#Plastic lattice crack#

The separation of elastic and plastic strain components in the context of computational modeling, whether by continuum ( 17) or atomistic ( 18) approaches, is more commonplace as these values are calculated for each element or from atom positions at every computation step during deformation. Furthermore, these measurements would be directly relatable to, and validate, model predictions of local elastoplasticity across the length scales, such as at the microstructural boundaries that often control the overall mechanical behavior.

#Plastic lattice crack#

At a crack tip, stress must accumulate to activate plastic flow for blunting and shielding by energy dissipating permanent deformation, without the mechanical potential energy becoming high enough to cause excessive crack propagation ( 16). This decomposition of strain would enable substantial broadening of our basic understanding of the elastic-plastic interplay, which is key to improving properties like the intrinsic toughness of materials. The ability to simultaneously track elastic, plastic, and total strains throughout deformation, with the necessary sub–100-nm resolution ( 15), may be considered the holy grail in characterizing the micromechanisms of materials mechanics.

The scale bar on the inset secondary electron image of the larger of the two cantilever dimensions used is 5 μm long the scale bar on the filtered (see Materials and Methods) secondary electron image of the Pt speckle pattern is 500 nm long. ( B) Notched microcantilever bending of single-crystal pure tungsten was performed in a SEM with in situ correlative DIC and HR-EBSD strain mapping on a region of interest surrounding the notch. The present study further extracts the purely plastic deformation through correlative measurement of the total and elastic deformation throughout mechanical loading. ( A) The total deformation gradient F measurable by digital image correlation (DIC) that maps an initial material configuration, B 0, onto a final, B, may be decomposed ( 73) into a plastic deformation F p (e.g., dislocation glide along crystal planes, indicated by inverted T symbol) to an intermediate configuration, B p, followed by elastorotational deformation F e (stretching atomic bonds and rotating the crystal lattice see text S1) the latter is measurable by HR-EBSD. However, not only is this technique laborious, requiring cryogenic conditions to limit spurious plasticity by atomic diffusion, but it is also not universally applicable: Amorphous materials, for example, do not deform by dislocation motion, nor do brittle ceramics above a certain size ( 6) or metals under diffusion creep. ( 5) in the context of discrete dislocation models. TEM in situ mechanical testing may be used to measure pure plastic strains by counting the passage of individual dislocations, as recently proposed by Cui et al. No current method generally allows for this. To answer these questions experimentally, we need to be able to simultaneously measure, in a sample under load, both the local elastic strains (from which the stress state is deduced) and the pure plastic strains (to determine to what extent deformation mechanisms are activated, i.e., excluding the residual elastic contribution to common net elongation measurements of “plastic strain”) with sub–100-nm spatial resolution to resolve deformation features.

As a result, mechanisms are often broken down into excessively simple relationships, which lose sight of the real complexity involved ( 3) and, over half a century later, remain hotly debated topics ( 4). Physical models exist to describe these phenomena ( 1, 2) however, these are based on a relatively small number of experimental observations as transmission electron microscopy (TEM) is time-consuming and limited to small volumes, as well as to the unavoidable plane stress state of thin foil specimens.