions have not been appropriately characterized. Recently, the structures of La2WO6, La18W10O57 and La10W2O21 were solved using X-ray powder diffraction (XRPD) [8-10].

 

For the La6W2O15 compound phase transitions at 630 and 930 °C have been reported [1-3]. The structure of the high temperature phase α-La6W2O15 was determined ab-initio by XRPD [11]. The lower-temperature forms β and γ, however, couldn’t be determined due to the large number of reflections in the X-ray powder diffraction pattern and the relatively low symmetry of the system. The existing literature on γ-La6W2O15 only relates two sets of unit cell parameters [5-6], that almost match the XRPD pattern of γ-La6W2O15, but some weak peaks remain without indexation and can’t be explained by the presence of any impurity.

 

Here we present the structure solution using transmission electron microscopy of the complex structure of γ-La6W2O15. From zone axis precession electron diffraction the unit cell was determined to be monoclinic with cell parameters a=1.57 nm, b=1.21 nm, c=1.57 nm, β=110°. As an example, the [100] zone axis is presented on figure 1. Due to the low symmetry of the crystal system and the large unit cell, a huge number of reflections needed to be acquired, so that electron diffraction tomography was used to record the intensities. The cation positions were obtained but the distribution of the cations on the sites was not evident. Z-contrast imaging showed that disorder on some cationic sites has to be considered (fig.2).

 

[1] Kumaran et al, J Cryst Growth 292 (2006) 368-372

[2] Lacorre et al, Nature 404 (2000), 856-858

[3] Alonso et al, J Solid State Chem 177 (2004) 2470-2476

[4] Brixner et al, J Solid State Chem. 5 (1972) 186-190

[5] Yoshimura et al, Mater Res Bull 11 (1976) 151-158

[6] Yanoskii et al, Sov Phys Crystallogr 20(3), 354-355

[7] Ivanova et al, Inorg Mater (1970) 803-805

[8] Chambrier et al, J Solid State Chem 183 (2009) 209-214

[9] Chambrier et al, Inorganic Chemistry 48 (2009) 6566-6572

[10] Chambrier et al, Inorganic Chemistry 53 (2014) 147-159

[11] Chambrier et al, J Solid State Chem 183 (2010) 1297-1302

Stéphanie KODJIKIAN (Grenoble), Christophe LEPOITTEVIN, Holger KLEIN, Thomas SCHÖNENBERGER, Oleg LEBEDEV, Olivier LEYNAUD, Marie-Hélène CHAMBRIER, François GOUTENOIRE
15:00 - 15:15 #6780 - IM06-OP124 Short-range-order (SRO) in quenched Ni-rich Ni-Ti alloys.
IM06-OP124 Short-range-order (SRO) in quenched Ni-rich Ni-Ti alloys.

Binary Ni-Ti alloys have a wide application in industry and medicine due to their shape memory effect and superelasticity properties. These mechanical properties are known to be caused by a martensitic transformation of which the characteristics are strongly dependent on Ni4Ti3 precipitates formed during aging.

In this study a Ni-Ti alloy which is quenched immediately after the production and aged at room temperature is investigated. No precipitation is expected to form in the sample, which is confirmed by conventional TEM images. However, the alloy still shows a changing shape memory effect with ageing time at room temperature, which indicates there must exist some small structural changes not visible by conventional TEM. These are also expected from observed diffuse intensities (Fig. 1) arranged around particular geometrical loci in reciprocal space. Different techniques are proposed and used to identify these microdomains. The Cluster Model [1] that assigns the shape of the diffuse reciprocal intensity to that of microdomains is applied to analyze the results. In the present case the diffuse intensity can to a first order be approximated by {111}* reciprocal planes, which can be translated into atomic rows along the [111] crystallographic directions in the cubic Ni-Ti lattice. Such rows of pure Ni are also present in the crystal structure of Ni4Ti3 precipitates, as seen in Fig. 2. In other words, the diffuse intensity can be correlated with contiguous strings of Ni atoms in the cubic directions of the B2 matrix, which normally reveals a …-Ni-Ti-Ni-Ti-… sequence along these directions. In order to observe such atomic strings in real space, aberration corrected HAADF-STEM has been performed along a cubic direction. Simulations indicate that the clustering of heavier Ni atoms can be seen as an increment of appr. 2% of intensity of a single atomic column due to the Z-contrast nature of the HAADF-STEM imaging concept. The experimental image shown in Fig. 3a indeed shows some random but coagulated fluctuations in intensity of the columns, as can be seen from the line trace in Fig. 3b (as well as by slightly defocusing your eyes when looking at the picture). However, to what extend these can be attributed to the atomic clustering is still not clear. In the near future, also other advanced TEM techniques will be applied in order to further identify the columns containing contiguous strings of Ni atoms.

[1] D. van Dyck, R. de Ridder, G. van Tendeloo, and S. Amelinckx, “A cluster model for the transition state and its study by means of electron diffraction. III. Generalisations of the theory and relation to the SRO parameters,” Phys. Status Solidi A, vol. 43, no. 2, pp. 541–552, Oct. 1977.

Saeid POURBABAK (Antwerp, BELGIUM), Xiebing WANG, Bert VERLINDEN, Jan VAN HUMBEECK, Dominique SCHRYVERS
15:15 - 15:30 #5718 - IM06-OP115 Radial distribution function imaging by STEM diffraction: a method development in resolving the mysteries of amorphous materials.
Radial distribution function imaging by STEM diffraction: a method development in resolving the mysteries of amorphous materials.

Interpreting atomic structure of amorphous materials have attracted attentions for a century and, in recent years, especially heterogeneous nanoglasses have fueled the interest because of their unusual structure and properties [1]. However, only few experimental means offer a way to characterize the disordered structures. Atomic radial distribution function (RDF) is one of the important tools for the goal, which was first obtained from X-ray diffraction data of organic solids [2], and afterward extended to electron diffraction for metallic glasses [3]. RDF describes the probability to find certain atomic pairs as a function of the pair separation and consequently, provides structural information in the short- and medium-range. However, the traditional diffraction experiments only average large sample areas and lack spatial resolution especially at nanometer scale. Plenty information is hidden in the averaged signal.

 

In this work, we demonstrate a newly developed method, RDF-imaging, combining diffraction imaging in scanning transmission electron microscopy (STEM) mode [4] with RDF analysis and spectral-imaging analysis (e.g. multiple linear least square (MLLS) fitting) to achieve structural mapping of heterogeneous amorphous materials with 1 nm resolution. Figure 1 schematically shows the procedure for data acquisition and RDF calculation: a 4-dimensional (D) diffraction-image is acquired by recording diffraction patterns in STEM mode with quasi parallel nano-beam configuration (0.8 mrad convergence angle) and 1 nm spot size. RDFs are calculated from the diffraction patterns according to description in [3][5]. A 3D data cube of RDFs (RDF-cube) is constructed by relating pixel positions of the initial diffraction-image to the calculated RDFs. The RDF-cube can then be analyzed by MLLS fitting of reference spectra taking experimentally from pure phase or numerically from matrix decomposition by multivariate statistical analysis.

 

The method is extremely sensitive to atomic packing variation. Figure 2 shows an application to amorphous ZrO2 (a-ZrO2) and a-ZrFe multilayers. Not surprisingly, both the a-ZrO2 and the a-ZrFe phases are unambiguously distinguished by RDF-image (figure 2d,e), but also an interface layer between ZrO2 and ZrFe (figure 2c) is detected, which could not be identified in STEM-EELS and EDX maps. The atomic structure (figure 2a, red dashed-line) of the interfacial layers possesses the same atomic packing as that of the a-ZrO2 phase (figure 2a, blue line) but with a 0.04 Ǻ shrinkage of the average bonding distance.  The shift in bonding distance could be introduced by Fe replacing Zr atoms in the ZrO2 clusters due to diffusion of Fe atoms from the a-ZrFe into the ZrO2 layers.

 

Acknowledgement: Authors thank financial support from KNMF and Hi-C project.

References

[1] R Witte, T Feng, JX Fang, A Fischer, M Ghafari, R Kruk, R Brand, D Wang, H Hahn, H Gleiter, Appl Phys Lett 103 (2013), p. 073106.

[2] T Egami and S J L Billinge in “Underneath the Bragg peaks structural analysis of complex materials”, (Elsevier Ltd, Kidlington, Oxford, UK), p. 55.

[3] D J H Cockayne and D R Mckenzie, Acta Crystallographica Section A 44 (1988), p. 870.

[4] C Gammer, V BurakOzdol, C H Liebscher and A M Minor, Ultramicroscopy 155 (2015), p. 1.

[5] X Mu, S Neelamraju, W Sigle, C T Koch, N Toto, J C Schon, A Bach, D Fischer, M Jansen and P A van Aken, J. Appl. Cryst. 46  (2013), p. 1105.

Xiaoke MU (Eggenstein-Leopoldshafen, GERMANY), Di WANG, Tao FENG, Christian KÜBEL
15:30 - 15:45 #6738 - IM06-OP123 Study of amorphous silica by Electron Energy Loss Spectroscopy and electron diffraction PDF.
IM06-OP123 Study of amorphous silica by Electron Energy Loss Spectroscopy and electron diffraction PDF.

Amorphous silica (a-SiO2) has many industrial (glass former) and scientific applications. Structurally silica which has SiO4 tetrahedral units connected with bridging O at the corners. To understand and predict interesting properties of such materials depends on the knowledge of the detailed atomic structure. The loss of structural ordering in amorphous materials, carry sharp Bragg diffraction reflections to disappear and only diffuse diffraction pattern are observed. Study of amorphous materials cannot be performed by routine crystallography and techniques like Pair Distribution Function (PDF) can be used for structure analysis using X-Ray, neutron or electron diffraction (ED) [1]. ED  related PDF (e-PDF)  in TEM has the big advantage over X- Ray PDF technique that allows studying local structural ordering of amorphous materials in nm scale by collecting (ED) patterns in very short time (msec instead of 15-24 hours in X-Ray case).  Here we present how e-PDF analysis can be used to study in detail hydrothermal reactions flow process in amorphous silica.

Amorphous silica samples were heated with water in hydrothermal container under different time scales and were studied by X-ray, Electron Energy Loss Spectroscopy and e-PDF. Hydrothermal reaction on amorphous silica was performed with reaction times 6, 16, 312 hours [2].  Figure 1 show the Si L2,3 edges for studied compounds. Si peaks are always observed but its giving rise to a broadening of spectra as the intensity increases. Such broadening in the Si L peaks is probably due to Si tetrahedral distortion, since in SiO2 amorphous phase, distances between silicon and the four surrounding oxygen are slightly different. From the collected ED data, pair distribution function (G(r)) was calculated using dedicated ePDF software [3], developed to analyse ED patterns from amorphous and nanomaterials (Figure 2). Our analysis with ePDF has revealed that no peak (corresponding to interatomic distances) was found beyond 5 Å, which confirms that only short range ordering present in the material, even after several hours of hydrothermal reaction. Besides, small changes were observed in the PDF peak positions (corresponding to interatomic Si-Si, Si-O and O-O distances). All such peaks/interatomic distances match well with the distances existing within the SiO2 crystalline structure. Slight peak width change has also been observed with reaction time, which can possibly arise from local strain. Modelling of the amorphous silica reactive product is under process.

 

1. G. R. Anstis, Z. Liu, M. Lake, Ultramicroscopy, 26, (1988), 65

2. L. Khouchaf, A. Hamoudi and P. Cordier, Journal of H. Materials, 168, (2009) 1188

3. A. M. M. Abeykoon, H. Hu, L. Wu, Y. Zhu, S. J. L. Billinge, J. Appl. Cryst., 48, (2015) 244

Lahcen KHOUCHAF , Khalid BOULAHYA (Madrid, SPAIN), Das PARTHA, Janos LABAR, Viktoria VIS, Stavros NICOLOPOULOS