The technique of high-resolution transmission electron microscopy (HREM) produces images that contain information about the atomic structure of the specimen. Within additional, very stringent, constraints, the HREM image can contain information about atomic structure of crystal defects, including grain boundaries and interfaces. to extract this information from the image it is necessary to compare the experimental image with a simulated image calculated based upon an atomic model of the specimen.2 in this comparison, investigators have been aided by the use of quantitative techniques.

Atomistic simulations are often used to predict the atomic structure of crystal defects or to simulate the evolution of dynamic processes in crystals, e.g. radiation effects or dislocation motion and interaction. During the development of new models of interatomic interactions, the predictions of simulations are tested against experimental observations to validate new potentials. Grain boundary structure is a good test because atoms residing in the boundary experience environments (interatomic distances and angles) that are significantly different from those experienced by atoms residing in a perfect crystal lattice site.

The phenomenon of segregation is of long standing scientific interest and has been studied extensively, both theoretically as well as experimentally. For our investigations we have chosen the Σ5 symmetric tilt grain boundaries (STGB) in two face-centered cubic (FCC) metals, aluminum and copper. Both metals were doped with only 1 at% of the impurity species (Cu and Ag). One of our major goals in this study was to investigate the size effect on segregation of an impurity to distinct sites at the grain boundary. Therefore we have selected the Ag as an impurity in Cu and Cu as an impurity in Al. The latter one is of special interest for applications like interconnects in microcircuits, where one of the major controlling factors of electromigration is expected to be the diffusion or segregation of Cu atoms at Al grain boundaries.

The model grain boundaries have been fabricated with ultra-high vacuum diffusion bonding of single crystals which bonds the bicrystals under highly controlled environmental conditions.

α-sialon is the isostructural derivative based on the a-silicon nitride (α-Si3N4) structure by partial replacement of Si by Al and N by O. The structure is considered to be stabilized by the incorporation into the large interstices (r=0.146nm) of a metal like Li, Mg, Ca and most small rareearths. Large rare-earth metal ions have been considered not to be able to enter the interstices in α-sialon structure until several recent successes in accommodating large cerium ions (r=0.103nm). This gives rise to questions of the location and distribution of these large cerium ions within the asialon structure. It has been found that α-sialon grains formed in Ce-Si-Al-O-N system are usually rich in a new structural defect which has been considered to be initiated specifically by the presence of cerium in the lattice. This new intergrown defect is identified as domain boundary in our present study by transmission electron microscopy (TEM).

The understanding of the mechanism of oxidation of Zircaloy materials provides an important support for the corrosion control of the fuel claddings in the light water nuclear reactors. Many investigations are devoted to the study of the oxidation of these materials. One of the important aspects of these studies, is the analysis of the metal-oxide interface, which produces information about the nature of the oxide formed at the interface, at different stages of oxidation and the influence of the oxide structure and morphology on the formation and growth of the oxide.

In the present study, analytical electron microscopy (AEM) and high resolution transmission electron microscopy (HRTEM) are used to examine the metal-oxide interface of an un-irradiated Zircaloy-4 material, oxidized in autoclave, under pressurized water reactor conditions.

The TEM specimen preparation for the interface analysis is an important step of the investigation, since the transverse section required for such observation should be sufficiently thin exactly at the position of the interface.

Epitaxial ZnO thin films were produced on basal plane ScAlMgO4 (SAM) by chemical solution deposition (CSD) (fig. 1). ZnO crystallizes in the wurtzite structure and ScAlMgO4 is isostructural to YbFe2O4, a layered structure shown in fig. 2. The crystallographic orientation relationship was observed by XRD using θ-2θ scans and off-axis φ scans to be (0001)zno ∥ (0001)SAM and zno ∥ SAM. The lattice mismatch between film and substrate is 0,09%.

For TEM observations in cross-sectional geometry two substrates were glued together with the epitaxial film surfaces. Electron transparent specimens were produced by standard methods of mechanical thinning and ion beam etching. The HRTEM investigations were performed on a Philips CM300 FEG electron microscope with Ultra-Twin lens operated at 300 kV. After correction of axial coma and two-fold astigmatism images were recorded on a slow-scan lk × lk CCD camera with a sampling of 0.0193 nm/pixel.

For over 30 years geologists have been trying to better understand antiphase domains (APD) and boundaries (APB) in pigeonite in hopes of using them as markers for the thermal history of the rocks in which they are found. The ability to know the cooling history of igneous rocks is of great interest to geologists and pigeonite has received special attention on this matter because it has exsolution (precipitation) and antiphase domains (APD), both of which can be used as possible thermal markers. APDs in pigeonite arise because of the C2/c → P21/c transformation that occurs upon cooling. When multiple APDs nucleate, grow, and impinge upon one another, they are either in registry or have a translational discrepancy of ½(a+b). The size of the APDs can be used as a qualitative marker of cooling rates, since slowly cooled pigeonites favor large APDs and rapidly cooled pigeonites favor small APDs.

A TiN thin film was grown by the DC-Sputtering technique was characterized by EELS spectroscopy in order to obtain the optical quantities, ε1,ε2 and the refractive index η. Optical properties were obtained analyzing the low loss window located at an energy less to 50 eV respect to the elastic peak. A primary energy of 200 keV and a spot size of 440 nm with an average acquisition time of 5 min were used. The optical properties were calculated using the Kramers - Kronig relations.

Figure 1 shows the PEELS spectra in the low loss region for TiN thin film. As we can see, the most prominent plasmon loss peak characterizing TiN appears at an energy loss of 24 eV. The origin of the loss peak located at 17 eV is interpreted as surface plasmon. in order to obtain a better approximation, we use a near-free electron model and we have compared with the respective optical quantities calculated with the conventional formalism.

Vanadium carbide was purchased as powder at Aldrich with a 99.9% pure. Sample was prepared in the standard method for powder observation in a TEM .EELS experiments were carried out in a Phillips CM 200 STEM equipped with a Gatan 766 PEELS spectrometer. Experimental conditions for acquiring were the follows: a spot size of 500 nm, a chamber length of 400 mm and a detector aperture of 3 mm using a energy dispersion of 0.3 eV/channel with a beam energy of 200 KeV. Acquisition time was around 10 mn. taking an average of 100 spectra.

All the EELS spectra were corrected for background subtraction and multiple scattering deconvolution, before the fine structure analysis. The fine structure range was about 200 eV beyond the C-K edge located at energy loss of 283.8 eV and respect to the VL23-edge located at a energy loss of 513 eV.

Titanium Nitride thin film was obtained by the DC-Sputtering technique and characterized using the transmission EXEELFS formalism . The experiments were performed in a CM 200 transmission electron microscope with a GAT AN 766 PEELS attachment. A primary energy of 200 keV and a spot size of 440 nm with an average acquisition time of 5 min were used. The atomic distribution obtained using EXEELFS was corrected with the respective experimental phase shifts. Atomic parameters were compared with those corrected by EXAFS phase shits and the known crystallographic values.

The energy loss spectra for the sample, around the ionization Ti L23-edge and N K-edge corrected for background and multiple scattering process is shows in Figure 1. N K-edge appears at an energy loss of 400.4 eV and the Ti L23-edge is located at an energy loss of 464.3 eV. These values on energy are shifted 0.6 eV respect to the isolated transitions.

Ceramic of ZrB2 and its composite of the ceramics with SiC are used for ultrahigh temperature protection coatings in re-entry space vehicles because of high melting temperature of ZrB2 (3245°C). Chemical composition and microstructure of the ceramics are critical to thennal shock properties of the ceramics. Especially oxidation of boron and formation of boron oxide will degrade property of thermal resistance. Conventional TEM with Energy dispersive X-ray spectroscopy (EDS) is unable to study light elements like boron. However, parallel-recorded electron energy-loss spectroscopy (EELS) and energy-filtering GIF system associated with field-emission gun (FEG) TEM are able to carry out chemical imaging and chemical information of light elements. Bonding structure and oxidation states of an element (such as B) can be obtained by using EELS.

Figure A is a zero-loss bright-field TEM image showing an interface between ZrB2 and SiC. Figure B is a Plasmon-loss filtered TEM image of the same area. A plasmon peak at about 20 eV was used for this image.

The excellent mechanical properties of Ni based superalloys depend upon the presence of γ’ particles (LI2 structure). Their volume fraction, spatial distribution and size determine the mechanical strength of these alloys. Ni alloys for technological applications make use of large volume fractions of precipitates where processes of coarsening and coalescence take place during service leading in some cases to deterioration of properties. Addition of different alloying elements prevents accelerated coalescence by retarding diffusion and thus improving the mechanical properties of such alloys. Coalescence can also take place under the influence of an applied stress leading to the formation of rafts of the y' phase. For example the microstructure changes during creep deformation, depending on the alloy composition, with the corresponding formation of dislocation networks and rafts of different morphologies [1]. The γ-γ’ interfaces are also different depending on the alloy composition and most likely to the local distribution of alloying elements and their strain fields.

Recently LACBED (Large Angle Convergent Beam Electron Diffraction) studies for identifying the nature of stacking faults has been reported in [1,2]. Here we report the LACBED study for a microtwin whose images are usually similar to those of an intrinsic or an extrinsic stacking faults (for this discussion, see [3]).

Observations: Thin foils of cobalt with the thickness of about 180 nm (f.c.c phase, a=0.354 nm) were examined by a Philips CM200. Fig. 1 shows strong beam dark field images of microtwins or stacking faults. Fig. 2 shows the bright field LACBED pattern taken near the area marked as a circle in fig. 1. The specimen height, from the convergent point of beams, was about 0.0586 mm and the convergent angle was 0.615 degrees.

Calculations and analysis: Analysis of fig. 1 alone indicates the encircled fault an extrinsic stacking fault.

Recent experiments with X-rays and high energy electrons have shown that image recovery from diffracted intensities is possible for non-periodic objects using iterative algorithms. Application of these methods to biological molecules raises the crucial problem of radiation damage, which may be quantified by Q = ΔE σi/σe, the amount of energy deposited by inelastic events per elastic event. Neutrons, helium atoms and low energy electrons below most ionization thresholds produce the smallest values of Q (see for TMV imaged at 60 eV). For neutrons (λ = 10-2Å, and deuterated, 15N-abelled molecules) Q is ∼3000 times smaller (∼50 times for λ = 1.8Å) than for electrons (80- 500keV) and about 4x 106 times smaller than for soft X-rays (1.5Å). Since σe for neutrons is about 105 times smaller than for electrons (and about 10 times smaller than for soft X-rays), a 105 times higher neutron dose is required to obtain the same S/N in a phase contrast image compared with electrons, if other noise sources are absent.

Three-dimensional reconstruction in TEM is an established technique in both electron diffraction crystallography and electron tomography [1,2]. By stereo lattice imaging, we refer to a different task: determination of unit cell parameters in 3D using only direct-space lattice images at two or more (often widely varying) orientations [3,4], This note concerns the uncertainties associated with lattice parameter determinations by this technique. in particular, for decreasing crystal thickness the broadening of reciprocal lattice spots increases one’s chances to visualize lattice fringes. At the same time, however, it introduces uncertainties in reciprocal lattice vector length and direction.

In the following discussion, sources from which uncertainties arise are assumed to be uncorrected. Three sources of uncertainty for reciprocal lattice vectors which lie in the plane of the specimen stage (and hence can be imaged directly as fringes) are, in order of decreasing relative size as follows:

The phrase “fringe visibility map” here refers to high-resolution images of uniform-thickness single crystal foils showing locally hemispheric deformation (i.e. bent into the shape of a watchglass), and to various mathematical analogs thereof. As the availability of “diffraction information” in direct-space form increases, for example with the availability of HREM and z-contrast lattice images, fringe visibility maps may serve as the direct space analog to Kikuchi maps, i.e. to diffraction pattern maps of Kikuchi line pairs as a function of specimen orientation, and their various cartoonifications.

Figures 1 to 3 are schematic Mathematica-generated fringe visibility maps for face-centered, body-centered, and diamond face-centered cubic lattices, respectively. Here we have turned the “locally hemispheric” concept around and drawn fringes visible when one is viewing a spherical particle of the appropriate thickness, along varying beam directions (rather than lattice orientations) parallel to the local radius vector on that sphere.

Introduction: Modern TEM/STEM instruments with field emission gun and various analytical equipment allow the measurement of interfacial properties by a range of modes and techniques without exchanging the specimen. It is therefore useful to think of a criterion to assess the sensitivity of each technique, e.g. for the location of light atoms, by a commonly applicable formula.

Sensitivity measures: A “quantitative” TEM experiment is defined here as consisting of a digital comparison (e.g. via data difference or cross correlation) between experimental (E) and simulated (S) microscopy signal. The matching quality R = f(E,S) is to be optimised (minimum for data difference, maximum for cross correlation). The target quantity of the Q-TEM experiment is assumed to be the location or occupancy of a light atomic column in projection surrounded by heavier atoms in the core of an interface making use of dynamical enhancement effects.

Grain boundaries of covalent bonding materials such as silicon, germanium and SiC require rather physical consideration than that of geometry to discuss the result of their electron microscopic observation. Dangling bond, for example, is expected least in number in the stable boundary (least dangling bond rule). [1][2] in practice, further precise consideration must be done both in geometry and in physics in the covalent bonding materials boundary. in the present study, geometrical discussion is combined together with that of the atomic bonding about diamond, SiC and ZnO grain boundaries. Chemical structure image of the modern high-voltage transmission electron microscope enabled to do it. An asymmetrical Σ3 boundary of diamond and a 6H-SiC boundary with 70.5 degree rotation are mainly observed.

The sliced specimen, 0.3mm in thickness, was mechanically thinned down to 0.02 mm and was finally polished using an argon ion milling. Final thickness of the specimen was estimated to be 3 nm in the region close to the edge of the small hole.

Grain boundaries in fee metals with low stacking fault energy are known to undergo extended relaxations that can at times lead to a thin layer of a different structure. in Cu, for example, it has been found that Σ3{112}} boundaries relax into a 9R phase or even a bec phase. in this work, we have used high resolution electron microscopy to investigate the atomic structure of Σ99 grain boundaries in mazed bicrystal films of Au. On {001} Ge single crystal substrates, Au bicrystals can be grown in two equivalent {110} orientation variants, related to each other by a 90° rotation about the surface normal. There is equal probability to nucleate either orientation variant in early stages of film formation and since there is no reason for preferential growth, the morphology of the bicrystal film consists of many grains in only two orientations. All grain boundaries are of the type Σ99 and display tilt character.

Thin films 10-200nm thick were grown by physical vapor deposition on {001} Ge substrate at 280°C and followed by annealing at 200°C to relieve stresses.

A variety of properties of polycrystalline TiO2, such as conductivity, creep and grain boundary diffusion, are strongly dependent on the atomic-level structure and chemistry of grain boundaries. TiO2 has been studied as a model dielectric system because of its relatively simple structure and well-understood point defect chemistry. Defect segregation in grain boundaries of polycrystalline and bi-crystal TiO2 have been studied by several groups and significant variations in solute segregation levels from boundary to boundary were observed. in this paper, we address this issue of anisotropic grain boundary segregation. We have measured solute segregation as a function of grain boundary misorientation to determine any correlation between segregation and misorientation.

Yttria-doped TiO2 polycrystalline samples were prepared by mixing 99.999% pure TiO2 powder with 0.1% mole percent 99.99% purity Y(NO3)3 (both powders are commercially available from Aldrich Co. or Alfa Co.), followed by uni-axial pressed to 200-400Pa and sintering at 1300 °C for about 5-7 hours.

We report a novel technique to accurately measure interfacial lattice displacement by forming an electron probe close to a specimen in a manner similar to that originally proposed by Gabor to record a hologram. This method is based on the quantitative analysis of the interference pattern of shadow images in coherent electron diffraction. The approach is unique in that there are no adjustable microscope parameters, the contrast is strong even when the fault is viewed edge-on, and a large number of shadow images of the fault corresponding to different Bragg reflections can be studied simultaneously. Since it is an interferometric technique, the spatial resolution of the measurement is not limited by the wavelength of the fast electrons. An accuracy down to 1pm has been demonstrated in measuring the displacement associated stacking faults and grain boundaries in superconductors. It is, to our knowledge, the highest that has been achieved in measurements of displacement vectors.