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Heat-treated FeCo-based magnetic alloys were characterized using a suite of electron microscopy techniques to gain insight into their structural properties. Electron channeling contrast imaging (ECCI) in the scanning electron microscope (SEM) found unique grains towards the outer edge of a FeCo sample with nonuniform background contrast. High-magnification ECCI imaging of these nonuniform grains revealed a weblike network of defects that were not observed in standard uniform background contrast grains. High-resolution electron backscattered diffraction (HR-EBSD) confirmed these defect structures to be dislocation networks and additionally found subgrain boundaries within the nonuniform contrast grains. The defect content within these grains suggests that they are unrecrystallized grains, and ECCI can be used as a rapid method to quantify unrecrystallized grains. To demonstrate the insight that can be garnered via ECCI on these unique grains, the sample was imaged before and after micro indentation. This experiment showed that slip bands propagate throughout the material until interacting with the dislocation networks, suggesting that these specific defects provide a barrier to plastic deformation. Taken together, these results show how ECCI can be used to better understand failure mechanisms in alloys and provides further evidence that dislocation networks play a critical role in the brittle failure of FeCo alloys.

Indium (In) and other low melting point metals are used as interconnects in a variety of hybridized circuits and a full understanding of the metallurgy of these interconnects is important to the reliability and performance of the devices. This paper shows that room temperature focused ion beam (FIB) preparation of cross-sections, using Ga+ or Xe+ can result in artifacts that obscure the true In microbump structure. The use of modified milling strategies to minimize the increased local sample temperature are shown to produce cross-sections that are representative of the In bump microstructure in some sample configurations. Furthermore, cooling of the sample to cryogenic temperatures is shown to reliably eliminate artifacts in FIB prepared cross-sections of In bumps allowing the true bump microstructure to be observed.

The transformation of unstable austenite to ferrite or α′ martensite as a result of exposure to Xe+ or Ga+ ions at room temperature was studied in a 304 stainless steel casting alloy. Controlled Xe+ and Ga+ ion beam exposures of the 304 were carried out at a variety of beam/sample geometries. It was found that both Ga+ and Xe+ ion irradiation resulted in the transformation of the austenite to either ferrite or α′ martensite. In this paper, we will refer to the transformation product as a BCC phase. The crystallographic orientation of the transformed area was controlled by the orientation of the austenite grain and was consistent with either the Nishiyama–Wasserman or the Kurdjumov–Sachs orientation relationships. On the basis of the Xe+ and Ga+ ion beam exposures, the transformation is not controlled by the chemical stabilization of the BCC phase by the ion species, but is a result of the disorder caused by the ion-induced recoil motion and subsequent return of the disordered region to a more energetically favorable phase.

The surfaces of textured polycrystalline N-type bismuth telluride and P-type antimony telluride materials were investigated using ex situ photoelectron emission microscopy (PEEM). PEEM enabled imaging of the work function for different oxidation times due to exposure to air across sample surfaces. The spatially averaged work function was also tracked as a function of air exposure time. N-type bismuth telluride showed an increase in the work function around grain boundaries relative to grain interiors during the early stages of air exposure-driven oxidation. At longer time exposure to air, the surface became homogenous after a ∼5 nm-thick oxide formed. X-ray photoemission spectroscopy was used to correlate changes in PEEM imaging in real space and work function evolution to the progressive growth of an oxide layer. The observed work function contrast is consistent with the pinning of electronic surface states due to the defects at a grain boundary.

Active slip systems in body centered cubic (BCC) metals remain ambiguous and controversial. In this work, active slip planes in tantalum are investigated using single crystal experiments and simulations. Four tantalum single crystals with [001], [011], [ 1 ¯ 11 ] and [ 1 ¯ 49 ] orientations along the loading direction are analyzed after quasi-static compression tests and dynamic Taylor impact tests. Mechanical behavior, deformed shape, texture evolution and slip traces of single crystals are analyzed and compared with crystal plasticity finite element simulations to investigate active slip systems in tantalum. Both experimental observations and modeling results support dominant dislocation slip along {112} planes in tantalum under compressive load.

Intermetallic alloys possess exceptional soft magnetic properties, including high permeability, low coercivity, and high saturation induction, but exhibit poor mechanical properties that make them impractical to bulk process and use at ideal compositions. We used laser-based Additive Manufacturing to process traditionally brittle Fe–Co and Fe–Si alloys in bulk form without macroscopic defects and at near-ideal compositions for electromagnetic applications. The binary Fe–50Co, as a model material, demonstrated simultaneous high strength (600–700 MPa) and high ductility (35%) in tension, corresponding to a ∼300% increase in strength and an order-of-magnitude improvement in ductility relative to conventionally processed material. Atomic-scale toughening and strengthening mechanisms, based on engineered multiscale microstructures, are proposed to explain the unusual combination of mechanical properties. This work presents an instance in which metal Additive Manufacturing processes are enabling, rather than limiting, the development of higher-performance alloys.

Control of the atomic structure, as measured by the extent of the embrittling B2 chemically ordered phase, is demonstrated in intermetallic alloys through additive manufacturing (AM) and characterized using high fidelity neutron diffraction. As a layer-by-layer rapid solidification process, AM was employed to suppress the extent of chemically ordered B2 phases in a soft ferromagnetic Fe-Co alloy, as a model material system of interest to electromagnetic applications. The extent of atomic ordering was found to be insensitive to the spatial location within specimens and suggests that the thermal conditions within only a few AM layers were most influential in controlling the microstructure, in agreement with the predictions from a thermal model for welding. Analysis of process parameter effects on ordering found that suppression of B2 phase was the result of an increased average cooling rate during processing. AM processing parameters, namely interlayer interval time and build velocity, were used to systematically control the relative fraction of ordered B2 phase in specimens from 0.49 to 0.72. Hardness of AM specimens was more than 150% higher than conventionally processed bulk material. Implications for tailoring microstructures of intermetallic alloys are discussed.

Structural disorder causes materials’ surface electronic properties, e.g., work function ([Formula: see text]), to vary spatially, yet it is challenging to prove exact causal relationships to underlying ensemble disorder, e.g., roughness or granularity. For polycrystalline Pt, nanoscale resolution photoemission threshold mapping reveals a spatially varying [Formula: see text] eV over a distribution of (111) vicinal grain surfaces prepared by sputter deposition and annealing. With regard to field emission and related phenomena, e.g., vacuum arc initiation, a salient feature of the [Formula: see text] distribution is that it is skewed with a long tail to values down to 5.4 eV, i.e., far below the mean, which is exponentially impactful to field emission via the Fowler–Nordheim relation. We show that the [Formula: see text] spatial variation and distribution can be explained by ensemble variations of granular tilts and surface slopes via a Smoluchowski smoothing model wherein local [Formula: see text] variations result from spatially varying densities of electric dipole moments, intrinsic to atomic steps, that locally modify [Formula: see text]. Atomic step-terrace structure is confirmed with scanning tunneling microscopy (STM) at several locations on our surfaces, and prior works showed STM evidence for atomic step dipoles at various metal surfaces. From our model, we find an atomic step edge dipole [Formula: see text] D/edge atom, which is comparable to values reported in studies that utilized other methods and materials. Our results elucidate a connection between macroscopic [Formula: see text] and the nanostructure that may contribute to the spread of reported [Formula: see text] for Pt and other surfaces and may be useful toward more complete descriptions of polycrystalline metals in the models of field emission and other related vacuum electronics phenomena, e.g., arc initiation.

Multiple experimental configurations for performing nanoscale orientation mapping are compared to determine their fidelity to the true microstructure of a sample. Transmission Kikuchi diffraction (TKD) experiments in a scanning electron microscope (SEM) and nanobeam diffraction (NBD) experiments in a transmission electron microscope (TEM) were performed on thin electrodeposited hard Au films with two different microstructures. The Au samples either had a grain size that is >50 or

In Part I, equal channel angular extrusion (ECAE) was demonstrated as a novel, simple-shear deformation process for producing bulk forms of the low ductility Fe–Co–2V (Hiperco 50A®) soft ferromagnetic alloy with refined grain sizes. Microstructures and mechanical properties were discussed. In this Part II contribution, the crystallographic textures and quasi-static magnetic properties of ECAE-processed Hiperco were characterized. The textures were of a simple-shear character defined by partial {110} and 〈111〉 fibers inclined relative to the extrusion direction, in agreement with the expectations for simple-shear deformation textures of BCC metals. These textures were observed throughout all processing conditions and only slightly reduced in intensity by subsequent recrystallization heat treatments. Characterization of the magnetic properties revealed a lower coercivity and higher permeability for ECAE-processed Hiperco specimens relative to the conventionally processed and annealed Hiperco bar. The effects of the resultant microstructure and texture on the coercivity and permeability magnetic properties are discussed.

Processing of the low workability Fe-Co-1.5V (Hiperco® equivalent) alloy is demonstrated using the Laser Engineered Net Shaping (LENS) metals additive manufacturing technique. As an innovative and highly localized solidification process, LENS is shown to overcome workability issues that arise during conventional thermomechanical processing, enabling the production of bulk, near net-shape forms of the Fe-Co alloy. Bulk LENS structures appeared to be ductile with no significant macroscopic defects. Atomic ordering was evaluated and significantly reduced in as-built LENS specimens relative to an annealed condition, tailorable through selection of processing parameters. Fine equiaxed grain structures were observed in as-built specimens following solidification, which then evolved toward a highly heterogeneous bimodal grain structure after annealing. The microstructure evolution in Fe-Co is discussed in the context of classical solidification theory and selective grain boundary pinning processes. Magnetic properties were also assessed and shown to fall within the extremes of conventionally processed Hiperco® alloys. Hiperco® is a registered trademark of Carpenter Technologies, Readings, PA.

Active slip systems in body centered cubic (BCC) metals remain ambiguous and controversial. In this work, active slip planes in tantalum are investigated using single crystal experiments and simulations. Four tantalum single crystals with [100], [110], [111] and [149] orientations along the loading direction are analyzed after quasi-static compression tests and dynamic Taylor impact tests. Mechanical behaviors, deformed shapes, texture evolutions and slip traces of single crystals are analyzed and compared with crystal plasticity finite element simulations to investigate active slip systems in tantalum. Both experimental observations and modeling results support dominant dislocation slip along {112} planes in tantalum.

Ferroelastic domain walls provide opportunities for deterministically controlling mechanical, optical, electrical, and thermal energy. Domain wall characterization in micro- and nanoscale systems, where their spacing may be of the order of 100 nm or less is presently limited to only a few techniques, such as piezoresponse force microscopy and transmission electron microscopy. These respective techniques cannot, however, independently characterize domain polarization orientation and domain wall motion in technologically relevant capacitor structures or in a non-destructive manner, thus presenting a limitation of their utility. In this work, we show how backscatter scanning electron microscopy utilizing channeling contrast yield can image the ferroelastic domain structure of ferroelectric films with domain wall spacing as narrow as 10 nm. Combined with electron backscatter diffraction to identify grain orientations, this technique provides information on domain orientation and domain wall type that cannot be readily measured using conventional non-destructive methods. In addition to grain orientation identification, this technique enables dynamic domain structure changes to be observed in functioning capacitors utilizing electrodes that are transparent to the high-energy backscattered electrons. This non-destructive, high-resolution domain imaging technique is applicable to a wide variety of ferroelectric thin films and a multitude of material systems where nanometer-scale crystallographic twin characterization is required.

The objective of this study was to evaluate effects of low temperature on hydrogen-assisted crack propagation in forged 304L austenitic stainless steel. Fracture initiation toughness and crack-growth resistance curves were measured using fracture mechanics specimens that were thermally precharged with 140 wppm hydrogen and tested at 293 K or 223 K (20 °C or −50 °C). Fracture initiation toughness for hydrogen-precharged forgings decreased by at least 50 to 80 pct relative to non-charged forgings. With hydrogen, low-temperature fracture initiation toughness decreased by 35 to 50 pct relative to room-temperature toughness. Crack growth without hydrogen at both temperatures was microstructure-independent and indistinguishable from blunting, while with hydrogen microcracks formed by growth and coalescence of microvoids. Initiation of microvoids in the presence of hydrogen occurred where localized deformation bands intersected grain boundaries and other deformation bands. Low temperature additionally promoted fracture initiation at annealing twin boundaries in the presence of hydrogen, which competed with deformation band intersections and grain boundaries as sites of microvoid formation and fracture initiation. A common ingredient for fracture initiation was stress concentration that arose from the intersection of deformation bands with these microstructural obstacles. The localized deformation responsible for producing stress concentrations at obstacles was intensified by low temperature and hydrogen. Crack orientation and forging strength were found to have a minor effect on fracture initiation toughness of hydrogen-supersaturated 304L forgings.

The article provides information about an upcoming conference from the program chair. The Microscopy Society of America (MSA), the Microanalysis Society (MAS), and the International Metallographic Society (IMS) invite participation in Microscopy & Microanalysis 2016 in Columbus, Ohio, July 24 through July 28, 2016.

This paper describes the evolution of friction-induced microstructures underneath the wear surfaces of nickel single crystals, and their role in influencing the steady state coefficient of friction. Using a custom built tribometer, friction measurements were made on {001} and {011} crystal faces along several crystallographic directions in unidirectional linear mode with a Si3N4 ball. Cross-sections of wear scars were prepared by focused ion beam (FIB) microscopy, and electron backscattered diffraction (EBSD) and transmission electron microscopy (TEM) were used to analyze the evolution of crystallographic texture and recrystallization. The formation of a thin (10–15 nm) ultra-nanocrystalline layer in the {011} orientation resulted in significant reductions in the coefficient of friction for that configuration, from the initial maximum of 0.64 to a steady state value of 0.32, after approximately 500 cycles of linear, unidirectional sliding. Crystal plasticity theory was used to describe the evolution of dislocation content in the worn material, which in turn provides an estimate of the characteristic grain size as a function of the imposed strain. An analysis of grain boundary sliding in ultra-fine-grained material suggests a mechanism for friction reduction. Model predictions of the contribution of grain boundary sliding (relative to plastic deformation) to friction provide a phenomenological description of the lubrication due to ultra-nanocrystalline surface microstructures.

The high-cycle fatigue life of nanocrystalline and ultrafine-grained Ni-Fe was examined for five distinct grain sizes ranging from approximately 50–600 nm. The fatigue properties were strongly dependent on grain size, with the endurance limit changing by a factor of 4 over this narrow range of grain size. The dataset suggests a breakdown in fatigue improvement for the smallest grain sizes

Vacuum assisted ex situ lift out may be used for fast, easy, and reproducible plan view specimen preparation. Manipulation of samples via beveled hollow glass probes whose plane of interest is parallel to slotted grids allow for conventional FIB milling for S/TEM analysis.

Wear surface analysis , specifically in the subsurface regions, is essential for understanding the fundamental mechanisms of friction and material removal during wear. The advent of focused ion-beam systems (combined scanning electron and focused ion-beam microscopes) has made it possible to prepare site-specific cross sections of wear surfaces for scanning electron microscopy, electron backscattered diffraction, transmission electron microscopy, and X-ray microanalysis. During the past decade, these state-of-the-art electron microscopy techniques are being increasingly used to characterize wear-induced changes to microstructures and crystallographic textures.

“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.