The measurement of individual crystal orientations in heavily deformed materials is of great importance especially for the study of recrystallization mechanisms. The electron backscattering diffraction (EBSD) technique which is frequently used nowadays in the SEM for these kind of studies has three important drawbacks: first, its spatial resolution is limited to about 0.5 μm, second, the sample deformation is limited to only about 60 %, and third, the dislocation microstructure cannot be observed. All these problems can be overcome using a micro beam diffraction technique in the TEM. Two techniques can be distinguished, the Kikuchi pattern technique and the spot pattern technique. With a spatial resolution of down to about 10 nm they arc well suited for orientation measurements in highly deformed material. Nevertheless, since spot patterns are less sensible to deformation the spot technique is the ultimate choice in the case a most highly deformed material. An automatic procedure for the evaluation of spot patterns has been developed and, together with the Kikuchi technique, implemented in a computer program for the on-line analysis of crystal orientations at the TEM. The new procedure and its applications to highly deformed materials is the subject of this presentation. Spot patterns are less sensitive to the sample deformation than Kikuchi patterns due to the differences in the diffraction geometry and due to the relaxation of the Bragg diffraction conditions for spot patterns.

There is a range of techniques now available for the study of crystallographic orientation in materials. Among these, Orientation Imaging Microscopy (OIM) for the SEM is being increasingly applied. This technique is, however, limited by the spatial resolution of the SEM (∼0.5 μm). TEM orientation measurements and imaging provides the necessary alternative to SEM measurements due to the much higher spatial resolution of the TEM (∼10 μm). In TEM crystal orientation measurements Kikuchi patterns are preferred over selected area spot patterns due to higher accuracy and spatial resolution. The earliest means of automation of TEM orientation measurements involved the use of digital scanners and beam deflection measurment for computerized semi-automatic indexing of the patterns. In recent years fully automatic on-line orientation determination has been developed. This has come about as a result of the advances in image processing for automatic indexing of backscatter Kikuchi patterns in the SEM, and the emergence of slow-scan CCD cameras with a large dynamic range.

A UHV diffraction camera has been built which can be used in either transmission or reflection mode. Fig. 1 shows a schematic view of the instrument. The optical axis of the instrument is vertical, with gun below and a downward-looking CCD camera on top. The pressure with gun operating is 1×10−10 Torr. An ion pump and oil-free turbo system is used. The electron gun is a fully bakeable custom UHV RHEED gun with LaB emitter and two magnetic lenses. This allows the beam to be focused to an area of about 1 micron on the surface of the sample. The gun works in the energy range of l-50keV. A double tilt stage, with sample exchange, has been built which allows azimuth rotation by 360° and tilt by ° 15°. The sample can be resistively heated on the stage up to 1300°C. The sample holder can also be used with 3mm diameter TEM grids for transmission diffraction. An external lens-coupled liquid-nitrogen cooled Princeton CCD camera with 1024×1024 pixels is used for image acquisition, viewing a transmission phosphor through a lead-glass window.

The energy filter consists of three spherical grids serving as a retarding field energy filter. The design is similar to that of Horio [1]. The first and last grids are grounded, while the center grid holds the retarding voltage. High voltage stability up to 10kV has been achieved with this filter. Fig. 2 shows an energy loss spectrum obtained.from the specular spot of the Si(111) surface. The spectrum was obtained from the measured integrated intensities by numerical differentiation.

Ternary and quaternary heterostructures from III-V-semiconductors get more and more importance in the fabrication of high-speed/high frequency devices in telecommunication systems. One of the key parameters for the performance of such devices is the crystalline quality and especially the amount of tetragonal distortion in strained layers on a nanometer scale. Strain can cause problems for the growth of such layers like relaxation induced defects, especially for the case of a high value of lattice mismatch. However, strain and the associated influence on the band structure can be used consciously for the design of tailor-made heterostructures (band gap engineering). Therefore, the spatially resolved investigation of local crystal properties (tetragonal distortion or strain and strain relaxation) is a key tool for the characterization of strained layers.

Convergent beam electron diffraction (CBED) patterns and convergent beam imaging (CBIM) can be used to evaluate informations on the local crystalline structure. E.g. the position of the High Order Laue Zone (HOLZ) lines in the CBED patterns is sensitive to the local strain and therefore can be used to determine strain and relaxation effects in heterostructures quantitatively. But in practice the applicability of CBED is often limited by a lack of ultimate spatial resolution and/or of sensitivity.

To have an understanding of many imaging and diffraction effects relating to high-energy electrons it is essential to take into account the vibrations of the atoms of the solid being examined. Einstein's model of the vibrations takes the atoms to be independent oscillators. Debye's model includes the effects of correlated atomic motion and is more accurate in describing simple solids. Electrons which exchange energy with the solid through a change in the vibrational state of the solid are incoherent with respect to the incident electron wave and may be considered as being absorbed out of the incident wave field. Yoshioka and Kainuma showed how to calculate the components of the absorptive potential, which is imaginary rather than real, and pointed out that they depend on the direction of the incident wave. The components of the imaginary potential vgh(im) are not simply a function of g-h, unlike those of the real potential describing elastic scattering. Here g and h represent vectors in reciprocal space. Using Debye's model and assuming that the speeds of the three components of acoustic waves in the solid are equal we find

It is well known that cubic SrTiO3 possesses a typical perovskite structure whose reciprocal unit cell is shown in Fig. 1 where X, M, and R designate particular points at Brillouin zone boundaries with the Γ point being the zone origin. The instability of the phonon mode designated by the R point leads to the cubic-to-tetragonal phase transition of SrTiO3 at Tc=103 K. By using the same unit-cell as the cubic SrTiO3, reflections (h k 1) / 2 with h, k, and 1 all odd are superlattice reflections characteristic of tetragonal SrTiO3. Previous investigations of diffuse scattering were carried out by neutron and xray diffraction at Temperatures below 240 K. In the present work, observation of the diffuse scattering was extended to a broad Temperature range of 20 ∼ 700 K.

Selected-area electron diffraction patterns (EDPs) reveal splitting and streaking of the superlattice reflections as shown in Fig. 2. We can see streaks along the [1 0 0]* direction, e.g., the streak connecting (5 3-5) / 2 and (7 3-5) / 2, the streak connecting (-3-5 7) / 2 and (-1-5 7) / 2. Splitting of the superreflections (7-1-1) / 2, (9-1-1) / 2, and (-7 1 1) / 2, along [0 1-1]* direction is evident. All the splitting and streaking behaviors shown in Fig. 2 and observed in other EDPs may be interpreted as being formed by the intersection of the Ewald sphere with reciprocal rods MRM shown in Fig. 1, which are along the <100>* direction passing through superlattice points.

Recently, it was reported that electron diffraction can be used to determine polymorphs in pharmaceutical industry [1]. The term polymorph refers to substances that have the same chemical formula but different crystal structures. Polymorphism has played an important role in drug development, production and litigation [2]. Synthesized and crystallized under different conditions, polymorphs may have different physical and chemical properties as well as bioavailabilities. The conventional techniques used to identify polymorphs .sometimes fail to differentiate definitively between two structurally similar polymorphs such as roxifiban, a very promising cardiovascular drug discovered recently at DuPont Pharmaceuticals [3,4]. Bulk and tapped densities of Form I and II, for example, were found to vary from 0.23 and 0.37 g/ml for Form I to 0.07 and 0.19 g/ml for Form II, respectively. These differences were of some concern given the low drug load in the solid dosage form. Unit cell parameters and atomic coordinates of crystalline polymorphs can routinely be determined by single-crystal xray diffraction techniques if suitably large crystals are available. However, many organic compounds tend to grow needle and plate-like crystals that never achieve a large enough volume for single-crystal studies. Such was the case with forms I and II of roxifiban.

Our investigations are aimed to exploit the potential of Electron Spectroscopic Imaging (ESI) using an LEO EM 912 Omega to perform element detection without any background signal included in the net elemental signal. In a first series of experiments we designed an approach for phosphorus (P) detection by using an internal mass-thickness marker in the shape of a carbon ramp combined with the two-window difference method in order to eliminate mass thickness signals [1]. However, pure carbon seems not to be ideal because this approach does not take into account the more complex elemental composition of biological samples ( H, N, O, P and S) compared to pure carbon within the reference. Consequently, we had to find an easy-to-handle and reproducible technique to include a combined Compositional- and Massthickness Marker (C:M-Marker) in the specimen under investigation. Fig. 1A shows the result of such an experiment. Turnip Yellow Mosaic Viruses (TYMV; bright particles) were layered on top of a carbon foil followed by a wash with distilled water and air drying. To form an internal C-M-Marker a 1%-solution of Bovine Serum Albumin (BSA), dialysed against H2O, was sprayed on the opposite side of the same sample resulting in localized aggregations of BSA with different thicknesses (bright region “M, Fig. 1A). It serves the same purpose like the carbon ramp [1]: a pre-edge image (115 eV) and a second one (150 eV, core edge image) are taken to calculate the intensity (I) background function I150 (I115) using only values from the C-M-Marker region. Then the P-map image can be calculated by subtracting the extrapolated contribution of the protein part from the core-edge image. For the case presented in Fig. 1A this procedure resulted in Fig. 1B where the TYMV intensity signals can still be seen whereas the BSA ramp vanished.

Several methods have been applied to the texture characterization of polycrystalline films, including X-ray diffraction and several SEM and TEM techniques. Few of these techniques combine high spatial resolution with the ability to characterize large sample areas in a short time. We have fabricated annular objective apertures for the TEM, extending earlier work of Heinemann and Poppa, to characterize the orientation and size of grains within polycrystalline thin films. This procedure is similar to conical dark field imaging which has been shown to provide microtextural information over large areas

Apertures were fabricated using a FEI FIB-200 focused ion beam (FIB) workstation. Bitmap templates of apertures were created using Adobe Illustrator™ and converted to stream files. These stream files were then used to controllably cut annular apertures into an existing platinum JEOL aperture strip with a 30 keV Ga+ FIB. The apertures were then placed in the objective aperture drive of a JEOL 2000FX transmission electron microscope and TiN thin films were imaged in bright- and dark-field imaging modes.

Phase identification using electron backscatter diffraction (EBSD) in the SEM has become a useful and important tool for the characterization of crystalline materials. Phase identification is accomplished using EBSD in the following manner. First, a high quality camera must be added to the SEM. Suitable cameras use slow scan CCD imagers coupled either by a lens or a fiber optic bundle to a phosphor screen that is situated near the sample. A EBSD pattern is collected and EDS or WDS is used to determine qualitatively the chemistry of the area. An automated routine is then used to extract the positions and widths of the lines in the pattern followed by a calculation of the unit cell volume. This information coupled with the chemistry of the sample is then used to search a database of crystal structures. Currently, the ICDD's Powder Diffraction file of over 100,000 compounds is used. Once a list of potential matches is found the patterns are indexed and then simulated to demonstrate that the phase has been identified. This paper will demonstrate use of EBSD for phase identification and then will speculate on future developments.

A particularly nice application of EBSD is the use of the technique for the identification of phases that form in welds. Figure 1a is an EBSD pattern obtained from a acicular phase in a superalloy weld. The phase was determined to be primarily Ti and Ni. Analysis of the patterns showed that the phase is Ni3Ti. Figure 1b shows the simulation for Ni3Ti overlaid on the experimental pattern demonstrating that the phase has been identified.

Electron backscatter diffraction patterns obtained from low symmetry crystals can present difficulties for automated pattern recognition. For example, when only a single zone axis is observed in the diffraction pattern and the Kikuchi bands used for indexing all pass through the zone axis then only two-dimensional crystallographic information is obtained. This results in there being a 180 degree ambiguity in the measured orientation. Furthermore, it is often observed that in automated indexing of quartz, an additional ambiguity is introduced when the zone axis imaged is [0001]. In this case it is difficult to distinguish between the true orientation and it’s twin. This ambiguity is overcome only if Kikuchi bands from the upper Laue zones are used in the indexing. However, in automated indexing it is normal for the computer to select the strongest Kikuchi bands visible and these will not be those from the upper Laue zones when the [0001] zone axis is imaged. This and other pseudo symmetries have been observed and have been overcome by applying correction procedures during data collection and in post collection analysis.

Texture determination of ceramic materials is important in ceramic processing since it will ultimately govern the resulting material properties. For instance, recent results have indicated that preferred particle orientation leads to anisotropic shrinkage of tape cast ceramics. In this work, we report results on the effect of pressing as well as shearing during tape casting followed by sintering on particle orientation and texture development in ceramic materials. Oxides with different particle shape anisotropy and crystal structures have been investigated including tetragonal rutile titania (TiO2) and hexagonal alumina (A12O3).

Ceramic samples were processed by either pressing or by tape casting of a high ceramic particle loading aqueous suspensions. The samples were then pyrolized at 425°C followed by sintering for one hour at 1200 and 1500°C for TiO2 and A12O3 respectively. High spatial resolution Electron Backscatter Diffraction (EBSD) (OPAL-Oxford) in a field emission SEM (LEO-982) has been used for texture analysis and for crystallographic orientation determination of particles. Between, 3000 to 5000 measurements per sample were recorded automatically to develop an orientation map.

Backscattered electron Kikuchi patterns (BEKP) were first observed by Alam et al. in 1954. J.A. Venables and C.J. Harland made the initial observation of BEKP and in the scanning electron microscope in 1973. In 1996, Goehner and Michael developed an electron backscatter diffraction (EBSD) system that uses a 1024 × 1024 pixel CCD camera coupled to a thin scintillator rather than photographic film. In their system, the quality of the raw patterns is improved by the use of “flat fielding” which normalizes the raw image to a “flat field” reference image that contains the image artifacts, including background, but not the crystallographic information. Automated pattern analysis is carried out using a Hough transform to locate bands and band edges in the pattern. The resulting crystallographic information is coupled with the elemental information from energy or wavelength dispersive x-ray spectrometry and the phase is identified is made through a link to a database such as the Powder Diffraction Files published by ICDD. An indexed pattern of the suspected phase is then synthesized for comparison to the unknown. This system is marketed commercially by Noran Instruments and offers the first practical method for rapid identification of unknown crystallographic phases in the SEM.

An important feature of polycrystalline materials is the orientation distribution of crystallites also known as crystallographic preferred orientation or texture [1]. Conventionally it is measured by x-ray diffraction, averaging over sample surfaces. With demands for more quantitative material characterization, both in engineering and earth sciences, new methods have been developed. Two dimensions are of interest: Averages over larger sample volumes give a better representation to estimate bulk physical properties. Here neutron diffraction is advantageous. Because of minimal absorption large sample volumes (1-50 cm3), rather than surfaces can be analyzed [2]. If textures are locally heterogeneous, it may be of importance to analyze small regions. With a synchrotron microfocus beam volumes as small as 5 μm3 can be characterized [3]. These methods have been quantified and are extensively applied in metallurgy and geology. They provide good statistics for appropriate sample grain size but they analyze bulk textures and contain no information on local orientation correlations. Furthermore, for geological samples with low crystal symmetry, the diffraction patterns are often very complex with many overlapping peaks, making identification difficult. In such cases orientation imaging, using electron backscattered diffraction patterns (EBSD) in the SEM is useful [4]. Though grain statistics are generally much inferior, new information can be gained.

Measurement of the geometry of triple junctions between grain boundaries in polycrystalline materials is used to generate large sets of dihedral angles from which maps of the grain boundary energy are extracted. A preliminary analysis has been performed for samples of magnesia and aluminum based on a three-parameter description of grain boundaries. An extended form of orientation imaging microscopy (OIM) was used to measure both triple junction geometry via image analysis in the SEM and local grain orientation via electron back scatter diffraction. Serial sectioning with registry of both in-plane images and successive sections characterizes triple junction tangents from which true dihedral angles are calculated. If there is local equilibrium at each triple junction, we may apply Herring's relation. By limiting grain boundary character to a (three parameter) specification of misorientation for the preliminary analysis, we can neglect the torque terms and apply the sine law to the three boundaries. This provides two independent relations per triple junction between grain boundary energies and dihedral angles. By discretizing the misorientation and employing multiscale statistical analysis on large data sets, (relative) grain boundary energy as a function of boundary character can be extracted from triple junction geometry. The results are discussed with respect to current understanding of grain boundary structure based on their crystallography. The results suggest that a three parameter characterization of grain boundaries (lattice disorientation) is not an adequate description of boundary character. A full analysis including torque terms and a five parameter boundary description is under development.

During plastic deformation of typical metals the original grains subdivide. Subdivision by cell formation is a well-known phenomena in metals with medium to high stacking fault energy, but the subdivision may also take place on larger length scales by: i) Formation of single-walled dense dislocation walls (DDWs) and/or double-walled microbands (MBs). The DDWs/MBs are seen as elongated, nearly straight, vertically parallel dislocation boundaries which delineate several cells, thus bounding cell blocks. Typical examples are shown in Fig. 1. ii) Subdivision on a grain scale, for example due to grain-grain interactions or due to the formation of transition bands separating matrix bands. This subdivision is typical for both single crystals and polycrystals. For a general description and recent overview, see [1, 2].

The grain subdivision depends on the orientation of the grains. The morphology may appear different in grains of different orientations and the DDW/MBs can be parallel or inclined to a slip plane.

Deformation and thermal processing operations can produce structurally heterogeneous and anisotropic alloys. Thick, hot-rolled aluminum plate is an excellent example of such a material. Due to the plate size and thermomechanical conditions experienced by the plate during processing, significant gradients in material structure arise. These gradients are especially strong in the through-thickness or normal direction (ND) of the plate. In AA 7050, an aluminum alloy developed for thick plate applications, gradients in microporosity, grain morphology and crystallographic texture are known to exist. These features have been linked to material properties and performance measures in this alloy. Recent advances in high speed machining technology have made it possible to consider thick AA 7050 plate for an increasingly large number of airframe components. Because of its inherent heterogeneity, however, designing flight-critical components machined from AA 7050 plate will require one to quantify and mathematically represent the variation of the relevant microstructural gradients.

This work describes the investigation of plastic strain in tensile deformed polycrystalline copper using the automated Electron BackScatter Pattern (EBSP) technique of electron diffraction in the Scanning Electron Microscope (SEM). A strong correlation was found between the known accumulated plastic strain in OFHC copper tensile specimens and quantitative measurements of intra-grain misorientation obtained from an analysis of automated-EBSP data.

Previous works on strain measurement using EBSP analysis have focused on the influence of dislocation density on the intensity profiles of diffraction bands. It is known that for some materials, like copper, deformation strain produces highly organized dislocation structures and thus the dislocation density is very inhomogeneous on the sub-micron scale probed by EBSP. This suggests that plastic strain analysis using the EBSP dislocation density analysis route will be difficult to reliably implement in these materials.

Bainitic low alloy steel has a complex microstructure exhibiting several types of boundaries. The boundaries in bainitic steel, although certain boundaries are absent with respect to the alloy composition and the manufacturing process, could be typically divided into 4 types; dislocation cell boundary, lath boundary, packet boundary, and prior austenite grain boundary, in increasing order of size. The size and distribution of the respective boundaries are an important factor which controls the mechanical properties of the steels, including brittle fracture. In the present research, the characteristics of the boundaries in the bainitic low alloy steels were investigated in view of misorientation between grains enclosed by the respective boundaries.

The alloys investigated were Mn-Mo-Ni low alloy forging steels having chemical compositions shown in TABLE 1. Steel-A was manufactured by the Vacuum Carbon Deoxidation(VCD) process. For the finer prior austenite grain size, Steel-B was produced by the aluminium addition and the silicon killing process. Before EBSD analysis, the microstructures of the alloys were observed using SEM and TEM. EBSD measurements were obtained using a Link OPAL system(Oxford) linked to a JEOL JSM 6300 SEM operating at 15KeV with the sample tilted at 70°.

A standard commercial EBSD system, comprising a CCD camera (Nordif) and acquisition software (HKL Technology) has been interfaced to a Philips XL30 FEGSEM, and the spatial and angular resolutions for EBSD compared with a conventional W-filament JEOL 6300 SEM equipped with a similar EBSD system.

In characterising a grain or subgrain structure, overlapping diffraction patterns are obtained at the boundaries, and the effective spatial resolution can be defined as the distance over which the patterns cannot be solved. The spatial resolution is best in the direction parallel to the specimen tilt axis (LA), typically some three times better than for the direction perpendicular to the axis. The resolution is conveniently determined by measuring the fraction of patterns which are solved (Ns) during a raster scan of the sample. For a microstructure of equiaxed grains of size D, LA is given approximately by D(1-Ns)/4. Figure 1 is a plot of Ns against D−1 for aluminium samples under optimum operating conditions, and it is seen that the spatial resolution in the FEGSEM is some three times better than in the W-filament SEM.

The effect of beam current on LA is shown in figure 2. For the W-filament SEM, LA is a strong function of the probe current.