Stacking-fault tetrahedra in deformed face-centred cubic metals

Loretto, M. H.; Clarebrough, L. M.; Segall, R. L.

doi:10.1080/14786436508224233

Cited by 109 publications

(26 citation statements)

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“…A general rule of thumb for metals is that 15% of plastic work AE plastic :•'SA'y s (A'y s is the shear increment in the slip system s) is stored energy [Loretto et al, 1964[Loretto et al, , 1965Hirth and Lothe, 1982]. The latter criterion, however, is incompatible with steady state flow, which corresponds to increasing plastic work without an increase in the dislocation density.…”

Section: Growthmentioning

confidence: 99%

Modeling dynamic recrystallization of olivine aggregates deformed in simple shear

Wenk¹,

Tomé²

1999

J. Geophys. Res.

101

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Abstract. Experiments by Zhang and Karato [1995] have shown that in simple shear dislocation creep of olivine at low strains, an asymmetric texture develops with a [100] maximum rotated away from the shear direction against the sense of shear. At large strain where recrystallization is pervasive, the texture pattern is symmetrical, and [100] is parallel to the shear direction. The deformation texture can be adequately modeled with a viscoplastic self-consistent polycrystal plasticity theory. This model can be expanded to include recrystallization, treating the process as a balance of boundary migration (growth of relatively underformed grains at the expense of highly deformed grains) and nucleation (strain-free nuclei replacing highly deformed grains). If nucleation dominates over growth, the model predicts a change from the asymmetric to the symmetric texture as recrystallization proceeds and stabilization in the "easy slip" orientation for the dominant (010)[100] slip system. This result is in accordance with the experiments and suggests that the most highly deformed orientation components dominate the recrystallization texture. The empirical model will be useful to simulate more adequately the development of anisotropy in the mantle where olivine is largely recrystallized. IntroductionSimple shear experiments of rock-forming minerals and analogs deformed by dislocation creep reveal some puzzling features. After moderate strains, the deformed original grains develop a characteristically asymmetric preferred orientation pattern with a monoclinic symmetry, including a mirror plane perpendicular to the shear plane and containing the shear direction. These experimental simple shear deformation textures (in this paper the term texture is used synonymous with lattice preferred orientation), such as those for quartz [ These orthorhombic textures cannot be predicted with polycrystal plasticity theory for deformation by dislocation glide such as Taylor, Sachs, or self-consistent models. Authors have referred to the textures as "easy slip" orientations [e.g., Schmid and Casey, 1986], implying that a microscopic slip plane of an active slip system is parallel to the macroscopic shear plane and a microscopic slip direction is parallel to the shear direction, and therefore deformation on such a slip system in simple shear is "easy." So far, there has been no explanation to how crystals rotate and remain within those orientations. Herwegh and Handy [1996] suggest a "rigid body rotation" into the easy slip orientation and an unexplained cessation of rotation once those special orientations are reached.Experimental

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Section: Growthmentioning

confidence: 99%

Modeling dynamic recrystallization of olivine aggregates deformed in simple shear

Wenk¹,

Tomé²

1999

J. Geophys. Res.

101

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“…Among these defects, stacking fault tetrahedra (SFT), bounded by stacking faults (SFs) and stair-rod dislocations, are peculiar types of volume defect typically found in face-centred cubic (FCC) crystals with low stacking fault energy 1,3 . Generally, SFT can be introduced by three methods, that is, quenching, high-energy particle irradiation and plastic deformation [3][4][5][6][7][8] . In 1959, Silcox and Hirsch 4 first observed SFT in quenched gold (Au).…”

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Atomic-scale dynamic process of deformation-induced stacking fault tetrahedra in gold nanocrystals

et al. 2013

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Stacking fault tetrahedra, the three-dimensional crystalline defects bounded by stacking faults and stair-rod dislocations, are often observed in quenched or irradiated face-centred cubic metals and alloys. All of the stacking fault tetrahedra experimentally observed to date are believed to originate from vacancies. Here we report, in contrast to the classical vacancyoriginated ones, a new kind of stacking fault tetrahedra formed via the interaction and crossslip of partial dislocations in gold nanocrystals. The complete atomic-scale processes of nucleation, migration and annihilation of the dislocation-originated stacking fault tetrahedra are revealed by in situ high-resolution observations and molecular dynamics simulations. The dislocation-originated stacking fault tetrahedra can undergo migration and annihilation due to mechanical loading in a manner that is not expected in bulk samples. These results uncover a unique deformation mechanism via dislocation interaction inside the confined volume of nanocrystals and have important implications regarding the size effect on the mechanical behaviour of small-volume materials.

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“…Haussermann and Wilkens 1966; Steeds 1967), and triangular Frank dislocation loops and stacking fault tetrahedra (e.g. Silcox and Hirsch 1959;Loretto, Clarebrough, and Segall 1965). The main advantages of the third method over the other two are that it is applicable to materials of a very wide range of stacking fault energy and involves only simple length measurements of defects that are easily recognized.…”

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confidence: 99%

“…The present authors have reconsidered this problem and, subject to the limitations of isotropic linear elasticity, have taken into account the major variables that may affect the values of y. It is the purpose of this note to present the results of this theory in a form in which values of y may easily be obtained from measurements of Frank dislocation loops and stacking fault tetrahedra without the resources of a large digital computer.Loretto, Clarebrough, and Segall (1965) have shown that triangular Frank dislocation loops and stacking fault tetrahedra are formed in f.c.c. metals and alloys when these are plastically deformed.…”

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The Determination of Stacking Fault Energy by the Tetrahedron Method

Humble¹,

Forwood²

1968

Aust. J. Phys.

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At present there are three methods for obtaining values of the stacking fault energy y of face-centred cubic (f.c.c.) materials by direct observation of dislocationstacking fault configurations in the electron microscope. These are based on measurements of extended three-fold dislocation nodes (e.g. Whelan 1958;Brown and ThOlen 1964), faulted dipole configurations (e.g. Haussermann and Wilkens 1966; Steeds 1967), and triangular Frank dislocation loops and stacking fault tetrahedra (e.g. Silcox and Hirsch 1959;Loretto, Clarebrough, and Segall 1965). The main advantages of the third method over the other two are that it is applicable to materials of a very wide range of stacking fault energy and involves only simple length measurements of defects that are easily recognized. However, it has suffered from the disadvantage that the values of y deduced from these measurements relied on an incomplete theory. The present authors have reconsidered this problem and, subject to the limitations of isotropic linear elasticity, have taken into account the major variables that may affect the values of y. It is the purpose of this note to present the results of this theory in a form in which values of y may easily be obtained from measurements of Frank dislocation loops and stacking fault tetrahedra without the resources of a large digital computer.Loretto, Clarebrough, and Segall (1965) have shown that triangular Frank dislocation loops and stacking fault tetrahedra are formed in f.c.c. metals and alloys when these are plastically deformed. Their experiments indicated that a dislocation mechanism was responsible for the formation of the triangular Frank loops and that all the stacking fault tetrahedra were formed by the dissociation of these in the manner originally suggested by Silcox and Hirsch (1959). Thus, by observing the size of the largest tetrahedron and the smallest Frank loop in a given plastically deformed material, it is possible to determine the critical edge length lc above which the transformation of loops to tetrahedra is energetically unfavourable.Several authors (Czjzek, Seeger, and Mader 1962; Jossang and Hirth 1966;Humble, Segall, and Head 1967) have computed the energy balance between triangular Frank loops and stacking fault tetrahedra as a function of defect size, but they have considered only the terms in the dislocation interaction energy and stacking fault energy. We have recently reformulated the problem considering the total energy of the defect (Humble and Forwood 1968). This formulation includes terms such as the kinetic energy of the moving Shockley dislocations T, the energy dissipated to the crystal lattice fL, and the work done by the stress W, as well as terms in the interaction energy Etd and fault energy E ty • The potential energy of a crystal containing the defect Etd + Ety -W was plotted as a function of dissociation, * Manuscript

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Stacking-fault tetrahedra in deformed face-centred cubic metals

Cited by 109 publications

References 10 publications

Modeling dynamic recrystallization of olivine aggregates deformed in simple shear

Modeling dynamic recrystallization of olivine aggregates deformed in simple shear

Atomic-scale dynamic process of deformation-induced stacking fault tetrahedra in gold nanocrystals

The Determination of Stacking Fault Energy by the Tetrahedron Method

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