Search

This chapter reviews structures and properties of various TiN-based heterostructures. The study of such systems is of vital importance in development of new materials with desired optimum properties. After discussion of the current literature, the recently published results of first-principles quantum molecular dynamics (QMD) calculations of heterostructures consisting of one monolayer of interfacial SiNx, SiC, BN and AlN inserted between several monolayers thick slabs of B1(NaCl)-TiN (001) and (111) in the temperature range of 0–1,400 K are reviewed in some detail. The results revealed that SiN(001) exists as pseudomorphic B1-SiN interfacial layer only at 0 K. At finite temperature, this heterostructure transforms into distorted octahedral SiN6 and tetrahedral SiN4 units aligned along the {110} directions. At 300 K, the aggregates of the SiNx units are close to a disordered, essentially amorphous SiN. After heating to 1,400 K and subsequent relaxation at 300 K, the interfacial layer corresponds to a strongly distorted Si3N4-like structure. The B1-SiN, Si3N4-like SiN and Si3N4-like Si2N3 interfaces between the TiN(111) slabs are stable in the whole temperature range considered here. The B1-Si3N4-like interfaces derived from SiN by the formation of Si-vacancies are unstable at finite temperatures. An estimate of interfacial formation energies showed that the most favorable configurations of the (111) interfaces are silicon atoms tetrahedrally coordinated to nitrogen. The most stable (001) B1-derived heterostructure with Si0.75 N interface consist of both tetrahedrally and octahedrally coordinated silicon atoms. The TiN(001)/B1-SiC/TiN(001) interface exists as pseudo-morphic B1-SiC layer between 0 K and 600 K. After heating to 900–1,400 K and subsequent static relaxations, the interfacial layer corresponds to a strongly distorted 3C-SiC-like structure oriented in the (111) direction in which the Si and C atoms are located in the same interfacial plane. The Si atoms form fourfold coordinated N-Si ≡ C3 configurations, whereas the C atoms are located in the Ti2 = C ≡ Si3 surrounding. All the (111) interfaces simulated at 0, 300, and 1,400 K have the same atomic configurations. For these interfaces, the Si and C layers correspond to the Si-C network in the (111) direction of 3C-SiC. The Si and C atoms are located in N-Si ≡ C3 and Ti3 ≡ C ≡ Si3 configurations, respectively. The BN(001) interfacial layer forms a disordered h-BN-like structure consisting of BN3 units in the whole temperature range considered here. Finally, the B1-AlN(001) interface is found to be stable within the whole temperature range.

The consept of self-energy figures prominently in the studies of interacting quantum many-particle systems. This quantity is generally defined in terms of the energy change that accompanies the addition of a single particle to the interacting system and incorporates exactly the effects of correlation and the statistics, Fermi-Dirac or Bose-Einstein. The single-particle self-energy is non-local, generally complex, and also energy-dependent. Its evaluation is ususally associated with various approximate, perturbative or decoupling, schemes as described1,2 in the literature.

Al-based alloys such as Al-Cu and Al-Mg have interesting superplastic properties which are usually assigned to sliding at grain boundaries. Even though the underlying mechanisms of superplastic behavior and grain boundary sliding are not understood at an atomic scale, some empirical relations between structural and superplastic properties are known. Among them is the observation, that grain boundary sliding and therefore superplasticity are favored in alloys with small grains that have wide grain boundary regions.

The synthesis of new phases under pressure and the study of their structure-properties relationships yield valuable information on the general laws of phase formation in multi-component alloys. They also provide an important insight in the conditions of occurrence of particular phases in systems which exhibit typical phase diagrams at ambient pressure [1]. Although pressure-temperature diagrams are available for most elements and some compounds [2], very little work has been done so far on the study of the possible competition between structural and chemical order effects in alloys as a function of pressure.

Engineering materials with desirable physical and technological properties requires understanding and predictive capability of materials behavior under varying external conditions, such as temperature and pressure. This immediately brings one face to face with the fundamental difficulty of establishing a connection between materials behavior at a microscopic level, where understanding is to be sought, and macroscopic behavior which needs to be predicted. Bridging the corresponding gap in length scales that separates the ends of this spectrum has been a goal intensely pursued by theoretical physicists, experimentalists, and metallurgists alike. Traditionally, the search for methods to bridge the length scale gap and to gain the needed predictive capability of materials properties has been conducted largely on a trial and error basis, guided by the skill of the metallurgist, large volumes of experimental data, and often ad hoc semi phenomenological models. This situation has persisted almost to this day, and it is only recently that significant changes have begun to take place. These changes have been brought about by a number of developments, some of long standing, others of more recent vintage.

Almost immediately after the discovery of quantum mechanics, physicists and metallurgist began the task of trying to understand the origins of alloy phase stability in terms of the underlying electronic structure. Hume-Rothery [1] used the chemical concepts of electronegativity and atomic size supplemented with the solid state physics notion of electron to atom ratio to formulate a set of rules that could be used to rationalize the complex array of phases that form when two elemental metals are mixed to form an alloy. The Hume-Rothery rulesand their progeny [2], [3] have proved very successful in providing metallurgists and alloy designers with a basis for understanding alloy phase stability.

Ni-Al alloys are promising candidates for new high-performance materials due to their outstanding physical properties. Moreover, this alloy system exhibits phenomena that are very interesting for theoretical study. In particular, both replacive and displacive transformations occur, and in a narrow range of composition, both processes can take place consecutively. Ni- rich alloys, with a high-temperature B2 phase, undergo a martensitic transformation in a narrow range of composition in which at low temperature, an ordered phase of Pt5 Ga3-type exists[l]. In this paper the driving force for this ordered phase will be examined in detail by using a method based on electronic structure calculations within the local density approximation.

In attempting to understand the thermodynamic properties of substitutional alloys, and specifically to construct alloy phase diagrams, it is natural to seek recourse in the vast body of knowledge acquired with respect to magnetic (spin) systems, in particular the 3D Ising model.