1. Medea Vasp Download Torrent
2. Vasp Potcar Download

Searching for new Heusler alloys with novel electronic and magnetic properties is a recent topic of interest to researchers. Using density functional theory based calculations, we study the effect of the replacement of Ga by Al at the A-site of A 2 BC-type alloy, Ga 2 MnNi. By analyzing the magnetic properties of the Ga 2-x Al x MnNi system (with values of x between 0 and 2), we find that in.

Cross-platform TBStudio is a powerful quantum technical software package to construct Tight-Binding (TB) model for nano-scale materials. Starting from the simplified linear combination of atomic orbitals method in combination with first-principles calculations (such as OpenMX or Vasp packages), one can construct a TB model in the two-centre approximation. Using Slater and Koster approach we calculate the TB Hamiltonian of the system and use a nonlinear fitting algorithm to find the best entries. This project is a small program using Qt for calculation and visualization of electron charge densities in crystals and molecules. It can create a 2D plane of charge densities or create a 1D representation of the charge densities along a given line, in both cases it can also export the graphical representations to various image formats. The program can import various file formats such as space 3D grids of charge density values, created in a quantum chemistry programs Gaussian, Crystal and VASP.

Electronic structure calculations have become a powerful foundation for computational materials engineering. Four major factors have enabled this unprecedented evolution, namely (i) the development of density functional theory (DFT), (ii) the creation of highly efficient computer programs to solve the Kohn-Sham equations, (iii) the integration of these programs into productivity-oriented computational environments, and (iv) the phenomenal increase of computing power. In this context, we describe recent applications of the Vienna Ab-initio Simulation Package (VASP) within the MedeA ® computational environment, which provides interoperability with a comprehensive range of modeling and simulation tools. The focus is on technological applications including microelectronic materials, Li-ion batteries, high-performance ceramics, silicon carbide, and Zr alloys for nuclear power generation. A discussion of current trends including high-throughput calculations concludes this article.Keywords:,.

More than ever before, our society depends on a perplexing multitude of materials to meet needs such as housing, heating/cooling, clean water, production of food, energy, infrastructure, communication, transportation, and health care, as well as allowing recreational and artistic activities. It is fair to say that until now the materials necessary for all these purposes have been developed by experimental methods. Although the fundamental physical and chemical laws that govern the properties of materials are known, the goal of designing materials with specific properties by a purely theoretical/computational approach remains elusive.

The experimental development of novel materials remains an expensive and time-consuming activity and will become even more so as the requirements for high performance, low cost, and environmental compatibility become more stringent. Hence, any technology which can focus and accelerate the improvement of existing and the development of new materials is highly valuable. This is indeed the role and the challenge for computational materials engineering, as has been expressed in the concept of “Integrated Computational Materials Engineering” (ICME).

Simulations on the macroscopic scale are a well-established engineering practice, especially in structural analysis, in computational fluid dynamics (CFD), and in electrical circuit design. In these highly successful macroscopic simulations, materials properties data have traditionally been taken from experiment. As the predictive power of atomistic simulations increases, computed materials property data can become the input for these macroscopic simulations. This opens the exciting and unprecedented opportunity to close the loop from system design to materials design as illustrated in. However, as explained above, it is unrealistic to expect that one can completely replace experiment by simulations. Rather, one needs to create the synergy between atomistic simulations and experiments in the positive spirit of yin-yang or taegeuk.

In complementary metal oxide semiconductor (CMOS) technology, steadily diminishing device sizes have mandated the introduction of high-k dielectrics such as hafnium dioxide, which are replacing pure silicon dioxide dielectric layers. As a consequence, to maintain a low threshold voltage for switching, the material for the gate metal has had to be changed. Titanium nitride has emerged as a suitable choice in this role.

A key requirement for energy efficient switching of CMOS devices is the alignment of the Fermi level ( i.e. The energy of the highest occupied states) of the metallic gate with the band edges of the semiconducting channel of the device, as illustrated in. Empirically it was found that annealing of the as-deposited TiN in an oxygen atmosphere increased the work function, as desired.

Secondary ion mass spectroscopy (SIMS) measurements showed that oxygen atoms had penetrated into the layer of TiN. It was thus concluded that the replacement of oxygen in TiN causes the increase of the work function. Using MedeA ®-VASP, detailed electronic structure calculations of models of the HfO 2/TiN interface revealed, however, that replacement of N by O inside the TiN layer did not change the work function ( cf. One could have concluded that computed results and experiment are in contradiction. Actually, this is not the case. While annealing of the stack in an oxygen-containing atmosphere leads to ingress of oxygen into the TiN layer, the O atoms inside the TiN layer are not the cause for the work function increase.

Rather, calculations revealed that, driven by the ingress of O atoms, the diffusion of N atoms, and the filling of O vacancies in the HfO 2 layer, the replacement of O atoms by N atoms exactly at the interface between HfO 2 and TiN caused a dramatic increase of the work function, thus reconciling computations and experiment. The origin of this behavior is the different chemical interaction of oxygen vs.

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Nitrogen with the transition metal atoms Hf and Ti. Changes in the distribution of electronic charges between the HfO 2 and TiN layers at the interface determine the effective work function. This detailed understanding and control of the chemistry at the interface is thus critical to fabrication processes of energy efficient transistors.

Silicon nitride is a fascinating ceramic material with a wide range of applications including engine parts and ignition systems in cars, ball bearings, for example in wind turbines, in rocket thrusters due to its resistance to thermal shocks, but also in medical orthopedic devices and in the semiconductor industry as insulator and diffusion barrier. As in many practical applications, interfaces play a critical role. An illustrative example in this context is the strength of an interface between aluminum and silicon nitride. Here, the coefficient of thermal expansion is computed on the ab initio level using the so-called quasi-harmonic approximation. From phonon calculations for a range of different lattice parameters one obtains the vibrational entropy. Combined with the electronic energy this in turn gives expressions for the Gibbs free energy as a function of temperature.

As the temperature is increased, the minimum of the Gibbs free energy shifts to larger lattice parameters. Analysis of this temperature dependence gives the coefficient of thermal expansion as used, for example, in the case of Mg 2SiO 4. Developed by K. Parlinski, the integration and automation of this capability within MedeA ® greatly facilitates this task. The volume change of active materials that accompanies charge and discharge of Li-ion batteries is a major source of degradation which limits the overall lifetime of such a battery.

While a zero-strain anode material exists, namely Li 4Ti 5O 12, there have not been any suitable zero-or low-strain materials for cathodes. By using systematic DFT calculations, three low-strain materials have been found within the class of LiMn xCr yMg zO 4. The most promising materials have been synthesized and characterized by X-ray diffraction and electrochemical techniques. The results are consistent with the ab initio predictions. According to the boron-carbon phase diagram, a boron carbide phase exists between approximately 8 at% and 20 at% C with a melting point reaching 2450°C. The crystal structure of boron carbides consists of icosahedra connected with short linear rods of three atoms.

However, the distribution of the C atoms in this structure is far from obvious. Clark and Hoard give a structure for B 4C where the icosahedra consist only of B atoms connected with C-C-C linear rods. For the boron-rich compound B 13C 2 the structure given by Larson shows B 12 icosahedra connected with linear rods of the composition C-B-C. The optical properties in the spectral range of visible and ultraviolet light are determined by electronic transitions from occupied to unoccupied states. Quantitative predictions of these states require a level of theory beyond standard density functional calculations.

So-called hybrid functionals such as HSE06, offer a practical approach to compute excitation energies. Using this approach in VASP and the optical analysis tools in MedeA ®, the computed refractive index of Y 2O 3 (yttria) is in good agreement with experimental data, as illustrated in. The formation of zirconium hydrides is of high concern in the operation of nuclear reactors.

Corrosion of zirconium alloys used in the core of nuclear power plants produces hydrogen, and a fraction of the hydrogen diffuses into the zirconium material. When the hydrogen concentration exceeds the terminal solid solubility, the excess hydrogen starts to precipitate as hydrides. This process may lead to embrittlement with crack formation due to lower ductility of the hydrides than that of the Zr matrix. VASP as integrated in the MedeA ® computational environment has been employed to study structural, thermodynamic, and elastic properties of the Zr-H system. The computational accuracy of this method is needed to quantify and determine the behavior of hydrogen in Zr. This becomes clear considering the small energy difference between the octahedral and tetrahedral sites for hydrogen in the Zr lattice. The electronic total energy difference between the sites is computed to be only 5.9 kJ/mol, with the tetrahedral site being energetically favorable.

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Vibrational effects can readily be added using MedeA ®-Phonon. Inclusion of vibrational effects change the energy difference between the sites to 0.5 kJ/mol at T = 0 K and to 8.6 kJ/mol at 600 K. The calculations also show that the hydrogen solubility increases under tensile strain and decreases under compressive strain.

This leads to hydrogen migration and accumulation in expanded regions of the Zr lattice resulting in hydride precipitation. Examples of regions under tensile stress can be at a Zr/ZrO 2 interface, at the front of a crack tip, or even in regions around Zr self-interstitial atoms. Furthermore, hydrogen is attracted to Zr vacancies and voids. The simulations show that up to six hydrogen atoms are strongly bound inside a single Zr vacancy.

Clustering of vacancies into dislocation loops can lead to regions with very high local hydrogen concentration. The simulations show that hydrogen inside the vacancy loops can delay or in some cases even prevent collapse of the loops. Each of these situations lead to regions highly supersaturated with hydrogen and could be potential nucleation sites of zirconium hydrides. A systematic study of the zirconium hydrides has been performed by successively filling tetrahedral sites in the zirconium lattice by hydrogen, probing a large number of configurations for H-Zr ratios between zero (pure α-Zr) up to complete filling of the sites at a ratio of 2.0 (ɛ-ZrH 2). Computation of the elastic properties of the hydrides is conveniently carried out using the automated approach. Some of the hydride structures display elastic instability, such as cubic δ-hydride with full hydrogen occupancy which can be stabilized by introducing vacancies on the hydrogen sites or by a tetragonal distortion into ɛ-ZrH 2. The elastic moduli of the most stable hydrides at each stoichiometry are shown in.

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The bulk modulus increases almost linearly with hydrogen concentration from pure α-Zr to ZrH 2. The shear moduli of the hydrides are similar to that of α-Zr while Young’s moduli of the hydrides typically are lower than for α-Zr. The clear exception is ZrH 1.25 which has high elastic moduli. This is identified as a γ-hydride of P4 2/mmc symmetry. The final example is the formation of nanotunnels on surfaces of silicon carbide.

The discovery of this type of surface features is the result of a close interaction between precise experimentation and systematic density functional calculations as detailed below. Silicon carbide is a fascinating ceramic material with a range of practical applications. It is a wide band gap semiconductor of interest for power electronics and high-frequency applications. The material remains operational at high temperatures and it is resistant to radiation. In nuclear energy technology, silicon carbide is used as fuel cladding, for example in the TRISO fuel.

Furthermore, silicon carbide is a bio-compatible ceramic of interest to medical applications. While the ingredients, silicon and carbon, are readily available, the synthesis of high-purity SiC wafers requires highly sophisticated approaches and the hardness of SiC makes processing difficult. On a Si-rich surface of cubic SiC the release of the surface stress gives rise to a remarkable behavior, when this surface is exposed to atomic hydrogen, namely the formation of nanotunnels.

When a clean SiC(001) Si-rich 3 × 2 reconstructed surface is exposed to atomic hydrogen, the most reactive Si atoms in the top layer react strongly exothermally by forming the structure 2H shown in. A plausible binding site for the subsequent H atoms seems to be the bridge position between the Si atoms in the third layer. In fact, this has been assumed in earlier investigations, but it leads to inconsistencies with experimental vibrational data. In particular, the vibrational modes of H-atoms in the bridge position are far too low to explain the presence of a high-frequency Si-H stretching mode which would be expected for Si atoms bonded to C atoms in the form C-Si-H. Using systematic ab initio calculations with VASP in MedeA ®, another reaction scheme has been identified as illustrated in Rather than binding in a bridge position in the third layer of Si atoms, H atoms can bind exothermally to Si atoms in the second layer forming a structure denoted 6H.

This reaction is remarkable, because it breaks Si-Si bonds between the second and third layer leading to an outward movement or “puckering” of the Si atoms in the second layer. Furthermore, the computed frequencies are also consistent with the notion that Si-H stretch frequencies are shifted to higher values if the Si atoms are bonded to C atoms. Earlier experiments using infrared spectroscopy showed absorption at 2118 and 2140 cm −1 (Δν = 22 cm −1). Computations using the nanotunnel model result in frequencies for the stretch modes Si1a-H and Si3a-H of 2087 cm −1 (not marked explicitly in ) and 2120 cm −1 (Δν = 33 cm −1) as discussed in Ref.

The earlier bridge-bonded model is inconsistent with these experimental data. Thus, ab initio calculations have been essential in the clarification of the remarkable nanotunnel structure of a silicon carbide surface. During the past decades we have witnessed steady progress in computational materials engineering of growing industrial value. It is probably fair to say that the ability to compute total energies of ensembles of any types of atoms using density functional theory is a cornerstone for this remarkable development. This fundamental capability has enabled in-depth understanding of rather complex systems and the prediction of a range of materials properties as illustrated in the previous section for selected cases. As example, for electronic applications we have presented a computational analysis of the interface between HfO 2 and TiN in the context of enhancing the efficiency of transistors with high-k dielectrics. The strength of metal/ceramic interfaces as a function of the interface chemistry has been illustrated here for the case of Al/Si 3N 4 together with the ab initio calculation of thermal expansion coefficients.

The accurate prediction of lattice parameters of compounds such as transition metal oxides in the spinel structure has enabled the design of low-strain cathode materials for Li-ion batteries. An application of this methodology to boron carbides using the cluster expansion method has helped to clarify the energetically preferred arrangement of boron and carbon atoms. Furthermore, the computation of the phonon dispersions of boron carbide has allowed the assignment of vibrational frequencies to specific atomic motions, such as the oscillation of boron atoms in linear C-B-C rods in boron carbides. Building on the results of DFT calculations, it is possible to predict optical properties such as the refractive index as a function of energy, which has been shown here for yttria. In the case of metal alloys, this ab initio approach provides detailed insight into the interaction of hydrogen atoms with a metal leading to the formation of hydrides and embrittlement, as illustrated here for the case of hydrogen in zirconium. In surface science, ab initio calculations have proven to be an invaluable tool to unravel complex surface reconstructions and to investigate the interaction of atoms and molecules with surfaces.

In fact, an important industrial application of ab initio solid state calculations is related to heterogeneous catalysis. Here, the application to a surface science problem has been demonstrated for the case of H atoms inducing the formation of nanotunnels on Si-rich 3C-SiC(001) 3 × 2 reconstructed surfaces. In combination with precise experimental determinations of the vibrational properties of this system, ab initio calculations have revealed energetically favorable structures and have allowed the characterization of these structures by aligning experimental and computed vibrational frequencies, thus leading to the discovery of novel nanotunnel features. One can say that during the past four or five decades, density functional calculations have evolved to a mature level. In fact, a recent systematic comparison of the major DFT codes resulted in a remarkable consistency of the computed results despite quite different algorithmic implementations. What is novel and perhaps revolutionary at present is the ease and rate new data can be computed for increasingly complex systems.

Stimulated by the Materials Genome Initiative in the USA, high-throughput calculations for hundreds of thousands of compounds have become possible and are driving research in a number of academic groups in the world. This is truly exciting and one can expect new algorithms, new computational procedures, and new forms of data analytics to be developed.

Large databases of computed results are being created for existing as well as hypothetical structures. Mining this richness of information will undoubtedly reveal new insights and novel materials. While this large volume of data is valuable, it will not resolve other major challenges of computational materials engineering, namely (i) the multi-scale and multi-physics aspects, (ii) the inherently non-equilibrium character of materials, and – last but not least – (iii) the accuracy of the ab initio calculations.

Let us consider these aspects as they provide a perspective on necessary future developments. In the overwhelming number of cases, materials engineering has a multi-scale and multi-physics character. For example, if one is interested in designing fracture resistant high-performance ceramics, one needs models of a polycrystalline material which are the result of a sintering process. Such models need to incorporate information about grain size, porosity, composition and properties of the crystalline grains with their defects, the structure of the intergranular interfaces, chemical segregation effects, residual stresses and perhaps charges trapped in defects.

The fracture process involves long-range strain fields, elastic and plastic deformations on the mesoscale, grain boundary sliding, crack propagation, bond breaking at crack tips, and diffusion processes. Bond breaking and chemical rearrangements may cause local changes in vibrational energy which may entail thermal transport phenomena. Quantitative and statistically relevant modeling and simulations of such a system are all but routine with current simulation technology and yet the above case is “simple” from an engineering point of view. A more complicated case would be stress corrosion cracking involving an aqueous phase, which adds electrochemical aspects. While a decomposition of such materials problems into coupled discrete and continuum models is tempting, the dynamic fluctuations of atomistic and continuum domains, the vast configurational space, and the many orders of time scales involved make such a direct approach challenging. Such examples point to the need for novel and innovative theoretical and computational approaches combining coarse-graining in length and time scales in moving from electronic structure calculations to a continuum description with “fine-graining” in regions such as crack tips, where atomic-scale phenomena are decisive.

Finally, there remains the question of accuracy of current ab initio methods. While quantum chemists have developed approaches, which - at least in principle - can be converged to any desired degree of accuracy, this is not the case for the current form of density functional calculations. No systematic and practical way is known today to converge to the exact density functional. The limitations of DFT-GGA calculations are fairly well known, but despite intense efforts by a number of leading research groups in the world, there is no systematic and practical ab initio many-body approach which would allow one to compute, for example, the temperature of solid-liquid phase transitions to within a few degrees even for systems such as pure silicon. This present situation is by no means unusual in the evolution of science and technology. Rather, it should be taken as stimulus for pushing the frontiers of science. One also has to keep in mind that engineering in all its forms is not and never will be perfect in all respects, but good engineering implies reliable control of the limits.

It is not necessary to predict the exact yield stress of a particular sample of a material. What is required is the knowledge of the upper bound where this sample will not fail. This general engineering principle is a good guide in computational materials engineering.

Rather than seeking the ultimate accuracy, one needs computational protocols and approaches, which give reliable boundaries. This means a clear understanding of the key physical and chemical mechanisms which determine the properties of a material. The true art of computational materials engineering is knowing which aspects can be neglected while keeping the key characteristics of the problem.

Sophisticated computational software environments with all their tools and capabilities facilitate this task, but in final analysis the best multi-scale and multi-physics tool remains the creative human mind abetted by the finest tools developed by the combined effort of the scientific and engineering disciplines. Left panel: schematic diagram of the energy levels across a stack of semiconducting Si (the channel of the transistor), the gate oxide consisting of SiO 2 and HfO 2, and the TiN gate metal. To achieve optimal energy consumption of a p-doped CMOS, the Fermi level of the metal has to be aligned with the top of the valence band of the semiconductor.

This is achieved by increasing the work function of TiN. Right panels: (a) atomistic model of HfO 2/TiN interface with some N replaced by O; (b) O atoms are replaced by N atoms only at the HfO 2/TiN interface. Panels (c) and (d) show the change of the work function due to the chemical changes in the system.