MD Simulations Of Deposition Process And Nanoindentation Deformation Behavior Of Cu, Ag And TiN Films

The structures and properties of thin films depend on their preparation process.The vapor deposition is the main preparation method. In addition, the thin filmdeposition is a powerful method to study substances with specific compositions andstructures. On the other hand, the mechanical properties of thin films are importantfactors constraining their performance range and life. Nanoindentation technique iswidely used in various tests of mechanical properties of thin films. In addition,nanoidentation process can induce an extremely small deformation of thin films,which makes nanoidentation to become a powerful method to study the microscopicdeformation behavior of thin films. Due to the current difficulty to carry outexperiments and theoretical research in the small scale, this thesis performedatomistic simulation study on the thin film deposition and microscopic deformationbehavior of thin films during nanoindentation.The thin film deposition was simplified as an atomic deposition process, andthen a molecular dynamics (MD) model consisting of a substrate, incident atoms,and a virtual wall was built to describe it. For describing the atomic collision in thedeposition process of energetic atoms, the embedded-atom method (EAM) potentialof the Cu-Ag system was modified by joining it with Ziegler-Biersack-Littmarkrepulsive potential at the small atomic distance, and a second nearest-neighbormodified EAM (2NN MEAM) potential of the Ti-N system was developed. Basedon the model and potentials, the deposition of Cu film on Ag substrate and Ag filmon Cu substrate under the condition of substrate temperature T=300K and incidentkinetic energy of incident atoms Ei=8eV was simulated, and the deposition of TiNfilm on TiN(100) substrate under the condition of T=300-700K、Ei=0.5-10eV wasalso simulated.The results show that the first deposited Ag layer grows on the Cu substrate bythe layer mode, while the first deposited Cu layer grows on the Ag substrate by theisland mode, determined by the surface energy difference between Cu and Ag. Thefilm is epitaxial up to the coverage reaching about10ML except for Cu/Ag(001).The interface is sharp due to the immiscibility of Ag with Cu. Atoms of Ag(001)substrate diffuse into the Cu film for a longer distance due to their lower surfaceenergy, stronger mobility and lower packing density of the (001) atomic plane. Theinterface displays a (9×9) hexagonal moiré pattern for both Ag/Cu(111) andCu/Ag(111), and the displacement of the superposed Cu atoms again the Ag filmcontributes to the corrugation of the interface. c(10×2) superstructure forms forAg/Cu(001). While for Cu/Ag(001) partial dislocations are activated near the interface, and a great number of stacking faults form in the Cu film with theseriously bent interface.When incident atomic N compacts the TiN(001) surface, incident atoms may beadsorbed or reflected, or surface atoms may be sputtered, determined by thedifferent energy transfer between them. The adsorption and reflection strengthenwith the increase of Ei, resulting in the decrease of the sticking coefficient of N onthe TiN(001) surface. The sticking coefficient of Ti on the TiN(001) surfaceapproximates unity under different values of Ei. The effect of T on the stickingcoefficient is faint. The film grows on the TiN(001) surface via the layer mode andTiN3is the smallest epitaxial island. Only N and Ti vacancies form in the films.Their concentration decreases with increasing T and Ei, resulting from theenhancement of thermal diffusion and kinetic energy assisted athermal diffusion. Toget the stoichiometric TiN film, the incident N:Ti flux ratio should be larger thanunity and be increased with the increasing Ei.Molecular statics model of nanoindentation on thin films was built based on themodel of virtual spherical indenter described by the power function. It wasdiscussed that the effect of the indenter radius R and stiffness constant k of thismodel on the results of nanoindentation simulations. The results indicate that withthe increase of k, the number of iteration steps for the same nominal indentation Hgradually increases, and the depth of indenter penetrated by contact atoms graduallydecreases until reaching a stable value. When k is small, the reduced elastic modulusErfitting to the relation between H and load based on the Hertian elastic expressionwill get smaller, which will be avoided by using the average indentation depth hconsisting of hpen. Locality of materials deformation during nanoindentation makesit more reasonable to measure the elastic deformation using H rather than theindentation strain. The upper fitting bound of Erunder different values of R shouldbe selected as the same value of H, if the critical depths of initial dislocationemission are selected as the upper bound, the value of Erwill decrease with thedecrease of R, resulting in an artificial size effect.The processes of nanoindentation on Cu(111) and Ag(111) thin films weresimulated. The results indicate that the abrupt load drop results from the emissionand slip of multi dislocations during loading, while during unloading dislocationsslip back to the emission source due to the atomic stress from the energetic stackingfaults, resulting in the reversible plasticity. After dislocations on the (111) plane isformed by cross-slip,(111) thin films will have irreversible plasticity. Thenucleation of a tetrahedral sessile lock accounts for the initial dislocation emissionof the Ag(111) film, different from the formation of the half V-shaped dislocationloop in the Cu(111) film. The sequent deformation is dominated by the dislocationsnucleated on the {111} slip planes not parallel to the indented surface. TBs harden the Ag films due to their effective obstacles to the motion of dislocations. Especially,the dislocations nucleated on the (111) slip planes dominate the deformation of filmswhen the TB approximates the indented surface. Furthermore the load for the initialyield is reduced by TBs due to the emission of glissile dislocations on the (111) slipplanes.Projected polygon method with boundary compensation was proposed, bywhich the area calculated is a physical contact area. Unlike the Oliver-Pharr methodused in experiments, this method is independent of the calculation of indentationdepth and the assumption of deformation modes. Hence, the hardness obtained usingthis method is a truly material response, while the Oliver-Pharr method does notwork in the nanometer regime.