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We estimate the free energy of ionic surfaces using an Einstein model for the phonons in the vicinity of the surface. The free energy of three of the low energy surfaces ((00l), (211) and (110)) of stannic oxide is calculated.

Biomaterials are hierarchical systems whose structure and properties represent a major challenge to simulation. We briefly discuss the remarkable experimental data now available both on the structure, formation and properties of biominerals. In the light of this we discuss current attempts to simulate biomaterials at atomic and meso length scales and timescales and the range of physical effects that are important in producing biomaterials. We emphasise the importance of obtaining robust forcefields for these systems and the sensitivity of simulations to the forcefields used. We conclude by suggesting future directions for the field and remaining problems to be solved.

Self-assembled monolayers of long-chain carboxylic acids are often used as substrates to promote the growth of oriented crystals. Recent work has shown that the length of the chain (odd or even number of carbon atoms) determines whether oriented growth is observed. We use molecular dynamics simulations to investigate whether the configuration of the headgroups is significantly different in the two cases. We conclude that there are differences between odd- and even-length chains, even at 300 K and in the presence of water for some packings of the monolayer. We discuss whether these differences are large enough to account for the different behavior.

We have modelled the surface diffusion and growth of BaO and SrO both in the homoepitaxial and heteroepitaxial BaO on SrO and SrO on BaO cases. The diffusion proceeds most favourably by an exchange mechanism involving the surface layer. When impurities are adsorbed on the surface this can lead to intermixing between the layers. This strongly suggests that ionic materials may not be grown on a substrate with a similar structure without significant intermixing. Island growth begins with the formation of individual clusters which grow and merge together.

Calcite crystals nucleate on the (01.2) face on a diverse range of organic substrates, including self-assembled monolayers, hydrogen-bonded ribbons, and polymer rafts. The (01.2) face of calcite is a polar surface. Therefore macroscopic crystal growth can only occur if the dipole moment is quenched. We demonstrate that the dipole moment can be quenched for a given polar direction by adsorption onto an organic substrate with arbitrary charge density. The density of ions in the outer calcium plane must be modified, by introducing rows of vacancies, to fulfill the condition of zero net dipole moment. Interfacial energies are calculated for interfaces between the polar (01.2) and (00.1) faces of calcite and stearic acid monolayers with a range of densities. It was found that, contrary to the experimental evidence, the (00.1) face has lower interfacial energy than the (01.2) face with monolayers with equivalent densities. We give an explanation for this discrepancy based on kinetic models.

Living organisms can control the size, shape, and structure of minerals. Attempts to reproduce this biological control in the laboratory often use Langmuir monolayers of long-chain carboxylic acids. We use large-scale molecular dynamics simulations to calculate the interfacial energies of calcite crystals grown on stearic (octadecanoic) acid monolayers. In light of these simulations we discuss the argument that the orientation of the growing mineral is controlled by the organic substrate acting as a template which the mineral must fit in order to grow.

Nucleation and growth models are well developed for nucleation on homogeneous substrates, and they can typically be described in terms of three energy parameters. Nucleation on substrates containing point–defect traps has been investigated, at the cost of introducing more energy parameters. This paper outlines the quantitative description of such growth models, using rate and rate–diffusion equations, in terms of energies for individual surface processes, with examples taken from metal–metal, metal–insulator and semiconductor growth. The challenge to modelling is to describe the large range of length and time–scales in thin–film fabrication and degradation, without relying on too many (unknown) material parameters, which often occur in combination. Separating them into elementary processes often proves to be a challenge. One typically requires selective nucleation using patterned substrates, in combination with controlled, self–organized, growth for reliable nanotechnology. Reconstructed semiconductor surfaces offer both a further challenge to modelling and an opportunity for future technology; these paradoxes are discussed briefly.

The atomic and electronic structures of three LixV2O5 phases have been investigated using plane wave calculations based on density functional theory. We show that reduction of the vanadium ions occurs during lithium intercalation. At low lithium concentrations this reduction is limited to specific ions. Changes to unit cell dimensions can be predicted correctly. Calculated cell potentials are in good agreement with experiment. Initial “atomistic” calculations were successfully used to find the most stable sites for lithium in the V2O5 lattice. Calculated activation energies suggest that lithium is highly mobile along the 010 direction.

We present a new method for calculating the perturbed phonon spectrum of defect crystals and show how this can be used to obtain entropies. The method considers a periodic structure in which a cluster of ions around the defect, placed in their relaxed positions, is embedded in a non-primitive unit cell. The normal modes of this structure may be used to calculate local modes and, since a reasonable representation of k space may be obtained by projection, considering only the Γ point of this cell, we may readily obtain values for the partition function and hence for thermodynamic quantities.

A theoretical study has been made of the defect structure of calcium sulphide. Empirical interatomic potentials are obtained from the analysis of crystal data for calcium and sodium sulphides. Defect energies are calculated using the Harwell HADES program which is based on a generalized Mott - Littleton method. The results predict the dominance of Schottky disorder in the crystal. The activation energy for the migration of both anions and cations is calculated. The results are in good agreement with the experimental data, suggesting that vacancy migration predominates.The results show that the ionic model works for sulphides, and its application to other chalcogenide systems is discussed.

A near-coincidence-site lattice theory is employed in conjunction with atomistic simulation to elucidate the interfacial stability and detailed interfacial structural configurations for BaO on MgO and CeO2 on Al2O3. The stability of these interfaces was found to be strongly dependent not only on the mismatch and the misorientational angles of the two incommensurate lattice structures but also on the initial orientations of the ions and their relaxations at the interfaces.

Ag-MgO and Au-MgO interfaces in the (100) and (110) orientations have been modelled using the MIDAS code. Image interactions between the ions and the metal and short-range interactions between the metal atoms and ion cores are included. The short-range interactions are calculated using accurate Hartree-Fock methods. Classical image interactions are modified by suppressing fluctuations of the induced charge density which have wavelengths longer than the Fermi wave-vector. This approach eliminates the singularities which plague classical image models. The answers agree well with calculations performed using the local-density approximation (including, unfortunately, the poor prediction of the wetting angle). The predicted interfacial bonding energy is 0.8 eV (1·81 Jm-2).

In the cooldown phase of the deposition process, plasma-sprayed ceramic barrier coatings generate large residual stresses. In this paper, a scheme for the analysis of the thermal stresses is given, which accounts for the nonuniformity of the temperature field, and for the microstructure of the coating. The considered microstructural data include the morphology and orientation pattern of the porosities. The case of a coated sphere is studied, as an application. In this connection, the coating structure is computer-simulated, and the temperature profile is numerically determined.

We calculate the excitation energies of the Vk centre in NaCl by an approach combining an Unrestricted Hartree Fock code with classical potentials to simulate the defect and the distorted lattice around it. This treatment gives better results than previous calculations based on the levels of the Cl-2 molecule. Results on the elastic dipole tensor are also discussed.

Methods are available which calculate the free energies of defects in ionic systems within the quasi-harmonic approximation. The Harwell SHEOL code is used to calculate the defect enthaplies and entropies for a number of alkali halides where reliable diffusion data are known. We discuss the trends in the calculations and the comparison with experiment.

The second-electron affinity of the oxygen atom is often required in modelling the behaviour of ceramics. It is usually taken to be about 8 eV, independent of the structure of the oxide being modelled. We show that this is not so.Quantum-mechanical calculations show that the true second-electron affinity of an in-crystal O2 - ion depends on both the chemical composition and the nuclear geometry. Moreover, for a binary oxide at its equilibrium geometry, this true second-electron affinity is about 40% greater than 8 eV. This discrepancy between the true second-electron affinity and the value of 8 eV derived from semi-empirical potentials determined by fitting to experimental crystal data arises from the structural dependence of the true value. In a semi-empirical treatment, this dependence becomes absorbed into the quantity regarded as the short-range cation-oxide interaction with the consequence that the resulting second-electron affinity is smaller than the true value. A further consequence is that these short-range cation-oxide interactions are dependent on the crystal structure. Thus, for example, a semi-empirical potential determined from a binary oxide with the rock salt structure cannot be transferred to a polymorph with the eightfold-coordinated CsCl structure.

Defect parameters are calculated for a wide range of defects in UO2. In addition to internal energies, entropies of formation have been calculated for some defects. Enthalpies are obtained for comparison with experimental data, where available. Defect processes studied include intrinsic defects, cation substitution, ion migration and defect clustering.

Computer simulation techniques have been used to calculate the energy and entropy changes, at constant volume, for the formation of a Frenkel defect, for the association of a cation vacancy with a cation impurity, and for the frequency of the exchange of a cation impurity with a nearest-neighbour cation vacancy in both silver chloride and silver bromide. The results are used to compute the corresponding Helmholtz energy changes at constant volume, which may (to a good approximation) be identified with the corresponding Gibbs energies at constant pressure, as derived from experimental measurements of ionic transport. The impurities considered were Cd2+, Na+ and Cs+ in AgCl and Cd2+ Na+, Rb+ and Cs+ in AgBr. The results were also used to calculate impurity-diffusion coefficients for these systems. While the results of these calculations possess many encouraging features, it is clear that the difficult problems involved in modelling the silver halides have not yet been solved entirely satisfactorily.

Calculations of the electronic structure of a pair of O-2 ions in a point-ion lattice are analysed in the manner of a shell model. Induced dipole moments on the highly polarizable ions are explicitly accounted for in fitting the repulsive potential to a Born form. The resulting potential, together with an empirical Mg2+-O2- potential, is used to calculate bulk and defect properties of MgO. The values obtained for bulk dielectric properties using this method are worse than those found from empirical potentials, but defect properties are very similar.

Anion and cation vacancy formation energies are calculated at the near interface region of the BaO(100)/MgO(100) interface. Image theory has been shown to correctly calculate the behaviour of vacancies approaching an MgO(100) surface. However, we show it fails to predict the behaviour of vacancies near the BaO(100)/MgO(100) interface. This is attributed to the misfit between the two materials.

In a previous paper it was shown that, for defects in transition metal oxides, there was evidence for an additional term in the potential energy beyond the usual monopole and dipole terms of the shell model. This appeared to be a quadrupole term in which the open 3d shells of the host ions, as well as any impurities, responded to local electric field gradients. Quantitative estimates are made which confirm our hypothesis and, in CoO, show that the effect is sufficient to bring theory into accord with experimental measurements. Such quadrupole interactions will also be relevant to other systems with host ions or impurities having incomplete electric shells.

We use atomistic modelling methods to predict correlation factors and motion energies for the diffusion of impurity ions in CoO and NiO. These use the shell model, and include the effects of impurity-vacancy binding, the temperature dependence of the extent of non-stoichiometry, and the explicit calculation of the ratios of critical jump frequencies. We further test the validity of some approximations to jump frequencies. Despite the success of the correlation factor calculations, the absolute motion energies do not agree well with experiment. We have identified the reason for this, associated with the open-shell transition metal ions; quantitative estimates are given in a subsequent paper.

Most molecular simulations assume a potential model. In some cases, where one is interested only in qualitative behaviour, the details of the potential are not important provided the basic form is physically sensible. However, if one requires quantitative results, it is essential to obtain an accurate potential model. In the past, such models have been obtained by empirical fitting to suitable bulk crystal properties. Increasingly, they are now being calculated. This paper reviews the relative merits of these approaches and suggests criteria that any "good" potential should fulfill.