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Grain growth in polycrystals is traditionally considered a capillarity-driven process, where grain boundaries (GBs) migrate toward their centers of curvature (i.e., mean curvature flow) with a velocity proportional to the local curvature (including extensions to account for anisotropic GB energy and mobility). Experimental and simulation evidence shows that this simplistic view is untrue. We demonstrate that the failure of the classical mean curvature flow description of grain growth mainly originates from the shear deformation naturally coupled with GB motion (i.e., shear coupling). Our findings are built on large-scale microstructure evolution simulations incorporating the fundamental (crystallography-respecting) microscopic mechanism of GB migration. The nature of the deviations from curvature flow revealed in our simulations is consistent with observations in recent experimental studies on different materials. This work also demonstrates how to incorporate the mechanical effects that are essential to the accurate prediction of microstructure evolution. 

Abstract This study uses high‐energy X‐ray diffraction microscopy of SrTiO₃to identify correlations between grain boundary (GB) area changes and the motion direction of neighboring GBs to investigate interfacial energy minimization mechanisms during grain growth. The local GB area changes were measured near triple lines (TLs) to isolate the effects of neighboring GBs. These area changes were then correlated to the migration direction and curvature of the neighboring GBs present at the TL, providing an alternative metric associated with lateral expansion for describing GB migration. Additionally, this study extracted GB dihedral angles, which reflect the relative GB energy, to test whether low energy GBs replace high energy GBs (i.e., GB replacement mechanism) and, thus, can be used to predict a GB's migration direction. The majority of GBs did not exhibit local area changes reflective of the GB replacement mechanism, and the dihedral angles were not reliable indicators of GB motion. However, the expansion and shrinkage of GBs moving away from their center of curvature was more often consistent with the grain boundary replacement mechanism. These results suggest that growth for certain GB configurations is governed by relative energy differences while others are governed by curvature. 

Abstract Determining suitable dopants with optimized doping concentration is critical to design efficient water splitting photocatalysts. However, there is currently a lack of fundamental knowledge to guide this process. Herein, we examine the impact of Al³⁺, Mg²⁺, and Ga³⁺on the photocatalytic performance of SrTiO₃and propose a defect compensation model to understand the doping effect. Doped SrTiO₃crystals were grown hydrothermally and treated in molten SrCl₂. The hydrogen production rates from 50 catalysts produced in this way were measured with a high‐throughput parallelized and automated photochemical reactor (PAPCR). The investigation revealed that all three dopants significantly enhance the photocatalytic reactivity. According to Brouwer diagrams computed using available reaction constants, the optimum reactivity is achieved when the concentration of acceptor dopants fully compensates the oxygen vacancy donors. The improved reactivity can be attributed to the reduction in free electron concentration, resulting in a space charge layer that is 1000 times longer. Consequently, this situation enhances the number of photogenerated charge carriers capable of being separated by the band bending and transported to the surface. 

Grain boundaries in polycrystalline materials migrate to reduce the total excess energy. It has recently been found that the factors governing migration rates of boundaries in bicrystals are insufficient to explain boundary migration in polycrystals. We first review our current understanding of the atomistic mechanisms of grain boundary migration based on simulations and high-resolution transmission electron microscopy observations. We then review our current understanding at the continuum scale based on simulations and observations using high-energy diffraction microscopy. We conclude that detailed comparisons of experimental observations with atomistic simulations of migration in polycrystals (rather than bicrystals) are required to better understand the mechanisms of grain boundary migration, that the driving force for grain boundary migration in polycrystals must include factors other than curvature, and that current simulations of grain growth are insufficient for reproducing experimental observations, possibly because of an inadequate representation of the driving force.