bluuebunny/crossref_metadata_2025_split huggingface dataset:
Deformation mapping in a
TEM
: Dark Field Electron Holography, Nanobeam Electron Diffraction, Precession Electron Diffraction and GPA compared
Authors: Cooper David , Bernier Nicolas , Rouviere Jean‐Luc 
Published: December 2016

The properties of nanoscaled materials can be changed by applying strain and as such there is an interest in the accurate measurement of deformation with nm‐scale resolution. This was until recently considered as a difficult problem. However, the last ten years has seen a great deal of development in techniques that can be used to measure deformation with the required resolution [1,2]. Today there are many different approaches which can be used to recover valuable information about the deformation. Each of these techniques has strengths and weaknesses and requires different set ups in the electron microscope [2]. In this presentation we will present dark field electron holography, nanobeam electron diffraction (NBED), precession diffraction (NPED) and the geometrical phase analysis (GPA) of TEM and high angle annular dark field (HAADF) scanning transmission electron microscopy (STEM) images. We will then discuss which technique is most suitable for different types of materials problems and benchmark their performance with respect to the accuracy of the measurements, precision and spatial resolution.
Figure 1(a) shows a HAADF STEM image of a 100‐nm‐thick Si test specimen containing 10‐nm‐thick SiGe layers with different Ge concentrations. As the specimen has been grown epitaxially we expect no deformation to be measured in the ex direction. However, due to the expanded lattice parameter of the SiGe layers relative to the Si reference, tensile deformation is expected in the ez direction. Figure 1 shows deformation maps that have been acquired by (b) GPA of HAADF images (c) dark holography and (d) precession diffraction. These are compared to finite element simulations that are shown in Figure 1(e). These results reveal that all of the different techniques provide accurate measurements of the deformation. Figure 2(a) shows a HAADF STEM image of a SiGe test device structure with a gate length of 35 nm. Finite element simulations showing the expected deformation in the struture is shown in Figures 2(b) and (c). Deformation maps are shown in Figure 2(d) and (e) for precession diffraction, (f) and (g) for dark holography and (h) and (i) for GPA of HAADF STEM images [4]. Again, accurate measurements of the deformation are made, but the precision and spatial resolution depends on the experimental technique that has been used. As well as presenting state of the art results from a range of strained materials we will highlight improvements that are required for all of the different techniques in order to optimise their performance and provide the best possible measurements of deformation.
Acknowledgements : These experiments have been performed on the platform nanocharacterisation (PFNC) at Minatec. The work has been funded by the ERC Starting Grant 306365 « Holoview ».
Vectorial field tomography from single projection for cylindrically symmetric nanostructures
Authors: Phatak Charudatta , Knoop Ludvig De , Houdellier Florent , Hytch Martin , Masseboeuf Aurélien 
Published: December 2016

One‐dimensional (1D) nanostructures have been regarded as the most promising building blocks for nanoelectronics and nanocomposite material systems as well as for alternative energy applications [1]. Magnetic nanowires with circular cross‐section, are of utmost importance from theoretical and technological aspects [2]. 1D carbon‐based nanostructures such as carbon nanotubes are amongst the best candidates for field emission displays and new high‐brightness electron sources [3]. The confinement effects in 1D nanostructures can alter their properties and subsequently their behavior significantly. Hence it is necessary to understand the strong effect of their size on their three‐dimensional (3D) properties such as the magnetic and electric fields associated with nanowires and nanotubes completely before they can be used in applications. There are currently very few methods, which have the capability to visualize the complete 3D fields associated with nanowires. In this work, we show that using a combination of symmetry arguments and electron‐optical phase shift data obtained using TEM, it is possible to recover the entire 3D magnetic or electric field in and around nanowires and nanotubes from a single image.
The phase change φe+φm of an electron wave traveling through a thin foil can be expressed in terms of tomographic quantities [4]. The phase shift can be recovered experimentally using various techniques such as transport‐of‐intensity based methods or off‐axis electron holography. Nominally, determining the 3D magnetic field or 3D elestrostatic potential requires recording a series of phase shift images as the sample is tilted about its axis. However, in a 1D nanostructure such as a nanocylinder that is uniformly magnetized along its long axis, the magnetic field possesses cylindrical symmetry with respect to the long axis. Similarly, carbon nanotubes under applied bias exhibit a cylindrically symmetric potential and electric field. Exploiting these conditions of cylindrical symmetry, we can reconstruct the full 3D magnetic or electric field associated with a nanowire from a single phase image using the inverse Abel transform. Simulations were performed for a uniformly magnetized sphere, which also possesses cylindrical symmetry to analyze the fidelty of the reconstruction algorithm. Fig. 1(a) shows the magnetic phase shift of such a uniformly magnetized sphere, (b) shows the derivative with respective to the horizontal axis (∂φm/∂x) showing the cylindrical symmetry, and (c) the 3D reconstructed magnetic induction using the single image method developed in this work. Experiments were performed on nickel nanowires that were uniformly magnetized to reconstruct the 3D magnetic induction. Similarly the 3D electric field was reconstructed from in‐situ biased carbon cone nanotips under varying applied bias.
Mapping
2D
strain components from
STEM
moiré fringes
Authors: Hytch Martin , Gatel Christophe , Ishizuka Akimitsu , Ishizuka Kazuo 
Published: December 2016

Artificial moirés are created in a STEM by deliberately choosing a low magnification where the scan step is close to the crystalline periodicity (see Figure 1a) [1]. A moiré contrast then results from the interference between the scan and the crystal lattice. The technique has been developed to analyse strain [2], and has been applied to the study of strained‐silicon devices [3].

In reciprocal space, STEM moiré fringes can be understood as the convolution of the lattice created by the scan, characterized by the scan‐step, s in real‐space, or s
*
in reciprocal space, and the reciprocal lattice of the crystalline lattice, characterized by d
*
(Figure 1b). Interference between neighbouring periodicities gives rise to moiré fringes of periodicity
q
M
. In general, only a few moiré fringe periodicities will be present in the image as the scan reciprocal lattice is in reality multiplied by the MTF of the probe: the size of the probe limits the possible interference terms (indicated by the yellow area in Figure 1b). The moiré peridocity
q
M
is related vectorially to
s
*
, which is usually 1 or 2 pixel
−1
, and the
d
*
(or g‐vector) for a particular set of lattice fringes (Figure 1c).

Here we show how the strain information can be extracted using the concept of geometric phase, previously used for the analysis of high‐resolution TEM images [4]. The advantage of this formulation is that the moiré fringes do not need to be aligned exactly with the crystalline lattice, thus freeing up the experimental work. The periodicity of the moiré fringes (Figure 2a) is identified from the power spectrum of the image and the corresponding geometric phase determined (Figure 2b). The strain is in turn calculated from the geometric phase (Figure 2c). It is not necessary to know the exact calibration of the original image providing a reference region of known crystal parameter is present, as the vectorial relationship (Figure 1c) provides a strong constraint.
We have also developed a procedure to determine maps of the 2D strain tensor from two differently oriented SMFs [5] (Figure 3). The results from two separate moiré fringe images need to be aligned (Figure 3b) and combined to determine the full 2D in‐plane strain tensor. Whilst the examples here are simulated, we anticipate presenting experimental results from devices and piezo‐electric materials.
New approach for low dose electron diffraction tomography
Authors: Ghadimi Reza , Tietz Hans , Olyenikov Peter 
Published: December 2016

Due to strong interaction of electrons with matter, 3D electron diffraction tomography has been proven to be a reliable method to solve structures of very small crystals, compared to X‐ray diffraction; even for small protein crystals [1].
Several procedures for automated acquisition of tomographic diffraction data sets have been described [2, 3], using precession technique or discrete beam tilt perpendicular to the goniometer tilt axis. However, those mentioned methods are mainly used for less beam sensitive materials and are not suitable for low dose applications.

Shi et al. [4–7] refined the collection of 3D data sets under low dose condition (less than 10 e
−
/Å
2
). Here, the camera system is continuously acquiring diffraction data during a continuous tilt of the goniometer. This allows to completely scan the Fourier space. It turned out that the stability of the goniometer tilt speed and the limited flexibility of the tilt acquisition parameters are problematic.

Here, we present a newly developed automatic data acquisition system, combining real‐time direct control of the TEM‐deflection systems, the goniometer tilt and the acquisition of high‐resolution diffraction patterns with a synchronized CMOS camera. This can be realized by the TVIPS
U
niversal
S
can
G
enerator (USG), controlling eight TEM deflection coils, i. e. four beam deflection coils and four image deflection coils.

For a static goniometer alpha, the total beam tilt range (±3° to ±5°, depending on the TEM) can be fragmented by a defined beam tilt range, e.g. 0.5° (see Fig. 1). During the camera exposure time the beam is continuously tilted for this defined range (e.g. 0.5°). The continuous beam sweep during exposure time ensures the complete sampling of the Fourier space.

Several crystal structures (Carbamazepin: C
15
H
12
N
2
O, bismuth oxychloride: BiOCl, Mayenite: Ca
12
Al
14
O
33
) have been successfully solved by this method. The total acquisition time for 190 high resolution (1Å resolution) diffraction patterns is about 15 minutes, with a total electron dose of about 10 e
−
/Å
2
. The collected 3D electron diffraction data sets were processed using the EDT‐PROCESS software [8].
Electrical transport‐induced transformations in filled carbon nanotubes imaged in situ by scanning‐transmission electron microscopy
Authors: Lozano Juan G , Aslam Zabeada , Nicholls Rebecca J , Koos Antal A , Dillon Frank , Sarahan Michael , Nellist Peter D , Grobert Nicole 
Published: December 2016

Fe‐filled CNTs have been proposed as perfect candidates for a large and varied number of applications, ranging from the biological to optoelectronics or memory storage devices1. For any of these applications, it seems crucial to understand the interactions, possible phase changes or reactions that can take place as a consequence of exposure to their real working conditions.Conventional bright‐field transmission electron microscopy (BF‐TEM) has been the tool traditionally employed to observein‐situthe dynamical effects that take place when individual CNTs are exposed to high electrical currents and Joule heating. Despite being a much more powerful technique, scanning TEM‐annular dark field (STEM‐ADF) imaging has been scarcely for this purpose2. One of the main advantages of STEM‐ADF over BF‐TEM is that the intensity in the images is highly dependent on the atomic number of the species which are present and therefore the images show compositional in addition to structural information. This gives a much better understanding of current‐induced migration effects, formation of intermediate phases or alloying, or phase separation phenomena, something that using BF‐TEM alone would miss. In the STEM‐ADF configuration it is also possible to perform complementary analytical spectroscopy simultaneously with imaging and with comparable spatial resolution.Here we show the significant advantages of combiningin‐situexperiments with STEM‐ADF and related analytical techniques to gain new insights into the electrical transport induced transformations in P, N doped Fe‐filled carbon nanotubes. It has been possible to monitor in real time a multistage process (Figure 1) in which the Fe filling reacts with nitrogen to form an intermediate alloy phase, which then decomposes into smaller particles. The presence of N2gas within the inner channel of the tube is found to be crucial to this process.
Multiphase material sample preparation using broad‐beam Ar ion milling for
EBSD
analyses
Authors: Nowakowski Pawel , Schlenker James , Ray Mary , Fischione Paul 
Published: December 2016

When compared to monophase materials, multiphase materials can offer an improvement in properties or introduce new properties. For example, tungsten, carbides, and nitrides do not have many applications as monophase materials; these metals are brittle, which can override their other desirable properties, such as hardness, high temperature stability, and wear. However, when these materials are dispersed in a soft matrix, such as low‐carbon steel, nickel, or cobalt, and form a composite, the multiphase materials are then characterized by desirable mechanical properties, such as ductility, strength, and resistance to creep and fatigue [1,2]. Other types of multiphase materials are protective surface coatings, such as tribological and antioxidation applications, and microelectronics devices [3].
Knowledge of the phases present in the material, their distribution, and their fractions is fundamental to materials characterization. A common investigative approach for engineered materials is electron backscatter diffraction (EBSD) combined with energy dispersive spectroscopy (EDS). EBSD provides information about crystallographic orientations or misorientations and allows the study of grain boundaries, deformation, and recrystallization. EDS can define the chemical composition of the materials. When combined, these techniques permit phase identification with high accuracy [4].
Sample preparation of multiphase materials for EBSD and EDS can be challenging, especially when the different phases have very different characteristics. For example, if one phase is hard and brittle and the second phase is soft and ductile, the rate of material removal during mechanical polishing can vary widely. In addition, the hard particles removed from the first phase can act as a grinding medium and tear the soft matrix of the second phase. The result is a surface inappropriate for EBSD analyses. Electrolytic etching is another sample preparation technique that is not ideal for multiphase materials; it often requires the use of toxic or dangerous chemicals, it can cause differential etching of the sample surface, and it can give the sample surface a varied topography, which will produce a shadowing effect in the EBSD/EDS analysis.
The goal of this work is to illustrate how the use of low energy, broad‐beam argon ion milling can improve and facilitate multiphase material sample preparation. Several examples are discussed, including hard particles within a soft matrix and a multiphase protective layer on a metal substrate.
Direct imaging of hydrogen atomic columns in hydride phases in titanium grade 2
Authors: Daněk Kamil , Gärtnerová Viera , Němec Martin 
Published: December 2016

Despite diligent work on titanium hydrides the crystal structures of some titanium hydride phases are still not well established. Transmission electron microscopes (TEMs) can support crystallographic data by direct imaging of unknown structures with superior resolution. Direct imaging of hydrogen atomic columns is extremely challenging due to their weak scattering. However, an introduction of TEMs with corrected spherical aberration has made such experiments feasible. In this work, direct imaging of hydrogen atomic columns in titanium hydrides was performed in order to provide additional support to incomplete crystallographic data for titanium hydrides.
The commercially pure titanium grade 2 annealed at 900 °C for 1 h was examined in this work. Samples for transmission electron microscopy (TEM) were prepared by mechanical grinding of 3 mm discs and subsequent thinning by twin‐jet electropolishing in perchloric acid solution using Struers TenuPol‐5. The titanium grade 2 contains 0.24 at.% of hydrogen, which can create hydride phases. Face‐centered tetragonal γ‐hydride phase was found in a lentil‐like shape in the α‐Ti matrix and in a shape of long thin lamellae at either the α‐Ti matrix/β‐Ti phase interface or originating at α‐Ti matrix grain boundaries. This phase introduces hydrogen‐induced diffracted reflections that do not coincide with the diffracted reflections coming from Ti atoms. For that reason, the presence of the hydrogen can be observed even using the selected area electron diffraction (SAED) technique. The hydrogen‐induced reflections translate also to a phase‐contrast pattern (PCP). Compared to the titanium columns, the hydrogen columns produce much broader peaks/dips that enable us to differentiate between the positions filled with hydrogen and the vacant columns in the structure.
Annular bright‐field scanning transmission electron microscopy (ABF‐STEM) and high resolution transmission electron microscopy (HRTEM) techniques were employed. As both techniques rely on phase‐contrast imaging, they can produce imaging artifacts that can be easily mistaken for genuine structural features caused by the presence of the hydrogen. To increase robustness of our results we support them with simulations based on theoretical models of the hydrides. Multislice method was used to simulate the dynamic scattering effects. This helped us to find optimal imaging conditions with regard to resolution, maximal contrast for the hydrogen positions and the absence of imaging artifacts for a wide range of thickness and defocus values.
Imaging by Zernike phase plates in the
TEM
Authors: Edgcombe Chris 
Published: December 2016

Abbe lens theory can be used to calculate the images produced by a system made up of a simple phase object, a round lens and a Zernike phase plate. When all these have rotational symmetry, 1D Fourier‐Bessel transforms are suitable.

Given a disc object producing a uniform phase shift over its diameter and a phase plate providing a simple step in phase by α at a defined spatial frequency q
0
, the resulting image intensity can be defined for any size of object phase change. The distortion of the image by the central hole in the Zernike plate depends on a parameter B which is proportional to the product of the diameter of the object and q
0
and also depends on the lens focal length; when B is less than 1, the distortion is small (Fig. 1). The range of object phase that produces a monotonic variation of intensity is found to depend strongly on the phase change α introduced by the phase plate, being greatest when α is between π and 2π (Fig. 2).

For this object and plate, the weak phase approximation (WPA) is useful in giving an indication of the radial distribution of intensity. It does not predict well the range of monotonic variation of intensity with object phase.

In the search to reduce the ‘ringing’ in the image produced by large objects with a stepped plate, we analysed also the behaviour when the phase change at the plate increases to its maximum value α over a range of radius (instead of a step). When the phase change is proportional to radius for q < q
0
, we find that as object diameter increases, the intensity is reduced gracefully when α ∼ π/2, but that contrast reversal can occur when α ∼ 3π/2.

The intensity has also been calculated in WPA for an object phase distribution typical of a spherical object. With this object, the image intensity varies continuously with object radius, as expected, and is reduced as B increases above 1. The behaviour when the phase profile of the plate is ramped is similar to that for the disc object with the same phase profile for the plate (Fig. 3).
The advantages for different kinds of objects offered by plates with phase shifts α either less than or greater than π suggest that in practical use it will be desirable for phase plate holders to be fitted with plates of both thicknesses.
Size and orientation effects on the mechanical response of a
Ni‐Ti‐Al‐Nb‐Hf
alloy
Authors: Ding Rengen , Zheng Lijing , Chiu Yulung 
Published: December 2016

NiTi‐Al based alloy is considered as a potential high temperature structural material [1]. The addition of Nb is able to improve the oxidation resistance of NiTi‐Al based alloys at ~ 800°C [2] and also their strengths at room and elevated temperatures [3]. A Ni‐43Ti‐4Al‐2Nb‐2Hf alloy has been demonstrated to be a promising lightweight high‐temperature structural material [4]. It is thus of significance to understand the effect of crystal orientations on its mechanical behaviours. In the past, mechanical testing of single crystal metals was performed on large crystal. However, the preparation of single crystals of this class of alloys is very difficult. Meanwhile, understanding material behaviour with decreasing sample size is necessary for exploring the potential application in nano‐technologies. In this study, electron backscattered diffraction (EBSD) was used to identify the grain orientations (e.g. close to [001], [011] and [111] directions) of a directional solidified Ni‐43Ti‐4Al‐2Nb‐2Hf alloy. Pillars with diameters ranged of 0.5 mm ~ 4 mm were machined using a focussed ion beam (FIB) microscope. Compression tests were performed using a Hysitron PI85 picoindenter. Specimens for transmission electron microscopy (TEM) were prepared along longitudinal direction of the pillars by FIB, and then examined in an FEI Tecnai F20 microscope.
The results show that its strength depends on the crystal orientations and the size of pillars (Fig. 1). The pillars were plastically deformed via dislocation gliding (Fig. 2) and twinning.