Dipole Annihilation Research Articles

Annihilation of edge dislocation dipoles by a glide mechanism

A glide mechanism for the annihilation of an edge dislocation dipole is proposed. It can operate in many crystal structures. Evidence is taken from transmission electron microscope investigations of plastically deformed NaCl. The mechanism can explain the decrease in accumulation rate of primary dislocations as observed during deformation beyond stage I of the work-hardening curve.

The enhancement of creep by irradiation

Three new processes are proposed for the enhancement of creep by irradiation. They all depend on the assumption that the rate controlling process in steady-state creep is the annihilation of edge dislocation dipoles in sub-grain boundaries by climb. The relaxation or misfit volume strain is much greater for an interstitial than for a vacancy and leads to a preferential loss of interstitials at isolated edge dislocations in irradiated materials. This biassed loss of interstitials constitutes the first process and causes edge dislocations of opposite sign present in a boundary to climb in opposite directions and thereby increases the rate at which they collide with each other and annihilate. The resulting creep rate increases with stress less rapidly than does the corresponding creep process in the absence of irradiation. In the second process the positive relaxation volume strain of an interstitial ensures that it will be attracted into regions where the tensile fields of adjacent dislocations overlap and reinforce each other. This dislocation dipole size effect interaction always favours climb in a direction which leads to dislocation annihilation. In addition to the size effect interaction the tetragonal interaction associated with the dumb-bell configuration of interstitials. has been evaluated for both these processes. It is shown in the appendix that if the interstitial has a very small effective shear modulus, which is also a consequence of the dumb-bell configuration, the tetragonal interaction also becomes insignificant. In the third process, the overlapping shear stress fields of two neighbouring dislocations of opposite sign attract interstitials by the inhomogeneity interaction. This arises because a dumb-bell interstitial has a very low effective shear modulus compared to the surrounding material. The increased concentration of interstitials also leads to dislocation climb which is preponderantly in the direction which favours annihilation. This process is a second order effect which is always much less than the size effect which controls the first two processes. In a fourth process, the energy barriers which impede the absorption of interstitials (usually at jogs) on dislocation of different systems are affected differently by an applied stress, and this leads to a rate of creep which, like the established mechanism of stressinduced preferential absorption, is to a first approximation linear in both the rate of irradiation and the applied stress.

Dislocation Dipole Annihilation in Pure Sapphire (Al2O3)

Dislocation dipoles are often formed during deformation of sapphire (Al2O3) undergoing basal slip by edge-trapping of 1/3<1120> dislocations. Tne dipole concentration remains constant with increasing plastic deformation because the formation of new dipoles due to increasing plastic strain is offset by dipole annihilation by diffusive processes. Dipole annihilation in sapphire occurs in five distinguishable steps:(1) the dipole width fluctuates by self-climb, forming narrower and wider segments; (2) narrow segments transform into faulted dipoles by self-climb; (3) both perfect and faulted dipoles pinch off to form prismatic loops, primarily by self-climb; (4) faulted loops and dipoles rotate into edge orientation; and (5) loops grow or annihilate by climb. The kinetics of dipole annihilation thus depends on both self-climb and climb, that is, on both pipe-diffusion and bulk-diffusion kinetics, and so quantitative analysis of these changes during annealing will give valuable information on diffusion in sapphire. The objective of this paper is to present a brief qualitative analysis of the kinetics of dislocation dipole annihilation in pure sapphire.

Energies d'activation des divers mecanismes de deformation par fluage de monocristaux et de polycristaux d'uranium entre 150° et 770° C

Creep tests were performed at high temperature (150–770° C) on mono- and polycrystals of uranium in the α and β forms. The variation of the activation energy for creep with temperature led to the establishment of temperature ranges in which different mechanisms were rate-controlling. Micrographie observations made after creep and the stress-dependence of creep rate confirm the hypothesis based on activation energies. Below 325° C, the creep behaviour of mono- and polycrystals is similar. From 325° C to an upper temperature limit which depends on the structural condition of the metal and on its degree of purity, the creep deformation of uranium is controlled by cross-slip. From this limiting temperature up to 520° C, the creep of uranium involves two independent but jointly active mechanisms: these are the motion of screw dislocations by cross-slip and dislocation climb by a diffusion mechanism. Above 520° C and up to the α/β transition, creep of polycrystals is controlled by climb of dislocations out of their slip-planes, following a pile-up process in the case of primary creep and dipole annihilation in the case of secondary creep. For monocrystals, creep depends on the climb of edge dislocations in sub-boundaries which existed before creep, and their rearrangement in these sub-boundaries. In the β phase, creep of polycrystals is controlled by the diffusional climb of dislocations out of their slip-planes.