Modeling Twin Clustering and Strain Localization in Hexagonal Close-Packed Metals

Timár, G.; Fonseca, João Quinta da

doi:10.1007/s11661-014-2547-z

Cited by 11 publications

(10 citation statements)

References 34 publications

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“…In principle, this finding is in good agreement with [30], which predicts the formation of twin clusters and the role of <a> slip strain location in Mg. In contrast, easy shear transfer from {10-12}<-1011> twins was not identified as important factor within the chain/clusters suggesting that twin nucleation by stress concentration from a neighbouring grain that has twinned is not a main driving force for twin formation.…”

Section: Discussionsupporting

confidence: 82%

“…For this reason, the prismatic <a> slip transfer analysis is of importance here. Recent crystal plasticity prediction have indeed suggested for Mg that strain localisation, in this case due to basal <a> slip, does result in intergranular stress concentration in those grains, which promote twin formation and clustering [30]. Figure 7a clearly displays a shift to higher values of the m'-value distribution, computed from within twin clusters/chains compared to the m'-values of the non-twinned grain family in respect of its neighbourhood, Figure 7a.…”

Section: Discussionmentioning

confidence: 80%

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3D characterisation of early twin formation in Ti‐4Al by diffraction contrast tomography

Nervo¹,

Ludwig²,

Fonseca³

et al. 2016

Proceedings of the 13th World Conference on Titanium

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Twin formation in alpha-Ti is a topic of much debate as the role of microstructure features and alloy chemistry on twin formation remains unclear. In the present work, early stage of twin formation, i.e. after 1% plastic compression, has been studied in 3D using diffraction contrast tomography. This technique has enabled to study a volume about 400 grains enabling statistical analysis. The family of twinned grains was compared with similarly orientated grains that had not twinned in respect of their neighbourhood and grain size. While the initial neighbourhood analysis did not identify any significant differences, 3D grain size analysis highlighted that grain size plays an important role during the early stage of twinning. Further, twin clusters/chains were identified, which form slightly imperfect 3D networks. A more detailed neighbourhood analysis of the parent grains associated with the twin clusters/chains demonstrate that strain localization in neighbouring grains well aligned for prismatic slip transfer promote the observed twin clustering.

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Section: Discussionsupporting

confidence: 82%

Section: Discussionmentioning

confidence: 80%

3D characterisation of early twin formation in Ti‐4Al by diffraction contrast tomography

Nervo¹,

Ludwig²,

Fonseca³

et al. 2016

Proceedings of the 13th World Conference on Titanium

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“…There is an increased difficulty in predicting the deformation and activity of different slip systems in HCP alloys compared to cubic alloys [52,95]. This is due the greater number of slip systems that can be active and the variations in their ease of activity.…”

Section: Plasticity Approachmentioning

confidence: 99%

Peak Broadening Anisotropy and the Contrast Factor in Metal Alloys

Simm

2018

Crystals

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Diffraction peak profile analysis (DPPA) is a valuable method to understand the microstructure and defects present in a crystalline material. Peak broadening anisotropy, where broadening of a diffraction peak doesn't change smoothly with 2θ or d-spacing, is an important aspect of these methods. There are numerous approaches to take to deal with this anisotropy in metal alloys, which can be used to gain information about the dislocation types present in a sample and the amount of planar faults. However, there are problems in determining which method to use and the potential errors that can result. This is particularly the case for hexagonal close packed (HCP) alloys. There is though a distinct advantage of broadening anisotropy in that it provides a unique and potentially valuable way to develop crystal plasticity and work-hardening models. In this work we use several practical examples of the use of DPPA to highlight the issues of broadening anisotropy. of 31where, D is the crystal size, K Sch is the Scherrer constant (often taken as 0.9 for spherical crystals), g is the reciprocal of the d-spacing of a peak, f m is a function related to the arrangement of dislocations (and represents how the arrangement of groups of dislocations influence their total strain), ρ is the dislocation density and C hkl is the average contrast factor of dislocations in grains that contribute to the hkl diffraction peak. In Equation 1 and other DPPA approaches, C hkl is assumed to be the only term that accounts for broadening anisotropy and its value is required to be able to obtain the dislocation density.TEM to be able to identify details of dislocations [15,16], especially in cases when g.b=0 does not lead to vanishing contrast or due to practicalities of rotating a sample. In the same manner, the diffraction broadening caused by a dislocation varies depending on the diffraction vector, and can be calculated by modelling the dislocation's displacement field. The contribution of a dislocation's displacement field to the broadening of different diffraction profiles is through what is known as the contrast (or orientation) factor of dislocations. The contrast factor of an individual dislocation is dependent on the angles between the diffraction vector and the vectors that define the dislocation: it's Burgers vector (b), slip plane normal (n), and dislocation line (s). The contrast factor of an individual dislocation can be calculated using the computer program ANIZC [17]. For example, the edge dislocation in Figure 2 has a contrast factor of 0.46 when g is [110] and parallel to the dislocations Burgers vector. The value is 0.00 when g is parallel to the slip line ([11 2]) and 0.05 when g is parallel to the slip plane normal ([1 11]).

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“…This can be seen in Figures 5a and 5b, where there seem to be microcracks that form along some grain boundaries, with small amounts of melt perhaps draining down and away. In this study we did not attempt to quantify Journal of Geophysical Research: Solid Earth 10.1002/2017JB014881 the microcracking, but strain localization and crack initiation in hcp metals have been modeled numerically by Timar and Fonseca (2014), at low temperatures. Moreover, anecdotally, the untextured, equiaxed sample rings machined at room temperature more easily than did the directionally solidified ones, which tended to chip.…”

Section: 1002/2017jb014881mentioning

confidence: 99%

“…After large-angle torsional deformation, the directionally solidified sample rings also showed considerably less integrity than the untextured, equiaxed ones. In this study we did not attempt to quantify Journal of Geophysical Research: Solid Earth 10.1002/2017JB014881 the microcracking, but strain localization and crack initiation in hcp metals have been modeled numerically by Timar and Fonseca (2014), at low temperatures. Obviously, under the high pressures in the inner core cracking is unlikely due to volumetric changes, but because of the unfavorable conditions for plastic flow of large, textured grains, one might expect lower ductility and more heterogeneous deformation (Karato, 2008), as we have observed.…”

Section: 1002/2017jb014881mentioning

confidence: 99%

Grain Boundary Sliding in High‐Temperature Deformation of Directionally Solidified hcp Zn Alloys and Implications for the Deformation Mechanism of Earth's Inner Core

Bergman

Lewis

et al. 2018

JGR Solid Earth

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Earth's inner core exhibits elastic anisotropy, but the cause of the texturing is uncertain: solidification, deformation, or some combination. In experiments presented here we deformed in torsion directionally solidified Zn‐3 wt pct Sn alloys at 0.97 of the melting temperature of pure Zn. We used a Zn‐rich alloy because Zn is hexagonal close‐packed at atmospheric pressure, as is likely Fe under inner core conditions. The directionally solidified alloys have the textured, columnar dendritic microstructure that has been proposed for the inner core. As a control we deformed in torsion the same alloy at the same temperature, but with a starting microstructure that is comparatively fine‐grained, equiaxed, and untextured. We find evidence that the directionally solidified alloys deform by means of grain boundary sliding, presumably accompanied by diffusion near the ends of the columnar crystals in order to preserve macroscopic shape. We also find that the directionally solidified alloys require approximately 3 times the torque to deform to the same strain than do the untextured alloys. Grain boundary sliding and greater hardness are due to insufficient slip systems to accommodate the strain, as a result of the texturing. This is supported by microcracking in the deformed directionally solidified alloys. Although the inner core is unlikely to exhibit semibrittle deformation, the experiments suggest that a directionally solidified inner core could have a higher viscosity than one with a finer‐grained, untextured microstructure and that grain boundary sliding may be a dominant deformation mechanism, in spite of the large grain size.

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Modeling Twin Clustering and Strain Localization in Hexagonal Close-Packed Metals

Cited by 11 publications

References 34 publications

3D characterisation of early twin formation in Ti‐4Al by diffraction contrast tomography

3D characterisation of early twin formation in Ti‐4Al by diffraction contrast tomography

Peak Broadening Anisotropy and the Contrast Factor in Metal Alloys

Grain Boundary Sliding in High‐Temperature Deformation of Directionally Solidified hcp Zn Alloys and Implications for the Deformation Mechanism of Earth's Inner Core

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