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The particle size distribution (PSD) and particle morphology of metal powders undoubtedly affects the quality of parts produced by additive manufacturing (AM). It is, therefore, crucial to accurately know the PSD and morphology of these powders. There exist several measurement techniques for these quantities, but since each method is based on different physical phenomena, which are sensitive to different aspects of a particle's shape and size, it is unclear how the measured PSDs and morphology compare to one another. In this study, five different techniques are used: sieve analysis, dynamic imaging analysis, laser diffraction analysis, X-ray computed tomography (XCT), and scanning electron microscopy. The first three are commonly used in the powder metallurgy field while the last two are laboratory-based tools capable of providing robust size and shape data. Nominally identical samples of stainless-steel powders were produced via riffling, and each technique was employed to measure effectively the same PSD and in some cases the morphology. In this paper, the differences among these measurement techniques are explored by a comparison of the measured results. Besides the random variations of the various measurement processes, the difference in the results is partly due to the fact that the particles are not perfectly spherical and that there are many multi-particles present. Each of these affect the principle of each method differently. Three-dimensional particle morphology and size data collected via XCT is used to provide insight regarding the discrepancies among other sizing and morphology measurement techniques. (Official contribution of the National Institute of Standards and Technology; not subject to copyright in the United States.)

We report a new hybrid integration scheme that offers for the first time a nanowire-on-lead approach, which enables independent electrical addressability, is scalable, and has superior spatial resolution in vertical nanowire arrays. The fabrication of these nanowire arrays is demonstrated to be scalable down to submicrometer site-to-site spacing and can be combined with standard integrated circuit fabrication technologies. We utilize these arrays to perform electrophysiological recordings from mouse and rat primary neurons and human induced pluripotent stem cell (hiPSC)-derived neurons, which revealed high signal-to-noise ratios and sensitivity to subthreshold postsynaptic potentials (PSPs). We measured electrical activity from rodent neurons from 8 days in vitro (DIV) to 14 DIV and from hiPSC-derived neurons at 6 weeks in vitro post culture with signal amplitudes up to 99 mV. Overall, our platform paves the way for longitudinal electrophysiological experiments on synaptic activity in human iPSC based disease models of neuronal networks, critical for understanding the mechanisms of neurological diseases and for developing drugs to treat them.

“High resolution SEM imaging” refers to the capability of discerning fine-scale spatial features of a specimen. Such features may be free-standing objects or structures embedded in a matrix. The definition of “fine-scale” depends on the application, which may involve sub-nanometer features in the most extreme cases. The most important factor determining the limit of spatial resolution is the footprint of the incident beam as it enters the specimen. Depending on the level of performance of the electron optics, the limiting beam diameter can be as small as 1 nm or even finer. However, the ultimate resolution performance is likely to be substantially poorer than the beam footprint and will be determined by one or more of several additional factors: (1) delocalization of the imaging signal, which consists of secondary electrons and/or backscattered electrons, due to the physics of the beam electron specimen interactions; (2) constraints imposed on the beam size needed to satisfy the Threshold Equation to establish the visibility for the contrast produced by the features of interest; (3) mechanical stability of the SEM; (4) mechanical stability of the specimen mounting; (5) the vacuum environment and specimen cleanliness necessary to avoid contamination of the specimen; (6) degradation of the specimen due to radiation damage; and (7) stray electromagnetic fields in the SEM environment. Recognizing these factors and minimizing or eliminating their impact is critical to achieving optimum high resolution imaging performance. Because achieving satisfactory high resolution SEM often involves operating at the performance limit of the instrument as well as the technique, the experience may vary from one specimen type to another, with different limiting factors manifesting themselves in different situations. Most importantly, because of the limitations on feature visibility imposed by the Threshold Current/Contrast Equation, for a given choice of operating conditions, there will always be a level of feature contrast below which specimen features will not be visible. Thus, there is always a possible “now you see it, now you don’t” experience lurking when we seek to operate at the limit of the SEM performance envelope.

A k-ratio is the ratio of a pair of characteristic X-ray line intensities, I, measured under similar experimental conditions for the unknown (unk) and standard (std): $$ k={I}_{unk}/{I}_{std} $$ The measured intensities can be associated with a single characteristic X-ray line (as is typically the case for wavelength spectrometers) or associated with a family of characteristic X-ray lines (as is typically the case for energy dispersive spectrometers.) The numerator of the k-ratio is typically the intensity measured from an unknown sample and the denominator is typically the intensity measured from a standard material—a material of known composition.

The incident beam energy, E0, is the parameter that determines which characteristic X-rays can be excited: the beam energy must exceed the critical excitation energy, Ec, for an atomic shell to initiate ionization and subsequent emission of characteristic X-rays. This dependence is parameterized with the “overvoltage” U0, defined as $$ {U}_0={E}_0/{E}_c $$ U0 must exceed unity for X-ray emission. The intensity, Ich, of characteristic X-ray generation follows an exponential relation: $$ {I}_{ch}={i}_Ba{\left({U}_0-1\right)}^n $$ where iB is the beam current, a and n are constants, with 1.5 ≤ n ≤ 2.

As illustrated in Fig. 16.1, the physical basis of energy dispersive X-ray spectrometry (EDS) with a semiconductor detector begins with photoelectric absorption of an X-ray photon in the active volume of the semiconductor (Si). The entire energy of the photon is transferred to a bound inner shell atomic electron, which is ejected with kinetic energy equal to the photon energy minus the shell ionization energy (binding energy), 1.838 keV for the Si K-shell and 0.098 keV for the Si L-shell. The ejected photoelectron undergoes inelastic scattering within the Si crystal. One of the consequences of the energy loss is the promotion of bound outer shell valence electrons to the conduction band of the semiconductor, leaving behind positively charged “holes” in the valence band. In the conduction band, the free electrons can move in response to a potential applied between the entrance surface electrode and the back surface electrode across the thickness of the Si crystal, while the positive holes in the conduction band drift in the opposite direction, resulting in the collection of electrons at the anode on the back surface of the EDS detector. This charge generation process requires approximately 3.6 eV per electron hole pair, so that the number of charge carriers is proportional to the original photon energy, Ep: $$ n={E}_p/3.6\ eV $$ For a Mn K-L3 photon with an energy of 5.895 keV, approximately 1638 electron–hole pairs are created, comprising a charge of 2.6 × 10−16 coulombs. Because the detector can respond to any photon energy from a threshold of approximately 50 eV to 30 keV or more, the process has been named “energy dispersive,” although in the spectrometry sense there is no actual dispersion such as occurs in a diffraction element spectrometer.

Information in SEM images about specimen properties is conveyed when contrast in the backscattered and/or secondary electron signals is created by differences in the interaction of the beam electrons between a specimen feature and its surroundings. The resulting differences in the backscattered and secondary electron signals (S) convey information about specimen properties through a variety of contrast mechanisms. Contrast (Ctr) is defined as $$ {C}_{\mathrm{tr}}=\left({S}_{\mathrm{max}}-{S}_{\mathrm{min}}\right)/{S}_{\mathrm{max}} $$ where is Smax is the larger of the signals. By this definition, 0 ≤ Ctr ≤ 1.

When studying the topographic features of a specimen, the microscopist has several useful software tools available. Qualitative stereomicroscopy provides a composite view from two images of the same area, prepared with different tilts relative to the optic axis, that gives a visual sensation of the specimen topography, as shown for a fractured galena crystal using the anaglyph method in Fig. 15.1 (software: Anaglyph Maker). The “3D Viewer” plugin in ImageJ-Fiji can take the same members of the stereo pair and render the three-dimensional surface, as shown in Fig. 15.2, which can then be rotated to “view” the surface from different orientations (Fig. 15.3).

«Trace analysis” refers to the measurement of constituents presents at low fractional levels. For SEM/EDS the following arbitrary but practical definitions have been chosen to designate various constituent classes according to these mass concentration (C) ranges

Before attempting automatic or manual peak identification, it is critical that the EDS system be properly calibrated to ensure that accurate energy values are measured for the characteristic X-ray peaks. Follow the vendor’s recommended procedure to rigorously establish the calibration. The calibration procedure typically involves measuring a known material such as copper that provides characteristic X-ray peaks at low photon energy (e.g., Cu L3-M5 at 0.928 keV) and at high photon energy (Cu K-L3 at 8.040 keV). Alternatively, a composite aluminum-copper target (e.g., a copper penny partially wrapped in aluminum foil and continuously scanned so as to excite both Al and Cu) can be used to provide the Al K-L3 (1.487 keV) as the low energy peak and Cu K-L3 for the high energy peak. After calibration, peaks occurring within this energy range (e.g., Ti K-L3 at 4.508 keV and Fe K-L3 at 6.400 keV) should be measured to confirm linearity. A well-calibrated EDS should produce measured photon energies within ±2.5 eV of the ideal value. Low photon energy peaks below 1 keV photon energy should also be measured, for example, O K (e.g., from MgO) and C K. For some EDS systems, non-linearity may be encountered in the low photon energy range. Figure 18.1 shows an EDS spectrum for CaCO3 in which the O K peak at 0.523 keV is found at the correct energy, but the C K peak at 0.282 keV shows a significant deviation below the correct energy due to non-linear response in this range.

The conventional SEM must operate with a pressure in the sample chamber below ~10−4 Pa (~10−6 torr), a condition determined by the need to satisfy four key instrumental operating conditions

Background: As part of a study into the in-service failure of the bearing surface of a large water pump, characterization was requested of the hard-facing alloy, which was observed to have separated from the stainless steel substrate, causing the failure.

Electron beams have made possible the development of the versatile, high performance electron microscopes described in the earlier chapters of this book. Techniques for the generation and application of electron beams are now well documented and understood, and a wide variety of images and data can be produced using readily available instruments. While the scanning electron microscope (SEM) is the most widely used tool for high performance imaging and microanalysis, it is not the only option and may not even always be the best instrument to choose to solve a particular problem. In this chapter we will discuss how, by replacing the beam of electrons with a beam of ions, it is possible to produce a high performance microscope which resembles an SEM in many respects and shares some of its capabilities but which also offers additional and important modes of operation.

By selecting the operating parameters of the SEM electron gun, lenses, and apertures, the microscopist controls the characteristics of the focused beam that reaches the specimen surface: energy (typically selected in the range 0.1–30 keV), diameter (0.5 nm to 1 μm or larger), beam current (1 pA to 1 μA), and convergence angle (semi-cone angle 0.001–0.05 rad). In a conventional high vacuum SEM (typically with the column and specimen chamber pressures reduced below 10−3 Pa), the residual atom density is so low that the beam electrons are statistically unlikely to encounter any atoms of the residual gas along the flight path from the electron source to the specimen, a distance of approximately 25 cm.

While amorphous substances such as glass are encountered both in natural and artificial materials, most inorganic materials are found to be crystalline on some scale, ranging from sub-nanometer to centimeter or larger. A crystal consists of a regular arrangement of atoms, the so-called «unit cell,» which is repeated in a two- or three-dimensional pattern. In the previous discussion of electron beam–specimen interactions, the crystal structure of the target was not considered as a variable in the electron range equation or in the Monte Carlo electron trajectory simulation. To a first order, the crystal structure does not have a strong effect on the electron–specimen interactions. However, through the phenomenon of channeling of charged particles through the crystal lattice, crystal orientation can cause small perturbations in the total electron backscattering coefficient that can be utilized to image crystallographic microstructure through the mechanism designated «electron channeling contrast,» also referred to as «orientation contrast» (Newbury et al. 1986). The characteristics of a crystal (e.g., interplanar angles and spacings) and its relative orientation can be determined through diffraction of the high-energy backscattered electrons (BSE) to form «electron backscatter diffraction patterns (EBSD).

Software is an essential tool for the scanning electron microscopist and X-ray microanalyst (SEMXM). In the past, software was an important optional means of augmenting the electron microscope and X-ray spectrometer, permitting powerful additional analysis and enabling new characterization methods that were not possible with bare instrumentation. Today, however, it is simply not possible to function as an SEMXM practitioner without using at least a minimal amount of software. A graphical user interface (GUI) is an integral part of how the operator controls the hardware on most modern microscopes, and in some cases it is the only interface. Even many seemingly analog controls such as focus knobs, magnification knobs, or stigmators are actually digital interfaces mounted on hand-panel controllers that connect to the microscope control computer via a USB interface.

Reading about a new subject is good but there is nothing like doing to reinforce understanding. With this in mind, the authors of this textbook have designed a number of practical exercises that reinforce the book’s subject matter. Some of these exercises can be performed with software you have available to you—either instrument vendor software or a spreadsheet like MS Excel or LibreOffice/OpenOffice Calc. Other exercises require functionality which may not be present in all instrument vendor’s software. Regardless, it is much easier to explain an exercise when everyone is working with the same tools.

The use of focused ion beams (FIB) in the field of electron microscopy for the preparation of site specific samples and for imaging has become very common. Site specific sample preparation of cross-section samples is probably the most common use of the focused ion beam tools, although there are uses for imaging with secondary electrons produced by the ion beam. These tools are generally referred to as FIB tools, but this name covers a large range of actual tools. There are single beam FIB tools which consist of the FIB column on a chamber and also the FIB/SEM platforms that include both a FIB column for sample preparation and an SEM column for observing the sample during preparation and for analyzing the sample post-preparation using all of the imaging modalities and analytical tools available on a standard SEM column. A vast majority of the FIB tools presently in use are equipped with liquid metal ion sources (LMIS) and the most common ion species used is Ga. Recent developments have produced plasma sources for high current ion beams. The gas field ion source (GFIS) is discussed in module 31 on helium ion microscopy in this book.

This module discusses the procedure used to perform a rigorous quantitative elemental microanalysis by SEM/EDS following the k-ratio/matrix correction protocol using the NIST DTSA-II software engine for bulk specimens. Bulk specimens have dimensions that are sufficiently large to contain the full range of the direct electron-excited X-ray production (typically 0.5–10 μm) as well as the range of secondary X-ray fluorescence induced by the propagation of the characteristic and continuum X-rays (typically 10–100 μm).

A conducting or semiconducting specimen must maintain good contact with electrical ground to dissipate the injected beam current. Without such an electrical path, even a highly conducting specimen such as a metal will show charging artifacts, in the extreme case acting as an electron mirror and reflecting the beam off the specimen. A typical strategy is to use an adhesive such as double-sided conducting tape to both grip the specimen to a support, for example, a stub or a planchet, as well as to make the necessary electrical path connection. Note that some adhesives may only be suitable for low magnification (scanned field dimensions greater than 100 × 100 μm, nominally less than 1,000× magnification) and intermediate magnification (scanned field dimensions between 100 μm x 100 μm, nominally less than 1,000X magnification and 10 μm × 10 μm, nominally less than 10,000× magnification) due to dimensional changes which may occur as the adhesive outgases in the SEM leading to image instability such as drift. Good practice is to adequately outgas the mounted specimen in the SEM airlock or a separate vacuum system to minimize contamination in the SEM as well as to minimize further dimensional shrinkage. Note that some adhesive media are also subject to dimensional change due to electron radiation damage during imaging, which can also lead to image drift.

The incident beam energy is one of the most useful parameters over which the microscopist has control because it determines the lateral and depth sampling of the specimen properties by the critical imaging signals. The Kanaya–Okayama electron range varies strongly with the incident beam energy: $$ {R}_{K-O}(nm)=\left(27.6\ A/{Z}^{0.89}\rho \right){E_0}^{1.67} $$ where A is the atomic weight (g/mol), Z is the atomic number, ρ is the density (g/cm3), and E0 (keV) is the incident beam energy, which is shown graphically in Fig. 11.1a–c.

There are two “zero-th level” assumptions that underpin the basis for quantitative electron-excited X-ray microanalysis

While X-ray analysis can be performed in the Variable Pressure Scanning Electron Microscope (VPSEM), it is not possible to perform uncompromised electron-excited X-ray microanalysis. The measured EDS spectrum is inevitably degraded by the effects of electron scattering with the atoms of the environmental gas in the specimen chamber before the beam reaches the specimen. The spectrum is always a composite of X-rays generated by the unscattered electrons that remain in the focused beam and strike the intended target mixed with X-rays generated by the gas-scattered electrons that land elsewhere, micrometers to millimeters from the microscopic target of interest.

The detection in SEM images of specimen features such as compositional differences, topography (shape, inclination, edges, etc.), and physical differences (crystal orientation, magnetic fields, electrical fields, etc.), depends on satisfying two criteria: (1) establishing the minimum conditions necessary to ensure that the contrast created by the beam–specimen interaction responding to differences in specimen features is statistically significant in the imaging signal (backscattered electrons [BSE], secondary electrons [SE], or a combination) compared to the inevitable random signal fluctuations (noise); and (2) applying appropriate signal processing and digital image processing to render the contrast information that exists in the signal visible to the observer viewing the final image display.

This article introduces readers to a web-based solution useful for interactive nanoscale measurements of centimeter-sized specimens. This solution is a client-server system that promotes collaborative measurements and discovery. The system consists of multiple computational modules that enable uploading microscopy images, extracting metadata, assembling many nanometer-resolution images into an image covering a centimeter-sized area, and interactive viewing and measuring of objects of interest at multiple length scales over terabyte-sized images. We illustrate the use of the system on images of aerosolized nanoparticles and dye particles on printing paper.

Nanoparticles adsorbed onto the surface of a drop can fully encapsulate the liquid, creating a robust and durable soft solid with superhydrophobic characteristics referred to as a liquid marble. Artificially created liquid marbles have been studied for about a decade but are already utilized in some hair and skin care products and have numerous other potential applications. These soft solids are usually formed in small quantity by depositing and rolling a drop of liquid on a layer of hydrophobic particles but can also be made in larger quantities in an industrial mixer. In this work, we demonstrate that microscale liquid marbles can also form through self-assembly during water condensation on a superhydrophobic surface covered with a loose layer of hydrophobic nanoparticles. Using in situ environmental scanning electron microscopy and optical microscopy, we study the dynamics of liquid marble formation and evaporation as well as their interaction with condensing water droplets. We demonstrate that the self-assembly of nanoparticle films into three-dimensional liquid marbles is driven by multiple coalescence events between partially covered droplets and is aided by surface flows causing rapid nanoparticle film redistribution. We also show that droplet and liquid marble coalescence can occur due to liquid-to-liquid contact or squeezing of the two objects into each other as a result of compressive forces from surrounding droplets and marbles. Irrelevant of the mechanism, coalescence of marbles and drops can cause their rapid movement across and rolling off the edge of the surface. We also demonstrate that the liquid marbles randomly moving across the surface can be captured and immobilized by hydrophilic surface patterns.

On behalf of the membership of the Microanalysis Society (MAS), welcome to Microscopy and Microanalysis 2011 and to Nashville. The Microscopy and Microanalysis (M&M) meeting was established in 1996 as the joint annual meeting of the Microscopy Society of America (MSA) and MAS. Since 2002, M&M has also served as the annual meeting of the International Metallographic Society (IMS), and it is now the premier meeting in our field, spanning the full spectrum of activities in microanalysis and the varied disciplines of microscopy. I thank my counterparts in the joining societies, MSA President Nestor Zaluzec and IMS President Nat Saenz, for all the time and effort they have invested to make this meeting a success.

The scanning electron microscope (SEM) and transmission electron microscope (TEM) are very powerful tools for the nanoscale characterization of materials. To make the best use of these capabilities, it is necessary to understand the nature of the uncertainties in the measurement process and to allow for errors during the data capture and data analysis stages of the measurement. By using replicated measurements and statistical analyses, it is possible to address repeatability and reproducibility issues and to assign a precision to the measurements. In general it is much harder to address accuracy issues in the experiment because some of the errors are systematic and non-statistical in nature and in almost all cases the true value of the measurand is unknown. In this work computer simulations are used to understand systematic errors and accuracy concerns in dimensional and chemical metrology in both the SEM and TEM. Multislice high-resolution TEM (HRTEM) simulations are used to produce synthetic cross-section micrographs of SiO2 gate dielectrics on Si substrates. These synthetic images are processed in the same manner as experimental images to extract a film thickness value from the micrographs. The absolute error in the dielectric film thickness measurement is determined by comparing this 'measured' value with the known, true thickness intrinsic to the sample's structural model used for simulation. This process allows the accuracy of the measurement process to be assessed for a given set of microscope, sample, environmental, data capture and data processing parameters. Statistical design of experiment methods is used to screen the influence of variables including beam tilt, along-beam thickness, dielectric thickness, defocus, astigmatism and vibration on the accuracy. The most important main effects are beam tilt, defocus and vibration, and the most significant two-term interactions are (beam tilt–defocus) and (defocus–vibration). Three-dimensional Monte Carlo simulations in Java and Jython are used to simulate the electron transport and x-ray generation of energy-dispersive x-ray (XEDS) experiments in the SEM. This code permits the simulation of synthetic hyperspectral XEDS datasets from complex, chemically heterogeneous nanostructures. Analysis of these datasets reveals the effects of self-absorption, thickness and beam broadening due to multiple scattering on the accuracy of chemical measurements in the SEM and TEM. A synthetic XEDS datacube from a four-phase Raney nickel sample is presented. While the effects mentioned above degrade both TEM-based and SEM-based measurements, their magnitude is larger and more easily visualized in the SEM because of lower beam energies and thicker samples.

A coordinated experimental and theory/simulation effort has been carried out to investigate the physics of the Traveling Wave Direct Energy Converter (TWDEC), a scheme that has been proposed in the past for the direct conversion into electricity of the kinetic energy of an ion beam generated from fusion reactions. This effort has been focused in particular on the TWDEC process in the high density beam regime, thus accounting for the ion beam expansion due to its space charge.

(La, Sr)(Al, Ta)O3 (LSAT) has recently received increased attention as a substrate for the deposition of various oxide films. LSAT is usually considered as having a perovskite structure. However, the exact atomic structure of LSAT is not yet well known and has received little attention up to now. In this paper, a series of studies using x-ray diffraction, transmission electron microscopy, and energy dispersive spectrometry were conducted to determine the atomic structure of LSAT. It was found that LSAT consists of two types of domains: one with a simple cubic perovskite structure (disordered with a = 0.3868 nm) and the other with a face-centered-cubic structure (ordered with a = 0.773 nm). Computer image simulation of high-resolution lattice images suggests that the ordered structure is due to ordering of the Al and Ta ions.

A 300 keV scanning transmission electron microscope was used to obtain tilt-series images of two two-level copper integrated circuit samples. The center-to-center layer spacing obtained from the tilt series showed internal consistency at the level of 15%, and external validity at the level of 20%.

The complex structural behavior of the perovskite-related compound Sr 4 Nb 2 O 9 [=Sr(Sr 1/3 Nb 2/3 )O 3 ] has been investigated using electron, X-ray, and neutron powder diffraction. Analysis of well-equilibrated specimens annealed at various temperatures indicated that two thermodynamically stable polymorphs occur above and below T =1250°C. The high-temperature (HT) polymorph exhibits an average cubic structure ( Fm 3 m , a ≈2 a c , where “c” refers to the cubic ∼4 A perovskite unit cell) with 1:1 (NaCl-type) partial ordering of Sr 2+ and Nb 5+ on the B sites. The low-temperature (LT) phase is monoclinic (P2 1 / n , a ≈6 a c √3, b ≈2 a c √2, c ≈ a c √6, β≈90°) with a distinct, yet unknown, B-cation arrangement. Examination of the HT cubic polymorph using electron diffraction revealed diffuse intensity contours that were consistent with additional local ordering/cluster formation on the sites of the mixed (Sr/Nb) fcc cubic sublattice of the 1:1 ordered B-cation array. Concentration of intensity observed at certain points within the diffuse intensity contours was attributed to superstructure formation. Additionally, refinements of the HT structure using X-ray and neutron powder diffraction data indicated local displacements of both oxygen and A-site Sr atoms which resembled those associated with octahedral tilting. HT-Sr 4 Nb 2 O 9 specimens subsequently annealed at 900°C exhibited remarkably complex chemical and structural behavior: The contours of diffuse intensity were replaced by arrays of sharp superlattice reflections corresponding to at least four distinct metastable superstructures, all derived from the 1:1 B-cation arrangement. Compositional analysis indicated that these phases exhibit Sr/Nb ratios slightly different from 2/1; that is, they are not true polymorphs, but rather a series of structurally distinct phases with compositions near Sr 4 Nb 2 O 9 . The metastable superstructures were attributed to ordering of Sr and Nb on the mixed B sites in the 1:1 ordered array, possibly combined with either A-site or oxygen vacancy ordering, depending on the Sr/Nb ratio in the particular phase. The metastable phases transform to the stable LT-Sr 4 Nb 2 O 9 phase upon annealing in the 1100–1200°C temperature range. The dielectric properties of HT- and LT-Sr 4 Nb 2 O 9 were measured by capacitive methods at 1 MHz. HT-Sr 4 Nb 2 O 9 exhibited an ambient permittivity of 40 with a non-monotonic temperature dependence, while that for LT-Sr 4 Nb 2 O 9 was 30 with near-linear temperature dependence. The peculiar dielectric behavior of HT-Sr 4 Nb 2 O 9 was ascribed to the competitive responses of nanodomains having somewhat different structures and compositions.

Computational models of the Traveling Wave Direct Energy Converter (TWDEC) have been developed with the goals of: demonstrating the physics of charged particle deceleration and energy conversion, supporting the development of an experimental TWDEC test article, and optimizing its performance. The first model is a 2D axisymmetric, electrostatic, particle-in-cell (PIC) simulation which includes the presence of electrodes and of an external circuit load impedance. The second model is based on an analytical approach, modeling the ion bunches as point charges, resulting in fast simulation at the expense of a reduced accuracy in the modeling of the beam dynamics. The PIC model is designed to investigate the physics of the TWDEC while the second model provides the ability to optimize electrode spacing for the maximum particle deceleration efficiency. In a concurrent experimental effort, a TWDEC test article is being developed to provide the first direct measurements of energy converted from a decelerating ion beam.

Compared to the significant body of work devoted to surface engineering for promoting dropwise condensation heat transfer of steam, much less attention has been dedicated to fluids with lower interfacial tension. A vast array of low-surface tension fluids such as hydrocarbons, cryogens, and fluorinated refrigerants are used in a number of industrial applications, and the development of passive means for increasing their condensation heat transfer coefficients has potential for significant efficiency enhancements. Here we investigate condensation behavior of a variety of liquids with surface tensions in the range of 12 to 28 mN/m on three types of omniphobic surfaces: smooth oleophobic, re-entrant superomniphobic, and lubricant-impregnated surfaces. We demonstrate that although smooth oleophobic and lubricant-impregnated surfaces can promote dropwise condensation of the majority of these fluids, re-entrant omniphobic surfaces became flooded and reverted to filmwise condensation. We also demonstrate that on the lubricant-impregnated surfaces, the choice of lubricant and underlying surface texture play a crucial role in stabilizing the lubricant and reducing pinning of the condensate. With properly engineered surfaces to promote dropwise condensation of low-surface tension fluids, we demonstrate a four to eight-fold improvement in the heat transfer coefficient., National Science Foundation (U.S.). Graduate Research Fellowship Program, National Science Foundation (U.S.) (CAREER Award 0952564), MIT Energy Initiative

We report on the structural, transport, and magnetic properties of ${\mathrm{La}}_{0.67}{\mathrm{Sr}}_{0.33}{\mathrm{MnO}}_{x}$ thin films grown in vacuum by pulsed-laser deposition. The as-grown thin films have both the matrix ${\mathrm{La}}_{1.34}{\mathrm{Sr}}_{0.66}{\mathrm{MnO}}_{4}$ phase with ${\mathrm{K}}_{2}{\mathrm{NiF}}_{4}$ structure and an embedded MnO phase. The electrical transport and magnetic properties of the films are determined mainly by those of the matrix phase. By annealing, the as-grown thin films can be transformed into the normal ${\mathrm{La}}_{0.67}{\mathrm{Sr}}_{0.33}{\mathrm{MnO}}_{3}$ single phase, which shows the expected colossal magnetoresistance effect. Based on the composition of the matrix phase, and the structural, electrical, and magnetic properties of the films, we propose that the matrix phase is possibly electron doped with a mixed valence of ${\mathrm{Mn}}^{2+}{/\mathrm{M}\mathrm{n}}^{3+}$ instead of the ${\mathrm{Mn}}^{3+}{/\mathrm{M}\mathrm{n}}^{4+}$ as in the hole-doped case.