Analysis of spurious image forces in atomistic simulations of dislocations

Molecular dynamics simulations of dislocation/obstacle interactions are enhancing our physical understanding of plasticity. However, despite increasing computational power, the interaction between simulation cell boundaries and the long ranged fields of dislocations make spurious image effects inevitable. Here, these image effects are examined in detail, providing a general map of the spurious image stress as a function of simulation cell size, aspect ratio and bow-out for both nominally edge and screw dislocations. This is achieved using an approximate image solution of the resulting boundary value problem as well as an analytic model that captures most of the spurious image effects. A unique simulation cell shape is found to minimize spurious image effects for a fixed simulation volume (i.e. fixed total number of atoms) and specified initial dislocation line length. The results are used to estimate image stress effects in various literature studies involving dislocation bow-out. The image effects are non-negligible. Several case studies involving simulation cell dimensions are shown to converge due to a near-zero scaling of the image stress with respect to the simulation cell dimensions used. Finally, a direct comparison is made between a dislocation bow-out configuration under an applied load in a finite simulation cell and an image-free multiscale simulation of the same problem and the difference is shown to be consistent with our estimated image stresses. Overall, the results here provide guidance for both the development and interpretation of quantitative molecular dynamics studies involving curved dislocation structures

Robust atomistic calculation of dislocation line tension

The line tension Gamma of a dislocation is an important and fundamental property ubiquitous to continuum scale models of metal plasticity. However, the precise value of Gamma in a given material has proven difficult to assess, with literature values encompassing a wide range. Here results from a multiscale simulation and robust analysis of the dislocation line tension, for dislocation bow-out between pinning points, are presented for two widely-used interatomic potentials for Al. A central part of the analysis involves an effective Peierls stress applicable to curved dislocation structures that markedly differs from that of perfectly straight dislocations but is required to describe the bow-out both in loading and unloading. The line tensions for the two interatomic potentials are similar and provide robust numerical values for Al. Most importantly, the atomic results show notable differences with singular anisotropic elastic dislocation theory in that (i) the coefficient of the ln(L) scaling with dislocation length L differs and (ii) the ratio of screw to edge line tension is smaller than predicted by anisotropic elasticity. These differences are attributed to local dislocation core interactions that remain beyond the scope of elasticity theory. The many differing literature values for Gamma are attributed to various approximations and inaccuracies in previous approaches. The results here indicate that continuum line dislocation models, based on elasticity theory and various core-cut-off assumptions, may be fundamentally unable to reproduce full atomistic results, thus hampering the detailed predictive ability of such continuum models

How Thioredoxin Dissociates Its Mixed Disulfide

The dissociation mechanism of the thioredoxin (Trx) mixed disulfide complexes is unknown and has been debated for more than twenty years. Specifically, opposing arguments for the activation of the nucleophilic cysteine as a thiolate during the dissociation of the complex have been put forward. As a key model, the complex between Trx and its endogenous substrate, arsenate reductase (ArsC), was used. In this structure, a Cys29Trx-Cys89ArsC intermediate disulfide is formed by the nucleophilic attack of Cys29Trx on the exposed Cys82ArsC-Cys89ArsC in oxidized ArsC. With theoretical reactivity analysis, molecular dynamics simulations, and biochemical complex formation experiments with Cys-mutants, Trx mixed disulfide dissociation was studied. We observed that the conformational changes around the intermediate disulfide bring Cys32Trx in contact with Cys29Trx. Cys32Trx is activated for its nucleophilic attack by hydrogen bonds, and Cys32Trx is found to be more reactive than Cys82ArsC. Additionally, Cys32Trx directs its nucleophilic attack on the more susceptible Cys29Trx and not on Cys89ArsC. This multidisciplinary approach provides fresh insights into a universal thiol/disulfide exchange reaction mechanism that results in reduced substrate and oxidized Trx

Atomistic dislocation core energies and calibration of non-singular discrete dislocation dynamics

The total energy of an atomistic dislocation includes contributions from the inelastic/large-distortion 'core' region. Capturing this inelastic 'core' energy is important, especially for dislocations with a curvature in the 10–100 nm scale. Current implementations of discrete dislocation dynamics (DDD) mesoscale simulations either approximate or neglect the core energy and so do not provide consistency with fully-atomistic studies. Using established interatomic potentials for FCC metals, the total dislocation energy is computed directly in atomistic simulations of straight dislocations and a core energy at any desired cut-off core radius is obtained as a function of dislocation character. A proper introduction of the atomistic core energy into the ParaDiS DDD code that uses a non-singular theory (Cai et al 2006 J. Mech. Phys. Solids 54 561–87) is then presented. The resulting atomistically-informed ParaDiS DDD is used to simulate the periodic bow-out of edge and screw dislocations in near-elastically-isotropic aluminum at various length and stress, with comparisons to fully-atomistic simulations. Generally good agreement is obtained between DDD and atomistics, with the best agreement achieved using a non-singular regularization parameter in the range of a = 5 – 10b. The analysis is then extended to compute the core energy of the Shockley partial dislocations that arise in the dissociation of perfect dislocations in fcc metals