Introduction

The D″ region, which lies just above the core-mantle boundary (CMB), is of primary importance in geodynamics because it represents the thermal boundary layer, where heat is transferred by diffusion from the core to the convective mantle. Its viscosity plays a key role in controlling the amount of heat that can be extracted from the core, with significant implications regarding the dynamics of the mantle. The discovery of a phase transition from bridgmanite to a post-perovskite (ppv) phase at pressure (P) and temperature (T) conditions close to those of the CMB has opened new perspectives regarding the interpretation of the D″ structure and dynamics. This phase, which exhibits a very peculiar layered structure, has attracted much attention. Based on first-principles calculations, Ammann et al.1 showed that the diffusion of Mg2+ and Si4+ is extremely anisotropic in post-perovskite, with high diffusion rates along <100>. This finding has led to the idea that the D″ layer could be weaker than the overlying mantle due to high diffusion creep rates in ppv, although creep rates may be controlled by diffusion in the slowest direction2. However, this hypothesis is difficult to reconcile with the strong seismic anisotropy that is the signature of the D″ layer (as first observed in ScS3 and Sdiff4 phases) and is usually indicative of dislocation creep rather than diffusion creep.

Making the link between seismic anisotropy observations and putative crystal preferred orientations (CPO) in ppv is not straightforward. In horizontally propagating phases, the observation that horizontally polarized shear waves propagate faster than vertically polarized ones (VSH > VSV) agrees with models involving dominant slip in (010) or (001) in ppv5,6. If only dominant slip on (100) seems to be excluded by seismological observations so far, the respective role of (010)6,7 and (001) slip8,9 remains debatable. From the mineral physics point of view, the hypothesis of slip in (010), which seems intuitive given the layered structure of post-perovskite, is difficult to assess experimentally. Indeed, observation of experimentally produced CPO in diamond cell experiments has also led to conflicting results10,11, which may result from textures formed during phase transformation12,13. Additionally, beyond the geometry of crystal plasticity (from CPO), there is no way currently to obtain experimentally quantitative data of the rheology of silicate post-perovskite under relevant P, T (and strain rate!) conditions of the D″.

Multiscale numerical modelling represents an alternative that is currently able to describe plasticity by dislocations of high-pressure minerals. Recent applications to wadsleyite14, ringwoodite15, periclase16 and bridgmanite17 have shown their ability to reproduce laboratory experiment data. The main implication of these studies (both numerical and experimental) is that pressure has a strong effect on the lattice friction opposed to dislocation glide. In bridgmanite, Hirel et al.18 showed that lattice friction increases monotonically throughout the lower mantle to reach, in the lowermost mantle, values of approximately 15 GPa. In this context, the behaviour of post-perovskite appears remarkable. Shearing the Si octahedral layers (for instance, with the [100](001) slip system) yields comparable values of lattice friction to those of bridgmanite19. However, shearing the structure parallel to the structural layering is much easier: lattice friction opposed to the glide of [100] screw dislocations in the (010) plane is one order of magnitude lower than that in (001). This is the reason why we focus on this slip system in the present study. Lattice friction describes the anisotropic mechanical resistance of the material at 0 K only; it is still necessary to model the thermal activation of dislocation glide to determine how lattice friction evolves in the conditions of the D″ layer.

In this work, we present a theoretical study based on full atomistic modelling of plastic deformation, demonstrating easy glide of [100](010) dislocations in MgSiO3 post-perovskite (Mg-ppv) at finite temperature and strain rate of the lowermost mantle.

Results

Dislocation core structure and Peierls barrier

The core structure of screw dislocations with the Burgers vector [100] is computed at the atomistic level using density functional theory (DFT). The atomic configuration of a straight dislocation line is shown in Fig. 1a. The stable core configuration is centred between two neighbouring Mg atoms and mainly spread in {011}. In (010), equivalent stable core configurations are found every ½[001], i.e. separated by a distance a′ of 3.07 Å. These cores correspond therefore to alternative variants labelled (I) spread in (011) and (II) spread in , as displayed in Fig. 1b. Although DFT represents high accuracy calculations, we also employ a pairwise potential to compute the minimum energy path (MEP) between two stable cores associated with dislocation glide in (010) and to carefully investigate finite size effects (as described in the supplementary materials). The MEP between configurations (I) and (II) is computed using the nudged elastic band (NEB) algorithm20,21,22. The observed path reproduces the peculiar <011> zig-zag trajectory (Fig. 1b) reported in a previous atomic-scale study19. The maximum of the Peierls potential corresponds to the MEP dislocation image when the (011) glide trajectory switches to . This high energy configuration is associated with 39 meV/b and 77.7 meV/b energy barriers for pairwise potential and ab initio simulations, respectively (Fig. 2b). Based on these results, the Peierls stress σp can be estimated from the maximum slope of the Peierls potential. Thus, DFT calculations lead to σp = 2.1 GPa, whereas the empirical potential simulations predict a lower value of 1 GPa (Fig. 2b). This apparent discrepancy is likely related to the known drawback of the pairwise potential parameterization used in this study. Indeed, both simulation techniques provide identical dislocation core structures, but the empirical potential underestimates the elastic stiffness coefficients C55 and C66 by a factor of 2 and, consequently, the corresponding anisotropic shear modulus μ (173 GPa vs 324 GPa) of the post-perovskite23. Qualitatively, one finds therefore the ratio σp/μ ~5·10−3 regardless of the level of atomic description.

Figure 1
figure 1

(a) Atomic structure of stable [100](011) dislocation core extracted from the two-cation thick (011) atomic layer highlighted in yellow on the right panel. (b) MgSiO3 post-perovskite structure viewed along [100]. Location of the low energy dislocation lines (I) and (II) is specified by the “screw” symbols. MEP of [100] dislocation gliding in the (010) plane (deduced from NEB simulations) is indicated by the green line. On the both panels, Si atoms and octahedra, forming rigid {010} layers, are shown in blue; Mg atoms, forming weaker {010} layers in grey. Red atoms on the left panel correspond to O atoms.

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Figure 2
figure 2

(a) Lattice friction: schematic illustration of a straight dislocation line gliding over the Peierls potential, which materializes the lattice friction. (b) Peierls barrier Vp calculated for [100](010), with DFT (dashed line) and pairwise potential (solid line). Evolution of the energy barrier with applied stress, i.e., Vp − σa ba′, deduced from NEB calculations with pairwise potential, is provided in colour (see the text for details).

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Thermal activation and kink-pair mechanism

At finite temperature, the actual motion of a dislocation occurs through the nucleation and propagation of kink-pairs, i.e., a dislocation does not move as a straight line but partly bows out over the Peierls potential, as illustrated in Fig. 3a. Unlike common silicates, dislocation glide in (010) of Mg-ppv is associated with a relatively low energy barrier (in metals for example, a σp/μ ratio of 10−3 is usually expected). Under an applied stress σa, a straight dislocation will move upwards of the Peierls barrier to reach an equilibrium configuration from which a bulge can form. Energy minimization of the bowed configuration is obtained from a balance between the local forces on the dislocation, classically called the line tension, i.e., the force resulting from the Peierls barrier and the applied force. However, assuming that the fluctuation of the Peierls potential is small compared to the energy at rest leads to the standard line tension (LT) model for describing the kink-pair mechanism24,25,26. Within the LT formulation, a screw dislocation line can be represented as a 1D function y(x), which describes its position y in the glide plane at each x coordinate along the dislocation line (Fig. 3a). Then, the dislocation line enthalpy HLT can be estimated according to the following expression:

Figure 3
figure 3

(a) Thermal activation of dislocation glide: schematic illustration of the kink-pair mechanism, which allows the dislocation line to pass from one stable position to the next at finite temperature. (b) Kink-pair formation activation enthalpy HLT and its evolution with applied stress, computed with the data acquired from ab initio and semi-empirical simulations.

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where Γ is the line tension (representing the stiffness of the dislocation line) and Vp(y(x)) − σaby(x) corresponds to the so-called “substrate enthalpy” in the 1D-Frenkel-Kontorova model27,28.

Computing line tension of a dislocation

To link the LT model with atomic-scale simulations, we follow the work of Dezerald et al.29 and discretize the integral in Eq. (1) into n segments {Yn} of length b:

where the sum over n accounts for the periodic boundary condition along the dislocation line.

To compute the line tension Γ, the energy cost associated with a dislocation bow-out consistent with the first stage of kink-pair formation should be estimated. Bending a dislocation line requires breaking up the 1b translational symmetry of the simulation cell. Thus, the length of the supercell along the dislocation line is increased up to 2b. We consider a dislocation line that consists of two segments of length b: segment S1 remains in the Peierls valley (I), while the other segment S2 bows out towards the next valley (II), as shown in Fig. 4c. To compute this process at the atomic scale, the evolution of atomic displacements Δx along [100] during the dislocation glide from (I) to (II) is analysed, relying on MEP structural information obtained from NEB simulations. Dealing with a complex material, we mostly focus on the cation sublattice and allow anions to adapt to the local displacement of cations. Along the MEP, one Mg and four Si atoms exhibit the largest displacements Δx along the dislocation line among the cations in the crystal (Fig. 4a,b). The selected Mg atom is located directly between the two Peierls valleys and bears the maximum displacement amplitude (approximately 0.8 Å). The Δx amplitudes of the four Si atoms are two times smaller (Fig. 4b). Once the evolution of atomic positions along the MEP is defined, one can create the dislocation bow-out while applying the exact displacements Δx consistent with the MEP (at zero stress) to the selected atoms belonging to the segment S2 of a dislocation line (Fig. 4d). The corresponding cations of segment S1 are fixed to their regular positions in the Peierls valley, constraining the degrees of freedom along the dislocation line. The computed change in energy, , related to the gradual disposition of line segment S2 is shown in Fig. 4b as a function of the reaction coordinate along [001]. Fitting the curvature of quadratic function ΔELT provides the line tension Γ = 9.2 eV/Å for DFT simulations (vs Γ = 7.1 eV/Å for empirical potential). The anisotropic line tension Γel = 2.08 eV/Å, calculated within Stroh formalism30 using the set of elastic constant Cij of Mg-ppv31 computed with the same generalized gradient approximation (GGA) and pseudopotentials, is notably lower. Previous studies of bcc metals based on a similar simulation approach for computing LT at the atomic scale29,32 report comparable discrepancies between the LT values predicted atomistically and from elastic theory. Indeed, the latter does not account for the large effect of the dislocation core contribution, which leads to drastically underestimated Γel values.

Figure 4
figure 4

Line tension calculation scheme.

(a) The five cations, four Si and one Mg, that exhibit the biggest displacements along [100] during dislocation glide are labelled here. Si atoms are indicated by blue spheres, Mg by grey, the anion sublattice is left out for clarity, the dashed line indicates the MEP followed by the dislocation during its glide, and the unit cell is indicated by the red rectangle. (b) Atomic displacements along [100], deduced for the five selected cations from NEB simulations. The plot provides displacements for a screw dislocation with positive Burgers vector; in the case of negative Burgers vector, atomic disposition occurs in the inverse way. (c) Schematic illustration of dislocation line bending, mimicking the instant of kink-pair formation, in a 2b simulation cell. Glide plane (010) is highlighted in grey. (d) Local structure of the cation sublattice produced by the screw dislocation (I) in the (010) plane. To create a bow-out, degrees of freedom along [100] are restricted for the cations located in the pink area: in segment S1, the selected atoms are fixed to their regular positions in the Peierls valley (I), while the corresponding atoms of segment S2 are forced to follow up to 5% of the MEP along a′ towards the next Peierls valley (II), as indicated by the arrows. The associated increase in elastic energy ∆E is provided in (b).

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Kink-pair activation enthalpy

Once line tension Γ is computed at atomic scale and the Peierls barrier VP is known, the equilibrium kink-pair shape at a given stress and the corresponding critical kink-pair enthalpy HLT can be calculated using Eq. (1). To solve Eq. (1), we rely on a trial function y(x) that describes the equilibrium shape of a symmetric kink-pair based on a combination of hyperbolic tangents28:

where, as previously mentioned, a′ is the periodicity of the Peierls potential and α and m are variable parameters.

The saddle point on the HLT(α,m) energy landscape ultimately defines both the enthalpy HLT and the equilibrium kink-pair configuration y(x). As illustrated in the Supplementary materials, kink-pair configurations are characterized by extremely large widths in the range of 35–40b (approximately 100 Å) resulting from the very low Peierls barrier. With applied stress, this width changes moderately while the height of the kink-pair decreases rapidly as the straight part of the dislocation line moves upwards along the Peierls barrier. Note that such a wide kink shape justifies our choice of the LT model with respect to unreasonable direct atomistic computation of bowed configuration lines (which would require too many atoms). Figure 3b shows the computed enthalpy HLT as a function of applied stress σa. As one expects from dislocation theory33, the kink-pair enthalpy is maximum under zero stress, with a value of 2.69 eV, corresponding to twice the energy of a single kink Hk, and it vanishes when the applied stress is equal to the Peierls stress. The normalized kink-pair enthalpy 2Hk/μb3 is found to be approximately 5 · 10−2, confirming the relatively low lattice friction borne by the [100](010) slip system in post-perovskite.

Discussion

The dislocation core structure computed here confirms the spreading of the [100] screw core in {011}, as observed in previous semi-empirical simulations19. The extension of the core (with a half-width of approximately 1.8 Å) is found to be in reasonable agreement with the first results of the dislocation core determined using the Peierls-Nabarro model31. In Carrez et al.31, the classical PN model, based on first-principles calculations (with the same GGA approximation and pseudopotentials as in this study) of generalized stacking fault energy, showed that the [100] screw dislocation should be compact in {011}, with a half-width of 1 Å. Discrepancies in core size between the present results and those of the PN model can be largely attributed to the GSF calculations method. As shown in Goryaeva et al.19, the GSF computed by Carrez et al.31 involved atomic layers above the actual spreading layer of the core, delimited by neighbouring Mg rows. Consequently, the GSF energies used in the PN model were overestimated, leading to a narrower core. Nevertheless, note that despite a spreading in {011}, the easiest glide plane of [100] screw dislocations is (010). This is the result of a glide alternating between (011) and , defining a global macroscopic glide plane (010). This non-standard behaviour could be evidenced only by full atomistic calculations. Based on calculation of the maximum height of the Peierls potential, we find a Peierls stress of 2.1 GPa, which is almost forty times smaller than the value reported in Carrez et al.31. The reason for this discrepancy is that the PN calculation of σp did not rely on the right lattice periodicity a′. The lattice periodicity a′ is found here to be ½[001] without any ambiguities based on the computation of the exact core energy. Moreover, even considering the correct lattice periodicity for the Peierls potential, the PN model would certainly not be accurate in evaluating the Peierls potential for such a rearrangement of atoms around the line during the glide. Indeed, the dislocation glide process (through a zig-zag scheme) violates one of the intrinsic hypotheses of the PN model, which is that the Peierls potential is evaluated in the plane of dislocation core spreading.

In post-perovskite, the Peierls stress of [100](010) glide at 120 GPa is remarkably low. This result is consistent with the observation of [100](010) dislocations in several experimental studies of low-pressure CaIrO3 and CaPtO3 post-perovskite analogues34,35,36,37. Because Mg-ppv is unquenchable, experimental studies remain scarce10,11. From the experimental point of view, in CaIrO3 post-perovskite, the few TEM studies36 do not show strong evidence of lattice friction for [100] dislocations (as observed for olivine deformed at low temperature, where dislocation lines tend to be aligned along a particular direction). This can be supported by several results found here, i.e., the low normalized values of Peierls stress and kink-pairs enthalpy, and from the computed line tension. As recently demonstrated for perovskite material17, kink-pair enthalpy evolution as a function of stress can be used to infer the evolution of the critical shear stress for dislocation glide as a function of temperature:

In Equation (4), p and q describe the evolution of the kink-pair enthalpy through the following relationship: ΔHLT(σa) = 2Hk (1 − (σa/σp)p)q. Temperature Ta, often called “athermal temperature”, corresponds to the critical temperature at which lattice friction vanishes. Generally, kink-pair energy 2Hk scales with Ta according to 2Hk = CkTa, where k is the Boltzmann constant; C is a function of a strain rate ε′, of the dislocation density ρ, and of the kink geometry (for more details see the corresponding section of the Supplementary Materials). C is classically found in a range of 20–30 (this is verified in metals38 and also in oxides39). Therefore, taking 2Hk = 2.69 eV, p = 0.73 and q = 1.31 from the DFT calculations (considering experimental conditions, i.e., strain rates of 10−5 s−1 and dislocation density of 1012 m−2), we find that lattice friction vanishes if the temperature is raised above 1,100 K.

More importantly, as demonstrated recently for bridgmanite17, the previous equation can be used in Earth mantle conditions by adjusting the scaling factor of Ta to strain rates characteristic for convection in the Earth’s mantle. Assuming a typical value of 10−16 s−1, the corresponding temperature evolution for the critical stress for the glide of [100] dislocations is as shown in Fig. 5.

Figure 5
figure 5

Evolution of the critical resolved shear stress (CRSS) with temperature for [100](010) in Mg-ppv at 120 GPa, compared with that for ½[110](100) and ½<110>{110} in periclase16 at 100 GPa.

CRSS for post-perovskite is computed based on the data inferred from DFT simulations. Strain rate and dislocation density ρ correspond to the lower mantle conditions (see the text for details).

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Implications

Our results on dislocation glide in Mg-ppv, including certain unexpected results, shed new light on the rheology of high-pressure mantle phases. Indeed, all recent studies, either experimental or theoretical, of wadsleyite14,40,41,42, ringwoodite15,43,44, periclase16,45 and bridgmanite17,46 consistently show that pressure in the transition zone and lower mantle range leads to a significant increase of lattice friction, which inhibits dislocation glide as a strain-producing mechanism. In particular, these results have important implications regarding the (non-)formation of seismic anisotropy from the deformation of the above-mentioned phases. In this context, it is surprising to find that crystal chemistry and the formation of a layered structure can lead to a completely different behaviour. Our results demonstrate that the presence of weak {010} Mg-layers containing a very short <100> lattice repeat of 2.5 Å leads to dislocation structures that can easily glide. We find that lattice friction is overcome at a critical temperature Ta far below the temperatures expected in the D″ layer (3,700–4,400 K47,48). This finding has several unexpected consequences. The relative ease of slip between Mg-ppv and periclase suggests that the latter could become the stronger phase in the D″ layer. We are not yet in a position to fully establish this fact because additional deformation mechanisms must be activated in Mg-ppv to ensure compatibility of plastic deformation in an aggregate. However, the fact that diffusion is also fast in this phase1 suggests that complementary deformation mechanisms involving diffusion should be easily activated. Mg-ppv being the dominant phase in this assemblage, it is expected that the D″ layer in regions dominated by the Mg-ppv should exhibit a very low viscosity compared to the overlying mantle.

The implications of such a low viscosity layer have already been considered and discussed49,50,51. The way that slabs behave when ultimately reaching the CMB is clearly affected, as is the broad dynamics of the CMB. However, the strongest implication is probably the enhancement of heat transfer from the core across the CMB, as earlier predicted by Buffett49 and more recently investigated numerically50,51. The most testable implication of our results is, of course, the strong (010) crystal preferred orientation, which should develop upon flow in this weak layer. This is an important parameter because the D″ layer has long been recognized as being highly anisotropic. Although no consensus has yet been reached (see, for instance, Cottaar et al.9), our finding that Mg-ppv exhibits dominant easy glide along (010) is consistent with the most recent studies of Nowacki et al.6 and Ford and Long7.

In addition to a low viscosity, a low lattice friction in Mg-ppv may have important implications regarding seismic wave attenuation. A seismic (body) wave corresponds to strains in the range of 10−8–10−6, with periods in the range of 1–10 s. These values correspond to stresses of a fraction of a MPa at most, applied at a strain rate of 10−6 s−1 or lower. Under these conditions, the athermal temperature Ta will be greater than in Fig. 5, but, in any case, lower than 1,400 K (the value corresponding to a strain-rate of 10−5 s−1, constraining the dislocation density at 108 m−2; see supplementary Figure S5). This result shows that Mg-ppv will be in the athermal regime under seismic loading conditions at temperatures of the D″ layer, with dislocations moving freely without lattice friction.

This situation has not been considered up to now for seismic attenuation because most discussions have been driven by the example of olivine52,53. Olivine exhibits lattice friction; thus, dislocations are prescribed to stay in their Peierls valleys, and dislocation damping can only result in a limited contribution from kink migration52,53. For this reason, the most important source of attenuation in olivine has been linked to diffusionally assisted grain boundary sliding54.

Under an applied stress σa and without lattice friction, a dislocation segment of length l will bow out with a curvature . This bowing gives rise to a reduction of the effective shear modulus, called the modulus defect or relaxation strength: , where ρ is the dislocation density55. Assuming that l scales with , one can conclude that Δ0 could be as high as 15% (as an upper bound). Consequently, a shear wave travelling through the post-perovskite containing dislocations could encounter a maximum velocity reduction approximately 7% compared to the ideal structure. In this regime, dislocation damping can be described using the vibrating string model55 which assumes that under an applied alternating stress, a dislocation characterized by a line tension can execute forced vibrations like a vibrating string. An alternating stress, such as the one associated with a seismic wave, will result in damping and energy dissipation. This model has two important consequences. At sufficiently high frequency, there exists a peak in tanϕ versus ω with a resonance at ω0 such that . The frequency ω0 is a function of Γ, i.e., the line tension of the dislocation, as previously computed, and of ml, the effective mass per unit length of the dislocation line. This effective mass ml can be computed by summing the squared displacements dqi of all atoms i in a simulation cell in which a dislocation has moved from one Peierls valley to the next one, i.e., by dQ = a′, using the following expression56:

In the previous expression, Mi corresponds to the mass of atom i; the effective mass of the dislocation thus incorporates the kinetic energy of surrounding atoms as if they were to respond adiabatically to dislocation motion. Computed from different cell sizes to account for finite size effects, we find (in units of atomic mass per unit length) ml ~ 9.7 u/Å, which results in ω0 ~ 3.3·104 s−1. This frequency is higher than that of seismic waves, but it could allow experimental verification in the laboratory. At lower frequencies corresponding to seismic waves, the internal friction is proportional to ρl4 and to the frequency ω.

Our proposition of post-perovskite being highly attenuating is consistent with the report of higher attenuation in the D″ by Anderson and Hart57 and Lawrence and Wysession58. However, attenuation in the D″ layer is still not well constrained; this issue deserves more attention in the future to verify our prediction.

Conclusions

In this study, we modelled the thermally activated mobility of [100](010) slip in Mg-ppv using the line tension model in conjunction with atomic-scale simulations.

We show that under pressure, temperature and strain-rate conditions of the lowermost lower mantle, there is no lattice friction opposed to the glide of these dislocations. This easy glide of [100](010) dislocations has several implications:

- Although the exact viscosity of Mg-ppv cannot be calculated from a single slip system, we can predict a weak behaviour comparable with, if not weaker than, that of periclase.

- This conclusion supports the scenarios that involve a weak layer in the D″ layer with, in particular, enhancement of heat transfer from the core.

- Easy glide along (010) suggests development of marked crystal preferred orientations characterized by alignment of the (010) planes.

- The high mobility of dislocations allows for energy dissipation (vibrating string model) when a seismic wave travels through dislocation-bearing Mg-ppv. We predict that deforming Mg-ppv should be characterized by strong seismic attenuation.

Methods

In this work, we employ an effective combination of the first-principles simulations with pairwise potential modelling. All simulations are performed with an external pressure of 120 GPa. Unit cell parameters for a Mg-ppv perfect crystal (Cmcm, Z = 4) at relevant conditions are a1 = 2.474 Å, a2 = 8.112 Å, a3 = 6.139 Å for the first-principles simulations and a1 = 2.521 Å, a2 = 8.124 Å, a3 = 6.050 Å for the semi-empirical calculations.

Ab initio calculations are performed based on DFT within the GGA, as derived by Perdew and Wang59, and the all-electron projector augment-wave (PAW) method, as implemented in VASP code60,61. The outmost core radius for Mg, Si and O is 2.0, 1.9 and 1.52 au, respectively. To achieve computational convergence, we apply a plane-wave cut-off of 600 eV. The first Brillouin zone is sampled using the Monkhorst-Pack scheme62, with a 10 × 1 × 1 k-point grid for 1b simulation cells containing 360 atoms (the exact geometry is described below) and with a 4 × 1 × 1 grid for 2b cells containing 720 atoms. The convergence energy is 10−3 meV/atom.

Atomistic simulations within the semi-empirical approach are carried out using the Buckingham form of a pairwise potential, with the parameterization derived by Oganov et al.63 for MgSiO3 perovskite. Transferability of this parameterization has been previously validated for modelling ground state properties and defects in the post-perovskite phase23. Molecular statics simulations are performed using the program package LAMMPS64, which relies on Ewald summation methods for Coulombic interactions. Optimization of dislocation core configurations are performed using a conjugate-gradient (cg) algorithm, followed by a Hessian-free truncated Newton (hftn) algorithm, until the maximum force on an atom drops below 10−9 eV/Å (1.602·10−18 N). The NEB simulations are performed via fire damped dynamics, as required by the minimization procedure implemented in LAMMPS. The MEP is sampled with 24 points (configuration images), which are bounded with a spring constant of 0.1 eV/Å.

All simulations are performed by employing a quadrupole arrangement of screw dislocations in fully periodic atomic arrays. Such simulation cells contain two dislocations with positive and two dislocations with negative Burgers vectors arranged as a rectangular checkerboard pattern. This geometry allows for cancelling the long-range displacement field produced by a dislocation65 and ensures that interaction of the dislocations remains at a quadrupolar level and that the net force on each core is zero due to the periodic arrangement66. The supercell is designed in such a way that [100] dislocation lines are parallel to x and [010] and [001] crystallographic directions are aligned with z and y, respectively. For core energy calculations and evaluation of the Peierls potential, the designed atomic systems are as thin as a single Burgers vector b along x, i.e., dislocation lines are straight and infinite due to the periodic boundary conditions. Because computation of line tension Γ requires bowing out the dislocation line and breaking the translational symmetry along x, we increase the length of the supercell along the dislocation line and employ 2b geometry, following the strategy proposed by Rodney and Proville67. For DFT simulations, we employ the smallest possible atomic array of 36 Å × 48 Å, with quadrupolar arrangement of dislocations, which is further reduced by half to a non-rectangular (but still fully periodic) cell containing a dislocation dipole, following the procedure described by Bigger et al.66. By applying periodic boundary conditions to such a dipole, the rectangular checkerboard pattern arrangement of dislocations (identical to that in the original rectangular cell containing four dislocations) is explicitly reproduced. The reduced simulation cells employed in this work for DFT simulations contain 360 and 720 atoms for 1b and 2b geometry, respectively. For the simulations performed with the pairwise potential, the size of the atomic arrays, containing a quadrupole of <100> screw dislocations, is gradually increased along y and z to track the size effect on the computed substrate enthalpy VP and line tension Γ. A typical rectangular simulation cell has a size of 97 Å × 97 Å and 3840 and 7680 atoms in the case of 1b and 2b geometry, respectively.

Finally, for periodic arrangements of opposite Burgers vector dislocations, we compute the elastic interaction term, in accord with anisotropic elastic theory68, and subtract from the energy computed via the NEB method. The exact location Yc of each dislocation image is defined through analysing the relative atomic displacements near the dislocation cores, which are further used to compute the disregistry function S(Y) using the following expression69:

where b is the Burgers vector, Yc is the coordinate of a dislocation centre, and ζ is an adjustable parameter.

Additional Information

How to cite this article: Goryaeva, A. M. et al. Low viscosity and high attenuation in MgSiO3 post-perovskite inferred from atomic-scale calculations. Sci. Rep. 6, 34771; doi: 10.1038/srep34771 (2016).