Cross-slip Process Research Articles

Dislocation constriction and cross-slip: Anab initiostudy

A novel model based on the Peierls framework of dislocations is developed. The new theory can deal with a dislocation spreading at more than one slip planes. As an example, we study dislocation cross-slip and constriction process of two fcc metals, Al and Ag. The energetic parameters entering the model are determined from ab initio calculations. We find that the screw dislocation in Al can cross-slip spontaneously in contrast with that in Ag, which splits into partials and cannot cross-slip without first being constricted. The dislocation response to an external stress is examined in detail. We determine dislocation constriction energy and critical stress for cross-slip, and from the latter, we estimate the cross-slip energy barrier for the straight screw dislocations.

Non-basal slip systems in HCP metals and alloys: source mechanisms

A possible source mechanism for non-basal 〈 c+ a〉 slip dislocations is proposed based on the formation of an attractive junction between glissile 〈 a〉 and sessile c dislocations from the prism plane into a pyramidal plane. The driving force for the junction formation, which comes from the long-range elastic interaction between c and 〈 a〉 dislocations, is relatively large in most hexagonal close-packed (hcp) metals. Beryllium, which has an unusually low Poisson's ratio, is an exception to this rule. The cross-slip process is energetically unfavorable in Mg and Ti, from a viewpoint of the change in anisotropic elastic line tension, but it becomes favorable in Ti at elevated temperatures above 300°C. Discussion is given on intrinsic stacking fault energies and kinetic aspects of the cross slip.

Large scale atomistic simulations of screw dislocation structure, annihilation and cross-slip in FCC Ni

Using QM-Sutton-Chen many-body potential, we have studied the 1/2 a〈1 1 0〉 screw dislocation in nickel (Ni) via molecular dynamics (MD) simulations. We have studied core energy and structure using a quadrupolar dislocation system with 3D periodic boundary conditions. The relaxed structures show dissociation into two partials on {1 1 1} planes. The equilibrium separation distance between the two partials is 2.5 nm, which is larger than the derived value according to experimental data, due to low stacking fault energy given by the QM-Sutton-Chen force field. From our calculations, the core energy for the 1/2 a〈1 1 0〉 screw dislocation is 0.5 eV/ b. We also studied motion and annihilation process of opposite signed dislocations. We build the dipole system with two combinations of dissociation planes: (a) two dislocations dissociated on intersecting slip planes and (b) two on parallel planes. The process of cross-slip and associated energy barriers are also calculated from these simulations.

Atomistic simulations of cross-slip of jogged screw dislocations in copper

We have performed atomic-scale simulations of cross-slip processes of screw dislocations in copper, simulating jog-free dislocations as well as different types of jogged screw dislocations. Minimum-energy paths and corresponding transition state energies are obtained using the nudged-elastic-band path technique. We find low barriers and effective masses for the conservative motion along the dislocations of elementary jogs on both ordinary {111}<110> and nonoctahedral {110}<110> slip systems. The jogs are found to be constricted and therefore effectively act as pre-existing constrictions; the cross-slip activation energy is thereby dramatically reduced, yielding values in agreement with experiment.

Comment on 'Atomistic simulation of cross-slip processes in model fcc structures'

Abstract Recently Rao et al. (1999, Phil. Mag. A, 79, 1167) claimed reasonable agreement between their results for cross-slip activation energies (CSAEs) in Ni (from atomistic simulation) and in Cu (from scaling of the Ni result) and experimental results. Furthermore, the atomistic result for Cu by Rasmussen et al. (1997, Phys. Rev. B, 56, 2977) was described as deviating ‘significantly’ from theirs. This comment demonstrates that Rao et al.'s scaling approach is too simple and that it is possible to obtain good agreement between atomistic estimates of CSAES with a more accurate, yet still approximate, scaling approach. It is noted that no experimentally derived value for the CSAE in Ni has been published.

Reply to Comment on 'Atomistic simulation of cross-slip processes in model fcc structures'

Reply to Comment on 'Atomistic simulation of cross-slip processes in model fcc structures'

Reply to comment on ‘atomistic simulation of cross-slip processes in model fcc structures’

Abstract The procedure used by Rao et al. (1999, Phil. Mag. A. 79, 1167) for evaluating the cross-slip energy in fcc Cu is explained in detail.

Comment on ‘Atomistic simulation of cross-slip processes in model fee structures’

Abstract Recently Rao et al. (1999, Phil. Mag. A, 79, 1167) claimed reasonable agreement between their results for cross-slip activation energies (CSAEs) in Ni (from atomistic simulation) and in Cu (from scaling of the Ni result) and experimental results. Furthermore, the atomistic result for Cu by Rasmussen et al. (1997, Phys. Rev. B, 56, 2977) was described as deviating ‘significantly’ from theirs. This comment demonstrates that Rao et al.'s scaling approach is too simple and that it is possible to obtain good agreement between atomistic estimates of CSAES with a more accurate, yet still approximate, scaling approach. It is noted that no experimentally derived value for the CSAE in Ni has been published.

Deformation and fracture of the PWA 1472 superalloy single crystal

The deformation and fracture of a PWA 1472 single crystal oriented in a soft orientation were studied by analyzing the evolution of slip bands and detailed fractography of tensile specimens tested at low (10 −4/s) and high (1.0/s) strain rates at room temperature. The results of this study indicated that the slip did not start on the system with the highest Schmid factor. This yielding anisotropy has been attributed to the deformation of the γ′-Ni 3Al precipitates. The plastic deformation was heterogeneous and occurred by the propagation of a Lüders band. Cleavage fracture occurred along the kink bands within a coarse shear band that developed after the activation of the secondary slip system. An increase in the strain rate did not affect the yield strength significantly, but it enhanced the deformation localization, and hence, decreased the ductility and increased the percentage of the cleavage facets. The effect of strain rate is discussed in terms of cross-slip processes.

Hydrogen effects on the character of dislocations in high-purity aluminum

The effect of hydrogen in solid solution on the edge/screw character of dislocations has been investigated by deforming samples of high-purity aluminum in a gaseous hydrogen environment in situ in a controlled environment transmission electron microscope. Solute hydrogen was found to stabilize edge segments of dislocations, which inhibited, and in some cases stopped dislocations from cross-slipping. Removal of hydrogen from the sample allowed the cross-slip process to continue. These observations provide an explanation for the macroscopic observation that hydrogen in solid solution can promote slip planarity.

Atomistic simulation of cross-slip processes in model fcc structures

Abstract Three-dimensional cross-slipped core structures of (a/2)[110] screw dislocations in model fcc structures are simulated using lattice statics within the embedded-atom method (EAM) formalism. Two parametric EAM potentials fitted to the elastic and structural properties of fcc Ni were used for the simulations. The two- and three-dimensional Green's function techniques newly developed by Rao et al, are used to relax the boundary forces in the simulations. Core structures and energetics of the constrictions occurring in the cross-slip process are studied. The core structure of the constrictions are diffuse, as opposed to a point constriction as envisaged by Stroh. The two constrictions formed by cross-slip onto a cross {111} plane have significantly different energy profiles, at variance with classical continuum theory of Stroh. This suggests that self-stress forces and atomistics dominate the energetics of the cross-slip process; the far-field elastic-energy contribution to cross-slip appears to be minimal. However, the Shockley partial separation distances near the constrictions as well as the variation in cross-slip energy with stacking-fault energy are in reasonable agreement with continuum predictions. Cross-slip energies estimated for Cu and Ni from these calculations show reasonable agreement with experimental data. The cross-slip energy shows a significantly weaker dependence on the Escaig stress compared with elasticity calculations. The activation volume for the cross-slip process is estimated to be of the order of 20b 3 at an applied Escaig stress of 10−3μ in Cu, an order of magnitude lower than experimental estimates and continuum predictions.

Atomistic simulation of cross-slip processes in model fcc structures

Atomistic simulation of cross-slip processes in model fcc structures

Simulation of Dislocation Dynamics in Ni3Al: A Study of Velocity Autocorrelations

AbstractThe yield strength anomaly in some L12 compounds has been linked to the thermally assisted cross slip of screw superdislocations. This work continues earlier efforts to understand the yield strength anomaly in L12 alloys using computer simulations of dislocation motion. Dislocations are modelled within isotropic elasticity theory, and simple rules are used to model the cross-slip process in the two dimensional geometry of the simulation. The velocity of a single dislocation in Ni3Al is studied as a function of the applied stress. The observed velocities vary nonlinearly with the applied stress. Further, dislocations are observed to become immobile for small applied loads. At high stresses, the dislocations are observed to advance relatively unhindered by the thermally activated cross slip process. Fluctuations in the velocity of the dislocations are studied, and their autocorrelation function shows an increased correlation time near a threshold stress. This threshold stress is identified with the critical stress proposed in earlier works.

Atomic Scale Simulation of Cross Slip and Screw Dislocation Dipole Annihilation

AbstractAtomistic simulations are used to study cross slip of a single screw dislocation as well as screw dislocation dipole annihilation in Cu. A configuration space path techniquex is applied to determine, without presumptions about the saddle point, the minimum energy path of transition for cross slip. The cross slip process is that proposed by Friedel and Escaig, and the energy of the in-plane constriction initiating cross slip is determined. A minimum stable dipole height much smaller than previously inferred from experimental studies is found. Relaxed screw dislocation dipoles adopt a skew configuration due to the anisotropy of Cu. The path technique is applied to investigate annihilation of stable screw dislocation dipoles, and the energy barrier for annihilation as a function of dipole height is determined for both homogeneous and heterogeneous cross slip leading to the annihilation. The results might be used as quantitative input into meso-/macro-scopical modelling approaches which rely on parameters deduced from either simulation or experiment.

Atomistic Simulations of Cross-Slip Processes in Model FCC Structures and L10 TiAl

AbstractThree dimensional (3D) cross-slipped core structures of a/2[110] screw dislocations in model FCC structures are simulated using lattice statics within the Embedded Atom Method (EAM) formalism using potentials fitted to the elastic and structural properties of FCC Ni as well as L10TiAl. 2 and 3-D Green's function (GF) techniques are used to relax the boundary forces in the simulations. The core structure of the constrictions are diffuse. At large separation distances between Shockley partials in the unconstricted configuration, the two constrictions formed by cross-slip onto a cross {111} plane have significantly different energy profiles suggesting that self-stress forces dominate the energetics of the cross-slip process. The variation in cross-slip energy with stacking-fault energy is in reasonable agreement with continuum predictions, excepting at high fault energy values as in L10TiAl. Cross-slip energies estimated for Cu, Ni and γ-TiAl from these calculations show reasonable agreement with experimental data. The cross-slip energy shows a significantly weaker dependence on Escaig stress as compared to elasticity calculations, with the activation volume for the cross-slip process being approximately 20b3 at an applied Escaig stress of 10-3μ in Cu.

In-situ direct measurement of activation energies for the generation of misfit dislocations in the InGaAs/GaAs (001) system

In-situ X-ray topography (XRT) studies of misfit dislocation generation and movement in epitaxial InGaAs strained-layer structures on (001) GaAs are described. Examination of the changes in dislocation structure during a series of successive post-growth in-vacuo sample anneals has, for the first time, yielded activation energies of 0.7 and 0.8 eV for the formation of α-and β-misfit dislocations (MDs) by the initial glide of substrate threading dislocations (TDs) in the InGaAs epilayer. The introduction of MDs by this method is supplemented by the presence of an additional MD generation process. The activation energy for this is found to be comparable to that required to initiate the glide of a TD. The XRT studies have also confirmed the existence of MD cross-slip events, where α to β cross-slip was found to have an activation energy of 1.2 eV and to be much more common than the reverse β-α cross-slip process.

On the role of incomplete Kear-Wilsdorf locks in the yield stress anomaly of Ni 3Al

In situ straining experiments in an Ni-25.9at.% Al alloy at room temperature indicate that the density of incomplete Kear-Wilsdorf locks is high during deformation. The model of mechanical properties which is based on their unlocking by an intrinsic cross-slip process is described, and compared with the alternative kink model. The results indicate that the alloy investigated has a higher complex stacking fault energy than stoichiometric Ni 3Al.

Atomistic process of dislocation cross-slip in Ni 3Al

In recent years, a great deal of interest has been given to the L1{sub 2} ordered intermetallic compound Ni{sub 3}Al because of its unusual flow behaviors. Based upon the measurements of mechanical properties and the observed dislocation configurations, several models have been proposed to explain the flow stress anomaly, including the anomalous increase of flow stress with temperature, the strain rate sensitivity, the pronounced orientation dependence and the tension/compression asymmetry. The most important common ground of these models is that the flow stress anomaly is the result of cross-slip of superdislocations form {l_brace}111{r_brace} to {l_brace}100{r_brace} planes, forming so-called Kear-Wilsdorf (KW) locks. Computer simulations of superdislocation core structure in Ni{sub 3}Al indicate that the immobility of KW locks is due to the non-planar core structure of 1/2 superpartials which always spread their cores onto {l_brace}110{r_brace} planes other than the {l_brace}100{r_brace} antiphase boundary (APB) planes. However, up to now the authors are still unaware of how the cross-slip takes place. The purpose of the present study is to reveal the atomistic process of the cross-slip of 1/2 superpartials from {l_brace}111{r_brace} to {l_brace}100{r_brace} planes in L1{sub 2} ordered materials.

A model for the anomalous mechanical properties of nickel-base L12 ordered alloys — II. Cross-slip processes and mechanical properties

A model based on the results of Part I, but taking into account more realistic cross-slip processes, is presented in this second part. It is based on a competition between two different cross-slip mechanisms onto the cube plane. Cross slip and glide over interatomic distances are shown to take place by a kink-pair (or double jog) mechanism, whereas cross slip and glide over larger distances are shown to take place by a cross-slip (or critical bow-out) mechanism. Assuming that the kink-pair mechanism is easier than the cross-slip mechanism, in agreement with microscopic observations, analytical expressions are obtained for the strain and temperature dependence of the yield-stress and strain-hardening coefficient, which are in fairly good agreement with experimental results. The role of the complex stacking fault energy, the orientation effects, and the tension-compression asymmetries, can be explained as in the PPV model. The stress-strain rate sensitivity and the reversibility of flow-stress upon a temperature decrease are also discussed.

Prismatic and basal slip in Ti3Al II. Dislocation interactions and cross-slip processes

Abstract In the first part, dislocation mechanisms controlling individual dislocation glide in prismatic and basal planes were studied and compared with results of atomistic calculations. In this second part, the plasticity of Ti3Al is studied in situ at an intermediate scale. The critical resolved shear stress (CRSS) of prismatic slip is shown to be controlled by the easy glide process of dislocations in the planes of high APB energy, and the CRSS of basal slip is shown to depend on the collective motion of several dislocation families forming a coplanar network. Other mechanical properties like work-hardening and ductility have been interpreted in terms of dislocation interactions, multiplications and cross-slip processes. The high Peierls-type frictional forces are thus surprisingly of minor importance for explaining the mechanical properties of Ti3Al.