Nucleation and propagation of dislocations in nano-Crystalline FCC metals
The normal perception regarding GBs (grain boundaries) is that they are found to have been playing the role of both, sinks and sources, for the dislocations in a condition in which there is a reduction in grain size to a nanometer regime in face-centred cubic (FCC) metals. Molecular dynamics (MD) computer simulations (Van Swygenhoven et al., 2001) is the basis of this mechanism, in which GB emits partial dislocation, which is absorbed in opposite and neighboring areas after travelling through the whole grain. It has been ascertained that in few materials (like A1) leading partial's emission is often followed with trailing partials. On the contrary, in other materials (like Cu and Ni) the entire grain witnesses SF (stacking fault) defect being transected rather than tailing partial. On the basis of absolute value of stable SF energy, a model has been put forward so that different dislocation characters evident in the MD can be explained and the association between splitting distance with the critical grain size for the emission of the trailing can be recognized (Van Swygenhoven et al., 2002). All the simulation results cannot be explained by this approach, which although is very attractive. The results published can help us in understanding that trailing and leading partial dislocations are only evident in the simulation of A1. Trailing partials were not found for Ni and CU (with higher and lower stable SF energy values, respectively, in regard to the potentials used). It is clear that the occurrence of partial or full dislocations cannot be verified by using absolute value of splitting distance and SF on the basis of nucleation criterion (Froseth et al., 2006).
It has been examined that, for the understanding of MD results, an individual must take into account both SF energies, i.e. stable as well as unstable state with the help of a generalized SF energy curve (Derlet et al., 2003a; 2003b). This approach enables us to understand the simple explanations towards all MD results, when the leading partial has been nucleated after which the ratio between unstable and stable SF energies is near to unite the energy barrier of the following partial which is relatively smaller for the material. It would also be explained as whenever the ratio is close to unity, the nucleation of following partial is expected to be in the timescale of MD simulation. This only happens in Al, not in Cu and Ni, which explains the reason behind the full dislocations as seen in the MD simulation of Al. So we cannot over look the fact that the observation related to the extended pile of faults in Cu as well as Ni could be the reason of the of MD simulation timescale. Furthermore, as long as the experiment goes the dislocation activity will nucleate both leading as well as trailing partial, ultimately full dislocation will happen. This is so as there in no experimental proof of increasing density of SFs after tensile distortion in nanocystalline (nc) Ni and Cu. With the help of electron microscopy examination it had been seen that there are isolated large SFs (Van Swygenhoven et al., 2003). Through lab examination of X-ray in nc-Ni, we have come to know that the act of reversibility of diffraction comes to peak during the process of plastic distortion. It demonstrates that there is lack of permanent residual diffraction network and helps us in supporting the ideology that nc-GBs will absorb leading as well as trailing partials that are predominantly emitted (Van Swygenhoven, 2002; Froseth et al., 2006).
There are numerous other very unusual and easily noticed details that are being observed through atomistic simulation other then issues related to trailing and leading partials (Froseth et al., 2006):
(1) After emission an extreme hydrostatic pressure is relived before nucleation (Yamakov et al., 2001; 2003);
(2) During emission, atomic shuffles and free volume migration, due to high stress, also occurs (Froseth el al, 2006).
(3) The distance that shows separation between trailing and leading partials is determined by both stress distribution of GB and the path on which dislocation travels (Derlet et al., 2003b; Forseth et al., 2006).
Many analytical models are generated in response to the results of the above mentioned observation. But in all the models the common elements are that they all carried out a small size dependent nucleation criterion having the ultimate goal to follow the pattern that were observed in Hall-Perch behavior (Yamakov et al., 2004), 10 times increased strain rate sensitivity (Froseth et al., 2004) and low value measured for the activation volume (10-20b) (Van Swygenhoven et al., 2004; Froseth et al., 2006).
Some significant facts related to propagation as well as nucleation are reveled through the above mentioned details related to dislocation in nc fcc metal which are of great importance to individual associated with mesoscale modeling (Froseth et al., 2006):
(1) Emitting of leading as well as trailing partials are not at same projection, not even at same GB, which propose that GB dislocation nucleation can't be a representative of the simple Frank-Read source (Froseth et al., 2006).
(2) Both nucleations as well as propagation are two different processes. It may happen that partial is nucleated but it is not propagated. This statement is valid in case of both leading and trailing partials (Froseth et al., 2006).
(3) GB ledges can certainly distract the dislocation propagation which is subject to the geometrical circumstances of the ledge along with the burgers vector. Dislocations do have the tendency to formulate into pinned on certain point of actions causing huge mess and in an amplified dislocation curvature (Froseth et al., 2006).
(4) According to Froseth et al. (2006) from the pinning point when dislocation goes through the deposition onto the ledge which takes place from part of the vector, making a new strain concentration, advising that dislocation is not an exhaustive procedure (Froseth et al., 2006).
(5) The thermally activated method relating to the consumption of time for the occurrence of unpinning is highly reliant, which is being provided with the evidence (Froseth et al., 2006).
While (1) because of the GB ledge removal upon nucleation a dislocation resource in the GB works only once and (2) leading and irregular fractional dislocations are mostly nucleated at distinctive areas. Regions that are wholly found through the structure of GB and mainly on the presence of GB ledges, represent a dislocation source which is not being conformed to the conventional FRANK-READ dislocation nucleation method as advised by researchers (Chen et al., 2003a; 2003b). According to Forseth et al. (2006) it is being concluded that while observing the previous simulations with the exact resolution as taken place currently the trailing partial is never transmitted as accurately on the same region as the leading partial, which is also being suggested by practitioners (Froseth et al., 2006).
Distance between the ledges could be referred to as the only length scale that could be aligned for the method of nucleation, however; it also cannot be verified fully because the leading and trailing partial dislocations nucleation took place at different GBs of distinctive disorientations. Due to the non-consideration, GBs structure depending on the grain size effect depicted a complex picture. In the former course of work (Van Swygenhoven et al., 2000), it has been presented to us that the structures of GB are basically not distinct in comparison with NC structures and coarse-grained size scale, where observations, of same misfit structures for identical GB disorientations and GB plane orientations, are being done (Van Swygenhoven et al., 2000).
It is duly expected that along with a larger number of high - energy GB's which were not present in coarse grain samples, various disorientations distributions would be implemented during the synthesis of metals along with grain sizes at the nano meter scale. Including a huge quantity of ledges which are close to the TJs or TJ lines, more GB disorientations will be examined. By forming the emission of dislocations from ledges along with their consequent elimination easier, not only shuffling but also stress-assisted free volume migration is supported by the occurrence of neighboring TJ/lines (Derlet et al., 2003; Van Swygenhovenet al, 2004). Pure shu-ing apart from dislocation emission (Van Swygenhoven et al., 2001) is possible by the removal of the ledges to make possible intergranular processes such as GB sliding which could be done by the exploration, which would help making an image towards the smallest grain sizes (Froseth et al., 2006).
After the emission as the evidences start to disappear, the pragmatic dislocation mechanism seems to be exhaustive after the first display. We also see that it is expected that the dislocation debris might eventually end in the development of latest sources whereas during the propagation this is deposited on GB's and due to the limited timescale proposed by the MD simulations, such an experiment has never been reported (Froseth et al., 2006).
In the propagation of a dislocation where either not to act or to…
"Nucleation And Propagation Of Dislocations In Nanocrystalline FCC Metals" (2011, April 18) Retrieved June 26, 2017, from http://www.paperdue.com/essay/nucleation-and-propagation-of-dislocations-119786
"Nucleation And Propagation Of Dislocations In Nanocrystalline FCC Metals" 18 April 2011. Web.26 June. 2017. < http://www.paperdue.com/essay/nucleation-and-propagation-of-dislocations-119786>
"Nucleation And Propagation Of Dislocations In Nanocrystalline FCC Metals", 18 April 2011, Accessed.26 June. 2017, http://www.paperdue.com/essay/nucleation-and-propagation-of-dislocations-119786