We have discovered that high-frequency ultrasound imaging, or ultrasound biomicroscopy, can be used to detect apoptosis in a number of experimental systems. We have shown that such detection with 30-40 MHz ultrasound is possible using cells in an in vitro system (AML-3 leukemia cells) made to undergo apoptosis in response to treatment with a variety of cancer killing chemotherapeutic drugs. We have shown that ultrasound biomicroscopy can also detect programmed cell death in tissues made to undergo apoptosis in response to photodynamic therapy, currently an experimental cancer treating regimen. Lastly, we have shown that this ultrasound imaging approach works in vivo, using living animals where apoptosis has been induced similarly using photodynamic therapy. Specifically, apoptotic cells and regions of apoptosis in tissues exhibit up to a 36-fold increase in ultrasound backscatter intensity permitting this type of cell death to be readily discriminated from surrounding viable tissue.

One of the common problems with current remote control systems is the quality of the digital image that reaches the remote user. The size, resolution, refresh rate, and latency all contribute to the quality of the remote users’ experience. We are working on a series of remote control and display systems that send images over dedicated and production networks to better understand the bandwidth and quality of service issues involved. The amount of bandwidth required is a function of compression, image size and frame rate. Different operations require quite different parameters. Focusing and astigmatism correction require faster frame rates but can be performed with reduced size. Imaging can stand slower frame rates but requires high resolution. The quality of service (QoS) is a measure of latency and how consistently the network delivers the required bandwidth.

“In-situ” TEM studies is currently going through resurgence for materials microcharacterization. One important aspect of these studies has always been the image recording process. Traditionally, the following analog image recording techniques have been used: i) “static” images taken by photographic films and ii) viewing the images with a TV camera and recording them onto a videotape [1]. On the other hand, with the advent of CCD cameras and its rapid evolution for digital “static” imaging in TEM [2-3], “real-time” digital imaging for “in-situ” experimentation has been in high demand. In this work, we report the utilization of a new CCD camera for both “real-time” and “static” image acquisitions. The performance of the camera is shown for heating/cooling “in-situ” TEM experiments performed on an aluminum alloy using a double tilt hot stage.

We have utilized a newly designed CCD camera that addresses imaging needs for “in-situ” applications.

Modern SSCs combine low noise with high sensitivity and provide digital images with high dynamic and excellent linearity. Since these images are immediately available in a PC, SSCs have become a non-dispensable tool for on-line image processing and software driven automated TEM tuning [1]. Many experiments performed nowadays in the area of high resolution, low dose, holographic reconstruction, and EFTEM [2] would not be possible without the digital input coming from SSCs.

Despite their excellent characteristics SSCs provide a restricted number of individual image points in respect to film. This turned out to be the most prominent disadvantage of SSCs especially for applications where highly resolved large specimen areas need to be captured. Tab. 1 shows a typical example for a biological application where an image detail of 4 nm (plasma membrane) needs to be resolved. Depending on the film material used, this resolution is matched more or less by a TEM magnification of ∼8,000x.

MacLispix, a public domain image processing system for the Macintosh, has been applied to a variety of image processing problems such as analysis of diffraction spots, uniform display of x-ray maps, determination of fractal dimension of particle outlines, and analysis of data cubes. Due to interest from the PC community, we have ported the software to Windows, renamed it Lispix', and distributed it for both platforms, along with example images, source and documentation.

Lispix reads TIFF files and raw files (no image header), both with pixel types of signed and unsigned 8, 16 and 32 bit integers, 4 and 8 byte IEEE standard floating point numbers, and 3x8 bit RGB color. Lispix has a variety of standard image processing operations, such as thresholding, edge finding (gradient), filtering, scaling, Linear Hough Transform, false coloring, RGB color overlays and particle measurement.

A complex three dimensional network structure results when blood fibrin polymerizes into branching threads. To understand the rheological behavior of fibrin clots and how this behavior is associated with network morphology, one must obtain quantitative measurements of morphometric parameters of the networks (e.g., fiber diameter, distance between branching points of fibers, and branching complexity) formed under various conditions. Over the past several years, researchers have measured the parameters manually from the 3D reconstruction of networks using stereo images generated with intermediate voltage electron microscopy. Our goal is to help automate this time consuming process as much as possible, and skeletonization is a promising tool. Skeletonization is a process of transforming an object into idealized thin lines. Although various techniques for skeleton computation have been developed, most of them require explicit boundary information of the objects being skeletonized, which is not directly available from raw images.

Position normalization is a method of manipulating data from any image source so that the effective information to noise ratio is dramatically increased. In this paper we present two very different examples of the application of position normalization used in the study of high current density superconductors. In the first example we use FESEM fractography to analyze grain boundary density and grain shape in micron-diameter filaments and then normalize the data obtained from individual grains to their positions with respect to the filament-matrix interface. In this way we extract the change of grain structure with position in apparently inhomogeneous microstructure. In the second example we analyze the backscattered electron intensity, BEI, from the cross-section of a superconductor filament with respect to its proximity to phase boundaries. In this way we extract small trends in the mean atomic number with a spatial resolution of better than 100 nm.

Many scientific questions encountered in electron microscopy require quantitative deductions from the observations. Comparisons of experimental observations with simulations are however, still relatively rare since measurements of intensity are normally difficult. In this paper we discuss the use of experimental observations of the effective extinction distances for zone axis silicon using a convergent beam STEM mode for comparison with a number of simulations. On the experimental side, the measurements were made with a STEM that provides accurate intensity measurements directly with a digital image recording system. Two theoretical schemes widely used in electron microscopy simulations, multislice simulation and Bloch-wave calculation, were employed for the simulations. In each case, both a TEM case and a STEM case were calculated for comparison.

The multislice simulations were carried out using codes available from Kirkland. For the TEM case with plane waves at normal incidence on the sample surface, the unscattered (0,0) exit beam gives the Bright Field (BF) intensity.

Al-Cu-Mg based alloys are of significant interest for aerospace and other applications, due to their low weight, mechanical strength and corrosion resistance. Their mechanical properties are based on a dispersion of S-phase precipitates, which have the composition Al2CuMg. The crystal structure of this inter-metallic phase has been studied using different diffraction techniques for more than five decades. While several models have been proposed for the structure of S-phase[l], only one of the previously proposed models were found to give a reasonable fit to our experimental data. This model, shown in Fig. 1 and given by Perlitz and Westgren (PW) [2] based on X-ray diffraction data, is orthorhombic with unit cell dimensions a = 0:4 nm, b = 0.923 nm, and c = 0.714 nm, space group Cmcm, containing 16 atoms in the ratio Al:Cu:Mg = 2:1:1.

Several HREM images of S-phase precipitates located near the edge of the foil, Fig. 2, recorded along the [100]s and [010]s directions, were digitized from film and used for analysis.

Recently out-of-focus, low angle electron diffraction and Foucault experiments have been carried out on superconducting specimens in a range of applied magnetic fields where the fluxons form a more or less regular lattice. Let us recall that in order to describe the effect on the electron beam in a transmission experiment of a fluxon in a tilted specimen, the fluxon itself has been approximated by a suitable bundle of straight flux tubes, relying on the important result that the phase shift of the flux tube can be calculated analytically even in the tilted specimen geometry. In this way, it is only necessary to convolute this phase shift with the chosen projected magnetic field distribution (in our case a London model with a phenomenological penetration depth of 50 nm) in order to obtain a fairly realistic model. Therefore, in order to interpret the main features of the experimental results,

For magnetic flux mapping in magnetic materials it is important to use a simple, reliable and quantitative method operating with resolution down to few nanometers scale. The traditional off-axes electron holography can reach, in principle, such a resolution. However, the time-consuming phase reconstruction procedure limits its applications to real-time dynamic studies of magnetic domains behavior under varied external field. A novel recently developed in-line holographic approach, based on magnetic interferograms generated by phase-coherent Foucault (PCF)-imaging, complies with most of these demands and is suitable for dynamic studies of domain structure. For instance, it was successfully applied for induction mapping of closure domains in soft Fe-films [1]. It was found, however, that the resolution of PCF-imaging was lower than that of off-axis electron holography, since it is limited by the observed fringe spacing [2].

We report that the real resolution of PCF-imaging can be improved significantly with help of a novel phase amplification method.

High-resolution electron microscopists are familiar with the concept of an “optimum defocus” for obtaining highresolution transmission electron microscope images. Scherzer recognized that it is possible to balance the phase changes imposed by the spherical aberration of the TEM objective lens by adjustments to lens defocus. Selection of this focus condition maximizes the same-phase transfer of structural information carried by electrons scattered from the specimen. The upper limit of spatial frequencies transferred with the same phase change determines the resolution of the microscope. The resolution and “optimum defocus” depend only on the electron wavelength of the microscope and the spherical aberration coefficient, CS, of its objective lens. Reduction of Cs is the major route to improved resolution.

With the advent of electron-optical systems able to generate negative spherical aberration (usually called “Cs correctors“), it has now become feasible to zero-out objective lens Cs in the high-resolution transmission electron microscope.

As a crystal gets smaller, diffraction spots stay visible for larger deviations from the Bragg condition. The upper limit of such deviations is connected to the threshold for getting lattice fringes in TEM images. This in turn allows one to quantify the probability of seeing cross lattice fringes along a certain zone axis. In this abstract we examine a simple semi-empirical model for the probability of detecting (001) zone cross-fringes of a spherical crystal of cubic lattice.

The upper limit for the deviation of crystal orientation from the exact Bragg condition, without losing cross fringes down a given zone, is expressed as the maximum half-angle θ1 between the zone and the electron beam. The solid angle σ subtended by a cone with this half-angle is proportional to the probability px that a randomly-oriented crystal will show the cross-fringes associated with that zone. A schematic, illustrating the principle used to calculate the probability of seeing cross-fringes, is given in Figure 1.

Factors that affect the visibility of lattice fringes include crystal orientation and thickness, as well as instrument response. As a crystal gets smaller, lattice fringes stay visible for larger deviations from the Bragg condition. Hence the persistence of fringes under tilt affects the abundance and range of lattice spacings (and angles) that one sees in an image of randomly-oriented crystals. A subset of the fringes in an image are “still-visible” after large (e.g. 35°) single or double axis tilts. If one is looking for “new-fringes” from the same crystal (e.g. to analyze it's 3D lattice parameters), rules for recognizing redundant fringes might also help out.

Here we examine a semi-empirical model for predicting the visibility of lattice fringes after tilt, by connecting the visibility to intersection of the corresponding crystal reciprocal-lattice spot with the illuminating Ewald sphere.

Creep mechanism determination on dental amalgams has been a subject of interest in recent research works. Most of these studies conclude that the creep behavior can be explained by the well-known expression: ɛ = A σnexp (-Q/RT). Reported values for n (potential coefficient of stress) are in the order of 3 for phase dispersed dental amalgams and 2 for conventional amalgams. It suggests a grain sliding creep mechanism for conventional amalgams and a climb dislocation mechanism for creep in phase dispersed amalgams due to interactions with the dispersed CuxSny particles typical of this kind of amalgam(l). For both cases, the creep process occurs inside the γ (Ag2Hg3) phase matrix of low melting point (310 °K) which is slightly higher than typical mouth temperatures of about 400 °K (37 °C). In this work some microestructural evidences are obtained to support the behavior described above by using High Resolution Transmission Electron Microscopy (HRTEM).

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 ∑3﹛ 112﹜ boundaries relax into a 9R phase [1]. In this work, we have used high resolution electron microscopy to investigate the atomic structure of ∑3 grain boundaries in mazed bicrystal films of Au. Using ﹛111﹜ Ge surfaces as a template, Au bicrystals can be grown in two orientation variants, related to each other by a 60° rotation about the surface normal. As described previously, such films have a strong tendency to facet onto the coherent twin plane parallel to the substrate [2], also known as “double positioning” [3]. If films are made very thin, the likelihood for such in-plane boundaries to lie in the foil decreases, and it becomes possible to observe the atomic structure of edge-on interfaces along <111>.

Ag-Ge forms a simple binary eutectic with only two solid phases - a Ag-rich fee phase and a Ge-rich diamond-cubic phase, both with very limited solid solubility for the other element. Because both phases are cubic and there are no intermediate compounds, precipitation of Ge in an Al matrix might be expected to be crystallographically and morphologically simple. However, previous investigations of similar alloys (Al-Ge and Al-Si) have shown that precipitation in such apparently simple eutectic systems can be surprisingly complex [1,2]. For example, it was found that Ge precipitation in Al-Ge alloys depends sensitively on the availability of lattice vacancies and that at least five different orientation relationships and even more types of morphologies can form during aging after a quench. In the course of that work it was discovered that twinning is a key factor in the development of precipitate morphologies. To understand the role of twinning in precipitation this study investigates the characteristics of precipitation in Al-Ge alloys.

Spectrum imaging, where a full spectrum is acquired at each pixel in a two-dimensional array, provides a new paradigm for materials characterization that combines the advantages of traditional paradigms for imaging and analysis in electron optical characterization. Traditionally, electron optical imaging (e.g., secondary electron imaging of electron-opaque or bright-field imaging of electron-transparent specimens) emphasizes contrast production among distinct features of the microstructure. An excellent survey of the specimen microstructure is achieved - many pixels are acquired - but typically no quantitative analysis of image intensities is performed. In contrast, electron optical analysis techniques traditionally concentrate on quantitative elemental analysis (e.g., electron probe microanalysis), or at least comprehensive qualitative analysis through full spectrum acquisition (e.g., energy-dispersive X-ray or electron energy-loss spectrometry). However, these analyses are typically performed at only a few locations on the specimen, which are pre-selected on the basis of the image contrast generated as described above.

Many crucial measurements in the semiconductor industry involve determining the root cause of an electrical failure, often requiring the capability of extracting microstructural and chemical information with nanometer resolution [1]. The microanalysis is achieved by stepping the focused probe over the region of interest in STEM mode and recording an EDX and EELS spectrum at each pixel. Even for relatively modest image sizes, the resultant spectrum-image may consist of more than 10,000 spectra. With such large data sets, the prospect of manually inspecting and quantifying each spectrum in real-time is impossible and even automatic background modeling followed by peak area integration over selected spectral regions-of-interest can be excessively time consuming. Other methods to extract useful physical information include “fingerprinting” characteristic ELNES shapes to known compounds or bonding environments, multiple least squares regression to a relatively small set of suspected components or peak-fitting to model systematic changes in electronic density of states.

As more x-ray energy dispersive spectroscopy (EDS) manufacturers begin to offer spectrum imaging (a complete x-ray spectrum from each pixel in an image), there is a clear need for robust and automated methods for quickly extracting the relevant information from the large spectrum image data sets. A typical spectrum image may consist of 100 x 100 pixels (10000 spectra) each with 1000 channels which (when stored at double precision) is 80 Mbytes. It is clear that a large four-dimensional data set such as this cannot be viewed in its entirety and the time to analyze individual spectra by hand is prohibitive. Conventional analysis of spectrum images by mapping energy windows is useful as a first pass only for finding the elements present and only if at sufficient concentrations. Additional problems with mapping include systematic overlaps of other x-ray peaks, changes in the background shape and displaying the maps so they faithfully portray the actual signal intensity.