5.2 Edge dislocation-void interaction
5.2.2 Discussion
It has been reported in previous studies that dislocation breakaway is always accompanied by vacancy absorption, i.e. by positive climb of the dislocation, and the creation of a pair of superjogs. This observation was due to the fact that the previous research was based on voids centred on the glide plane, but, as already mentioned, climb by vacancy absorption was not always found for the configurations studied here. Dislocation climb did occur in all cases, but, depending on void position relative to the glide plane, it could be negative as well as positive: the dislocation did not always absorb vacancies from the void but it sometimes left extra vacancies behind in the void. The latter happened in configurations -R/2 and -R, where the void equatorial plane is below the glide plane. As seen in figures 5.17 and 5.18, this effect is more obvious for big voids.
It seems that vacancy exchange between the void and the dislocation is not the cause for dislocation climb but rather the effect of it. As the positive edge dislocation enters the void, it ‘pushes’ its surface atoms forward above the glide plane; this leads to the creation of the shear step at the end of the interaction. This shearing would be expected to happen on the plane where the dislocation
‘slices’ the void, i.e. the glide plane. This is the situation on the entry surface where the dislocation is pure edge in character and cannot cross-slip. In order to minimise the dislocation energy, the dislocation ‘prefers’ to terminate at the void for as long as possible. Therefore, when it reaches the exit surface on the glide plane, instead of leaving the void, one or both of the screw arms cross- slip upwards or downwards (depending on the geometry of the configuration) towards the void equator and therefore prolong its termination at the void as long as possible. Subsequently, the shear step on the exit surface is located above or below the dislocation glide plane; hence the surprising result that climb occurs even for configurations R and -R.
Of course, climb is not necessarily enough for the dislocation to manage to reach the equator. When the applied stress is high enough, the dislocation may break away earlier. This is also assisted by the self-stress effect, as predicted by the Scattergood-Bacon model.
In the case of configuration 0, where the glide plane coincides with the void equator, no climb would be expected. However, this is the case where the highest vacancy absorption is observed (∼40 vacancies for 4 nm voids). This arises from two factors. The first is that, as mentioned in the introduction of this chapter, the void equator was not located exactly on the slip plane, but directly below it, initiating the climb process. The second is the fact that the voids are not exactly spherical, but are faceted. Therefore, as the dislocation meets specific facet steps, it ‘chooses’ to cross-slip on them to minimise the energy of the required surface step. This results in the form the resultant superjogs obtain, with segments lying in specific directions. If random vacancy absorption occurred, this would not be the case.
This conclusion on the origin of the climb mechanism is also enhanced by observation of the superjogs shape in figure 5.4 (a), reproduced from Osetskyet al. [108]. It can be seen that for big voids the superjogs created are a result of both positive and negative climb. The fact that the total number of vacancies absorbed is larger than those left behind is again due to the exact position of the void centre. This is shown in fig. 5.4 (b) for the case of a 3 nm void, where it can be seen that there are eight (1¯10) atomic planes above the dislocation glide plane, and seven below it. In other words, the void centre is located exactly above the glide plane, and, as explained above, this results in positive mean climb.
Figure 5.4: (a) Dislocation line viewed in [111] projection perpendicular to b after intersecting voids of different sizes. The void diameters and numbers of vacancies removed are indicated. (b) Position of atoms in six consecutive (¯1¯12) planes through the centre of a 3 nm void after dislocation breakaway. The horizontal line indicates the trace of the dislocation glide plane and the arrow indicates the exit step on the void surface. Owing to dislocation climb, the exit step is two (1¯10) planes above the glide plane. Both figures are reproduced from Osetsky et al. [108].
The fact that, with increasing void size, the climb presents more clearly both positive and negative components (fig. 5.4 (a)), indicates that for voids of even bigger diameter, positive and negative climb might be of equal magnitude and no vacancy exchange would occur between the void and the dislocation. This would make sense, considering that as the void size increases, the asymmetry created by the imperfect coincidence between the equator and the glide plane becomes less important.
The vacancy exchange effect is significantly stronger for the big, 4 nm, voids. This might result in a slight shift of their stress-strain profiles, resulting in the different strength order for the two void sizes.
Finally, with respect to the conclusion in Osetsky et al. [108] that after breakaway void size is reduced and the void becomes a weaker obstacle, it would seem that dislocation glide within a crystal can also provide a mechanism for void growth. In principle, according to the mechanism described above, a moving dislocation that would react with a number of voids of all configurations, would reduce the vacancy content of the part of the crystal above its glide plane, and increase the vacancy content underneath its glide plane.