Molecular Dynamics simulation of 〈 c + a 〉 dislocation core structure in hexagonal-close-packed metals

The core structures of 〈c+a〉 dislocations in hexagonal-close-packed (hcp) metals have been investigated by molecular dynamics (MD) simulation using a Lennard-Jones-type pair potential. The 〈c+a〉 edge dislocation has two types of core at 0 K; one is a perfect dislocation (type A), and the other has two 1/2 〈c+a〉 partials (type B). Type A transforms to type B by abruptly increasing temperature from 0 K to 293 K, while type B is stable in temperature range from 0 K to 293 K. In contrast, type A extends parallel to (0001) at 30 K, and this extended core is still stable at 293 K. These results suggest that the 〈c+a〉 edge dislocation glides on the as two 1/2 〈c+a〉 partial dislocations and becomes sessile, due to changes of the core structure. The 〈c+a〉 screw dislocation spreads over two planes at 0 K. The core transforms into a unsymmetrical structure at 293 K, which is spread over and , and core spreading occurs parallel to at 1000 K. A critical strain to move screw dislocations depends on the sense of shear strain. The dependence of the yield stress on the shear direction can be explained in terms of these core structures. This article is based on a presentation made in the symposium entitled “Dect Properties and Mechanical Behavior of HCP Metals and Alloys” at the TMS Annual Meeting, February 11–15, 2001, in New Orleans, Louisiana, under the auspices of the following ASM committees: Materials Science Critical Technology Sector, Structural Materials Division, Electronic, Magnetic & Photonic Materials Division, Chemistry & Physics of Materials Committee, Joint Nuclear Materials Committee, and Titanium Committee.     

Molecular dynamics simulation of 〈<Emphasis Type="Bold">c</Emphasis>+<Emphasis Type="Bold">a</Emphasis>〉 dislocation core structure in hexagonal-close-packed metals

The core structures of 〈c+a〉 dislocations in hexagonal-close-packed (hcp) metals have been investigated by molecular dynamics (MD) simulation using a Lennard-Jones-type pair potential. The 〈c + a〉 edge dislocation has two types of core at 0 K; one is a perfect dislocation (type A), and the other has two 1/2 〈c+a〉 partials (type B). Type A transforms to type B by abruptly increasing temperature from 0 K to 293 K, while type B is stable in temperature range from 0 K to 293 K. In contrast, type A extends parallel to (0001) at 30 K, and this extended core is still stable at 293 K. These results suggest that the 〈c+a〉 edge dislocation glides on the {11 2} as two 1/2 〈c+a〉 partial dislocations and becomes sessile due to changes of the core structure. The 〈c+a〉 screw dislocation spreads over two {10 1} planes at 0 K. The core transforms into a unsymmetrical structure at 293 K, which is spread over {11 2} and {10 1}, and core spreading occurs parallel to {11 2} at 1000 K. A critical strain to move screw dislocations depends on the sense of shear strain. The dependence of the yield stress on the shear direction can be explained in terms of these core structures. This article is based on a presentation made in the symposium entitled “Defect Properties and Mechanical Behavior of HCP Metals and Alloys” at the TMS Annual Meeting, February 11–15, 2001, in New Orleans, Louisiana, under the auspices of the following ASM committees: Materials Science Critical Technology Sector, Structural Materials Division, Electronic, Magnetic & Photonic Materials Division, Chemistry & Physics of Materials Committee, Joint Nuclear Materials Committee, and Titanium Committee.     

Transmission electron microscopy investigation of 〈c+a〉 dislocations in Mg and α-solid solution Mg-Li alloys

The ductility of Mg alloys is limited due to a shortage of independent slip systems. In particular, c-axis compression cannot be accommodated by any of the easy slip or twinning modes. Basal-textured samples of pure Mg and Mg-15 at. pct Li were examined for the presence of 〈c+a〉 dislocations by post-mortem transmission electron microscopy (TEM) after a small deformation, which forced the majority of grains to compress nearly parallel to their c-axes. A higher density and more uniform distribution of 〈c+a〉 dislocations is found in the Li-containing alloy. Because the 1/3〈11 3〉 {11 } pyramidal slip mode offers five independent slip systems, it provides a satisfying explanation for the enhanced ductility of α-solid solution Mg-Li alloys as compared to pure Mg. The issue of 〈c+a〉 dislocation dissociation and decomposition remains open from an experimental point of view. Theoretically, the most feasible configuration is a collinear dissociation into two 1/2〈c+a〉 partial dislocations, with an intervening stacking fault on the glide plane. It is speculated that Li additions may lower the fault’s energy and, thereby, increase the stability of this glissile configuration. This article is based on a presentation made in the symposium entitled “Defect Properties and Mechanical Behavior of HCP Metals and Alloys” at the TMS Annual Meeting, February 11–15, 2001, in New Orleans, Louisiana, under the auspices of the following ASM committees: Materials Science Critical Technology Sector, Structural Materials Division, Electronic, Magnetic & Photonic Materials Division, Chemistry & Physics of Materials Committee, Joint Nuclear Materials Committee, and Titanium Committee.     

Microscopic Properties of Long-Period Ordering in Al-Rich TiAl Alloys

The ordering mechanism of long-period superstructures (LPSs) in Al-rich TiAl alloys has been investigated by high-resolution transmission electron microscopy (HRTEM). The LPSs are classified in terms of arrangements of base clusters with different shapes and compositions formed in Ti-rich (002) layers of L1₀-TiAl matrix: square Ti₄Al, fat rhombus Ti₃Al, and lean rhombus Ti₂Al type clusters. The HRTEM observations revealed that antiphase boundaries of long-range-ordered LPS domains and short-range-ordered microdomains are constructed by various space-filling arrangements of the base clusters. Such a microscopic property characterized by the base clusters and their arrangements is markedly analogous to that of the * special-point ordering alloys such as Ni-Mo. This article is based on a presentation given in the symposium entitled “Materials Behavior: Far from Equilibrium” as part of the Golden Jubilee Celebration of Bhabha Atomic Research Centre, which occurred December 15–16, 2006 in Mumbai, India.

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Measuring Stress Distributions in Ti-6Al-4V Using Synchrotron X-Ray Diffraction

This article presents a quantitative strain analysis (QSA) study aimed at determining the distribution of stress states within a loaded Ti-6Al-4V specimen. Synchrotron X-rays were used to test a sample that was loaded to a uniaxial stress of 540 MPa in situ in the A2 experimental station at the Cornell High Energy Synchrotron Source (CHESS). Lattice-strain pole figures (SPFs) were measured and used to construct a lattice strain distribution function (LSDF) over the fundamental region of orientation space for each phase. A high-fidelity geometric model of the experiment was used to drastically improve the signal-to-noise ratio in the data. The three-dimensional stress states at every possible orientation of each α (hcp) and β (bcc) crystal within the aggregate were calculated using the LSDF and the single-crystal moduli. The stress components varied by 300 to 500 MPa over the orientation space; it was also found that, in general, the crystal stress states were not uniaxial. The maximum shear stress resolved on the basal and prismatic slip systems of all orientations within the α phase, was calculated to illustrate the utility of this approach for better identifying “hard” and “soft” orientations within the loaded aggregate. Orientations with low values of which are potential microcrack initiation sites during dwell fatigue conditions, are considered hard and were subsequently illustrated on an electron backscatter diffraction (EBSD) map. This article is based on a presentation given in the symposium entitled “Neutron and X-Ray Studies for Probing Materials Behavior” which occurred during the TMS Spring meeting in New Orleans, LA, March 9–13, 2008, under the auspices of the National Science Foundation, TMS, the TMS Structural Materials Division, and the TMS Advanced Characterization, Testing, and Simulation Committee.

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