This chapter describes the use of neutrons and X-rays as probes in the study of structural and dynamic properties of metallic materials. Crystalline materials are characterized by their diffraction peaks related to their average crystallographic structure. In real crystals, locally displaced atoms and chemically (or isotopically for neutrons) different species may lead not only to changes of peak shapes and positions, but also to additional (diffuse) scattering between Bragg peak, including scattering around the primary beam (small-angle scattering). All these features can be used to extract information about the state of a sample, its compositional and structural variations on a scale depending on the scattering, in static and time-resolved kinetic studies. Energy-resolved scattering also offers an insight into solid-state dynamics on a microscopic scale. Some of the most important methods will be described and illustrated by instructive examples. The presentation offers a combined view of neutron and X-ray scattering, with the necessary simplifications dictated by space limitations. The special properties of thermal neutrons and of hard X-rays (now widely available at synchrotron radiation sources), their mutual combination, and combinations with other methods, in particular electron microscopy, offer ample opportunity to better understand and control materials properties. After a brief introduction to scattering from real crystals and some general ideas about long-range strains and Bragg peaks, the vicinity of Bragg peaks (displacement scattering at large scattering angles), the scattering far away from Bragg peaks (chemical heterogeneities, short-range order), and, in greater detail, small-angle scattering (which is not sensitive to the extent of crystallinity, but to nanoscale variations of chemical composition and of magnetization, precipitation) will be described, along with classical and more recent applications related to short-range ordering and precipitation in bulk and nanostructured alloys. Some other fields are only briefly addressed (grazing-incidence studies of surfaces, radiography, absorption spectroscopies, coherent X-rays). The final section offers some information on the influence of defects on lattice dynamics and on (slow) diffusive motion in materials.

Attempts to understand the stabilization mechanism of structurally complex metallic alloy (CMA) phases dates back to 1936 when Mott and Jones interpreted the stabilization of Cu5Zn8 and Al4Cu9 gamma-brasses containing 52 atoms per unit cell at particular electrons per atom ratio e/a in terms of the contact of the free electron sphere with the Brillouin zone. The Mott and Jones theory has been thought to serve as a milestone in establishing the Hume-Rothery electron concentration rule empirically established in the late 1920s. However, we have soon realized that it can be hardly applied when alloys contain transition metal (TM) elements as major constituent elements. The exploration of the phase stabilization mechanism for TM-bearing alloys has remained unsettled in physical metallurgy. In particular, the discovery of the Al-Mn quasicrystal by Shechtman et al. in 1984 and that of thermodynamically stable quasicrystals by Tsai et al. at more or less constant e/a values have intensified the interest in the Hume-Rothery electron concentration rule for TM-bearing compounds. Among numerous topics in the field of the electron theory of metals, the present authors focused on the e/a-dependent phase stabilization mechanism of CMAs, in particular, those containing a large amount of TM elements by making full use of the FLAPW formalism, in which wave functions outside the muffin-tin (MT) spheres are expanded into plane waves over allowed reciprocal lattice vectors. They have established the new Hume-Rothery electron concentration rule for CMAs in possession of a pseudogap at the Fermi level, regardless of whether TM elements are involved or not. The method to determine reliably the e/a value for TM elements has been also proposed. The relevant data for 3d-, 4d- and 5d-TM elements are presented.

This Chapter describes the theoretical foundations of most of the major phases that occur due to the solid-state transformation of austenite. The phases include allotriomorphic ferrite, pearlite, Widmanstätten ferrite, bainite and martensite, covering in each case the thermodynamics, kinetics and evolution of morphological features, and where appropriate, the crystallography and shape deformation of the phase concerned. Elementary mathematical derivations are included without compromising on the seminal features of particular transformations. Case studies are then presented, which show the ability of the theory to predict useful steels, for example, the TRIP and TWIP alloys, bulk nanostructured bainite, and welding materials that mitigate the development of residual stress in constrained assemblies

Hydrogen in metals attracts interest from scientists since many decades. Most of the interesting properties are related to the small size of hydrogen: its interstitial diffusion accompanied by quantum mechanical tunnel transport results in an extraordinary high mobility of hydrogen atoms in materials. For metals, H diffusivity may reach values as known for ions in aqueous solutions. Thus, thermodynamic equilibrium is reached within comparably short times even at room temperature. Therefore, metal–hydrogen systems are often used as model systems to study physical or chemical properties and their change with concentration (see, for example Oates and Flanagan, 1981, 1981a or Pundt and Kirchheim (2006)). In 1937, Lacher (1937) already used Pd–H (Flanagan and Oates, 1981, 1991) to study solute–solute interactions and interpreted it in the framework of a quasi-chemical approach (Lacher, 1937). The quantum mechanical tunneling as a diffusion mechanism also for atoms in solids was first discovered and discussed for hydrogen tunneling in metals (Flynn and Stoneham, 1970; Völk and Alefeld, 1975; Birnbaum and Flynn, 1976). Völk and Alefeld (1978), Zabel and Peisl (1979, 1980), and Steyrer and Peisl (1986) studied hydrogen density modulations that are related to the sample geometry; and Zabel and his colleagues, as published by Miceli et al. (1985), Uher et al. (1987), Song et al. (1996, 2000), and Uher et al. (1987), firstly used metal–hydrogen systems to study the behavior of systems with reduced dimensions and modulated hydrogen affinity. Kirchheim (1988) and colleagues extensively studied metal–hydrogen systems as representative for solute/solvent systems. The high mobility of hydrogen further allows studying the impact of defects that usually annihilate at elevated temperatures, see Gottstein (2001). It was, therefore, suggested to use hydrogen as a probe for defects (Cahn, 1990; Flanagan et al., 2001a, 2001b; Kirchheim, 2004) and perform site energy spectroscopy by gradually increasing the hydrogen chemical potential.

This chapter is an overview on computational material science methods applied to understanding the process–microstructure–property relationships of metallic material systems, or computational metallurgy. The emphasis will be on modeling and predicting the stability (thermodynamics) and the temporal/spatial evolution (kinetics) of microstructures, or computational microstructural science using a combination of electronic/atomistic level first-principles calculations of structures and thermodynamic properties of individual structural phases in a microstructure and interfaces and mesoscale microstructure evolution models.

Recovery, recrystallization and grain growth are among the most important metallurgical heat treatment processes to soften cold worked metals and design desired microstructures and textures. Specifically the reduction in grain size can be efficiently achieved by recrystallization. While plastic cold working increases the stored energy of metals, mainly through dislocation accumulation, recovery and specifically recrystallization lead to it reduction. While recovery describes the gradual re-ordering and annihilation of the stored dislocations, primary recrystallization proceeds discontinuously by the formation and motion of high angle grain boundaries which discontinuously sweep the deformation substructure. Grain growth describes the process of competitive capillary driven coarsening of the average grain size. This chapter reviews the main mechanisms, lattice defects, and driving forces associated with recovery, recrystallization and grain growth and provides an introduction to the simulation of these phenomena.

More than fifty years have passed since Müller and his coworkers invented the atom-probe field ion microscope (APFIM) in 1968 (Müller et al., 1968). The APFIM was originally developed as a tool for surface science; however, from the emerging stage of the technique, physical metallurgists realized that its use could solve many critical problems on the microstructures of metallic materials, in particular steels. By the 1980s, several researchers started applying the atom probe (AP) technique for microstructural characterizations of various metallic materials. However, the recent successful implementation of pulsed laser to assist field evaporation for 3DAP analyses expanded the application areas to a wide variety of materials including semiconductors, insulator thin films, and even to bulk insulating ceramics. In addition, the development of site specific specimen preparation method using the focused ion beam (FIB) technique made it possible to prepare tips from all types of materials, including powder, thin films, and devices. The aim of this chapter is to provide the basic principles of the 3DAP technique for those who are about to start using this technique for microstructure characterization of materials including metals and alloys, ceramics and semiconductors.

Since publication of the 4th edition of Physical Metallurgy, there have been many breakthroughs in the characterization and analysis of microstructure, and probably none more important than the rapid evolution of three-dimensional (3D) materials analyses. This chapter provides an overview of microstructure in metallic materials that captures many of these new developments, with emphasis on microstructures developed during solid-state evolution. Including all microstructural forms and elements of microstructure is well beyond the scope of a single chapter; the focus here is thus on some representative examples centered about the characterization, description, and classification of 3D solid-state metallic microstructures.

This article is intended to acquaint the reader with the basic techniques that are currently available for acquiring maps of materials microstructure based on crystallographic orientation. Examples are given of the applications of the techniques with the aim of stimulating the imagination of the reader. Many readers will be familiar with the electron back scattered diffraction (EBSD) method because of its general availability but there is a wide variety of available methods, each of which is best suited to certain problems. All methods rely on generating diffraction patterns via the interaction of electrons or X-rays with (polycrystalline) matter. In some cases, especially electrons, patterns from individual crystallites are required whereas in others, especially with X-rays, overlapping patterns must be disaggregated. A working knowledge of the methods used to describe orientations is assumed. Likewise, a knowledge of the physics that governs the interaction of electrons or X-rays with matter is assumed and the origins of the diffraction patterns used in the methods described are not explained. Neutron diffraction is not included here because the method cannot afford sufficient spatial resolution to be useful for mapping. For the most part, the methods produce maps with the orientation specified at points on a regular grid. In general, the electron diffraction-based methods provide a direct measurement at each grid point whereas in the X-ray methods, the gridded map is derived from other diffraction data.

Transport in materials via random individual atomic migration steps (jumps) is called ‘diffusion’. General understanding and detailed theory of diffusional transport represent a classical field of material science. While in liquids and gases, diffusion rate ranges up to millimeters or even centimeters per second, transport in solid materials is rather slow. In densely packed metals near to the melting point, one can expect about a micrometer per second and this rate drops down to about a nanometer per second at half the melting temperature. At room temperature atomic migration is practically frozen, except for few exceptional cases of small interstitial impurities. Measurement of diffusion in solids needs therefore sensitive techniques and well-developed microscopy. Not surprising that a scientific proof of diffusion in solid metals came rather late in history of science (Roberts-Austen, 1896).