Research Group for Nuclear Materials Modeling
Location: JAEA > Nucl. Sci. and Eng. Direct. > Fuels and Mater. Eng. > Res. Group for Nucl. Mater. Modeling
(A) Interaction between grain boundary and dislocation
The interaction between dislocations and grain boundaries is the principal factor for determining the mechanical properties and the plastic deformation behavior of metals. It is possible to control the grain-boundary microstructure and the macroscopic behavior has been widely exploited for scientific and industrial applications. In atomic scale, however, specific interaction characteristics such as the reaction energy and pathway have yet to be revealed. We have investigated the interaction process between a dislocation and an energetically stable grain boundary, and the quantitative characteristics were determined via atomistic transition state analysis. As a result, the interaction energy is found to be 0.116 eV/A, which is much higher than the Peierls potential. The lattice dislocations subsequently experience anomalous dissociations on the grain boundary (Fig. A-1), which becomes a key factor for the previously unexplained dislocation disappearance and grain-boundary migration.
An incipient plastic deformation of several types of grain boundaries subjected to nanoindentation was investigated by atomistic simulations. Crystal defects such as grain boundaries undermine the nucleation resistance. In this paper, we examined the dislocation nucleation mechanism at the twin and several coincidence site lattice grain boundaries and the resulting weakening of the dislocation nucleation resistance. We found that for the twin and the relatively stable ƒ°11 grain boundary the primary slip deformation is activated on the grain boundary plane prior to the defect-free region because of the low fault energy of the grain boundaries during slip deformation. Subsequently, the secondary slip is activated from the grain boundary (Fig. A-2).
Fig. A-1 Interaction process of an edge dislocation and a twin boundary. (a) The reaction energies of four layers along the MEP of the interaction process per unit length of a dislocation. (b) Atomic configurations visualized by atomic shear stress.
Fig. A-2 Atomic images for the initial yield event , where atoms are visualized according to the centrosymmetry parameter. (a) The initial and final states corresponding to before and after the load drop are shown in (A) and (D), respectively. And the transition states between the initial and final states are shown by (B) and (C). (b) Crystallographic orientation of the displacement site lattice dislocation on the twin boundary and the slip dislocation.
(B) Hardening effects and unfaulting of radiation defects
Anomalous lattice defects such stacking fault tetrahedron (SFT) and self interstitial atom (SIA) loop are induced by irradiation, which contribute to the deformation properties. In this study the structural development under deformation is investigated by molecular dynamics simulations. As a result, it is found that the presence of SFT and the SIA loop cause the increase of yield stress as they act as the obstacle to the dislocation glide (Fig. B-1). An edge dislocation absorbs SIA loop during cutting SIA while a screw dislocation effects a change in the slip system of SIA (Fig. B-2).
The dislocation channeling observed in irradiated metals has been thought to be one of the key stress factors in irradiation assisted stress corrosion cracking (IASCC), in which the dislocation accumulation occurs and the stress concentration are generated around the grain boundaries. In the present study, the deformation characteristics of irradiated metals is investigated by atomistic simulations. The polycrystal models are prepared by Voronoi polyhedron division, and the interstitial and vacancy type defect structures are introduced as the radiation defects. Deformation modes of a whole system of a sequence of polycrystalline models are examined. Grain boundaries are preferentially deformed and that dislocations are intensively emitted from the heterogeneous triple grain junctions at the initial stage of plastic deformation. As a result, radiation defects on the slip systems near those of glide dislocations are annihilated during the slip motion of a number of dislocations (Fig. B-3).
Fig. B-1 Shear stress-strain responses of the interaction between radiation defect and edge/screw dislocation.
Fig. B-2 Atomic images of the dislocation configurations during the interaction between radiation defect and edge/screw dislocation.
Fig. B-3 Unfaulting of radiation defects under tensile deformation in polycrystalline model.
(C) Non-empirical prediction of impurity segregation in a-Fe from first principles
The segregation of impurities in a-Fe were investigated by first principle density functional theory calculations. The segregation tendencies of various elements observed in reactor pressure vessels were considered and the interaction characteristics between Fe and each impurity element were estimated by mean field approximation. Stable N-atom impurity clusters were subsequently chosen to evaluate the changes in free energy for clustering. These calculations show that Cu and Mn impurities embedded in a-Fe are more stable when they are in the segregated state. Conversely, Nb and Ta are stable in the separately solute state (Fig. C-1). The present estimates provide reliable suggestions for the segregation characteristics. We suggest that the segregation tendency is derived from the d-orbital interaction (Fig. C-2) and that the solubility limit is not necessarily correlated with the tendency of clustering formation.
Fig. C-1 Enthalpy changes from the isolated impurity to the segregated N-atom clusters
Fig. C-2 Local DOS of the typical segregated and nonsegregated impurities of (a) Cu and (b) Nb embedded in a-Fe, where the local DOS of the nearest neighbor Fe atom and those of single-component Cu and Nb at bcc phase are drawn.
T. Tsuru, et al., J. Appl. Phys. 107 (2010), 061805.
(D) Phase transformation of copper precipitation and its effect on obstacle strength in a-iron
The size- and spacing- dependent obstacle strength due to the Cu precipitation in a-Fe is investigated by atomistic simulations, in which the effect on phase transformation of Cu precipitation is considered by a conventional self-guided molecular dynamics (SGMD) method that has an advantage to enhance the conformational sampling efficiency in MD simulations. A sequence of molecular statics simulations of the interaction between a pure edge dislocation and spherical Cu precipitation are performed to investigate the obstacle strength associated with phase transformation. It was shown that the SGMD method can accelerate calculating the bcc to 9R structure transformation of a small precipitate, enabling the transformation without introducing any excess vacancies. Such metallographic structures increase the obstacle strength (Fig. D-1, D-2) through strong pinning effects as a result of the complicated atomic rearrangement within the Cu precipitation.
Fig. D-1 Shear stress-strain responses of a-Fe including D2 to D6 precipitation.
Fig. D-2 Atomic images of the dislocation configurations during the interaction between an edge dislocation and (a) coherent and (b) transformed Cu precipitations.
(E) Interactin between lattice defects and He cluster in a-Fe
Helium has a strong tendency to precipitate into thermally stable helium-vacancy clusters and helium bubbles, which causes embrittlement of ferritic steel in fusion reactor material. In this study the thermodynamic stability of diffusion of He in a-Fe are investigated by atomistic simulations based on both first-principles DFT and empirical potential. As the result the migration enthalpy on T-T migration in the  direction is one order smaller than that of T-O migration and therefore the T-T migration is dominant diffusion path of He in Fe (Fig. E-1). Additionally the stability of He atom cluster up to six atom is investigated by molecular dynamics simulation with 3-body potential. A cluster with several He atom becomes stable at the vacant lattice point while a single He is stable at the T-site of bcc metals.
Fig. E-1 Change in enthalpy during T-T and T-O-T migration fo He in a-Fe.