Small-Size-Induced Plasticity and Dislocation ... - ACS Publications

Subscriber access provided by UNIV OF CAMBRIDGE

Communication

Small Size-induced Plasticity and Dislocation Activities on Non-charge-balanced Slip System of Ionic MgO Pillars Ting-Chun Lin, Chao-Chun Yen, Shao-Yi Lin, Yi-Chung Huang, Chi-Huan Tung, Yu-Ting Hsiao, and Shou-Yi Chang Nano Lett., Just Accepted Manuscript • Publication Date (Web): 09 Jul 2018 Downloaded from http://pubs.acs.org on July 9, 2018

Just Accepted “Just Accepted” manuscripts have been peer-reviewed and accepted for publication. They are posted online prior to technical editing, formatting for publication and author proofing. The American Chemical Society provides “Just Accepted” as a service to the research community to expedite the dissemination of scientific material as soon as possible after acceptance. “Just Accepted” manuscripts appear in full in PDF format accompanied by an HTML abstract. “Just Accepted” manuscripts have been fully peer reviewed, but should not be considered the official version of record. They are citable by the Digital Object Identifier (DOI®). “Just Accepted” is an optional service offered to authors. Therefore, the “Just Accepted” Web site may not include all articles that will be published in the journal. After a manuscript is technically edited and formatted, it will be removed from the “Just Accepted” Web site and published as an ASAP article. Note that technical editing may introduce minor changes to the manuscript text and/or graphics which could affect content, and all legal disclaimers and ethical guidelines that apply to the journal pertain. ACS cannot be held responsible for errors or consequences arising from the use of information contained in these “Just Accepted” manuscripts.

is published by the American Chemical Society. 1155 Sixteenth Street N.W., Washington, DC 20036 Published by American Chemical Society. Copyright © American Chemical Society. However, no copyright claim is made to original U.S. Government works, or works produced by employees of any Commonwealth realm Crown government in the course of their duties.

Page 1 of 29 1 2 3 4 5 6 7 8 9 10 11 12 13 14 15 16 17 18 19 20 21 22 23 24 25 26 27 28 29 30 31 32 33 34 35 36 37 38 39 40 41 42 43 44 45 46 47 48 49 50 51 52 53 54 55 56 57 58 59 60

Nano Letters

Small Size-induced Plasticity and Dislocation Activities on Noncharge-balanced Slip System of Ionic MgO Pillars

Ting-Chun Lin1, Chao-Chun Yen2, Shao-Yi Lin3, Yi-Chung Huang1, Chi-Huan Tung2, Yu-Ting Hsiao2, and Shou-Yi Chang2,*

1

Department of Materials Science and Engineering, National Chung Hsing University, Taichung 40227, Taiwan

2

Department of Materials Science and Engineering, National Tsing Hua University, Hsinchu 30013, Taiwan

3

Department of Mechanical and Computer-Aided Engineering, National Formosa University, Yunlin 63201, Taiwan

ABSTRACT. We observed small size-induced hardening and plasticity of brittle ionic MgO owing to abnormally triggered dislocation gliding on a non-charge-balanced slip system. The indentation tests of 〈111〉 MgO pillars revealed an increased hardness with decreasing pillar size, and the tips of the pillars ≤ 200 nm were plastically deformed. The in-situ compression tests of 〈111〉 MgO nanopillars in TEM verified aligned dislocation-mediated plasticity on the {111}〈110〉 or {100}〈110〉 rather than the charge-balanced {110}〈110〉 slip system.

KEYWORDS. Ionic nanopillar; size effect; plasticity; dislocation activity; slip system

1 Environment ACS Paragon Plus

Nano Letters 1 2 3 4 5 6 7 8 9 10 11 12 13 14 15 16 17 18 19 20 21 22 23 24 25 26 27 28 29 30 31 32 33 34 35 36 37 38 39 40 41 42 43 44 45 46 47 48 49 50 51 52 53 54 55 56 57 58 59 60

Ionic crystals consist of cations and anions of high binding energy (600−1500 kJ/mol, 3−8 eV/atom) and accordingly have a high stiffness and hardness [1]. However, when ionic materials are subject to plastic deformation, the relative positions of cations and anions to retain an electroneutrality (charge balance) need to be considered in determining dislocation slip systems. For face-centered cubic (fcc) ionic crystals, charge-balanced (non-close packed) {110}〈110〉, rather than non-charge-balanced {111}〈110〉 (close-packed) or {100}〈110〉, has been conventionally recognized as the main slip system [1-4]. The instability of the chargebalanced, non-close packed {110}〈110〉 slip system that results from a reduced number of slip system, a doubled length of Burgers vector and hence the increased activation energy for dislocation motion causes the detrimental brittleness and low plasticity (< 0.6%) of bulk ionic materials [3, 5-6]. Only at a high temperature and/or under an extremely high pressure, the transition of slip system into non-charge-balanced {100}〈110〉 would be activated achieve a high plasticity of > 2% for ionic materials [7, 8]. Different mechanical behaviors of nanosized materials from those of bulks have been found, in consequence of large specific surface area and unique surface bond reconstruction [9-12]. Extraordinary dislocation activities at the nanoscale, such as mechanical annealing, dislocation starvation and the easy dissociation of perfect dislocations into partials, have been observed [13-16]. Particularly, a small size-induced plasticity has been discovered in covalent fcc GaAs pillars when the pillar size reaches the nanoscale [17]. The intersection of two slip bands of different orientations causes the early fracture of microsized GaAs pillars, whereas single-directional slip band-accommodated deformation improves both the strength and the ductility of nanosized pillars [17]. Attributable to a stress shielding effect (the alignment of dislocations) to reduce system energy [18, 19], the density of accommodated dislocations in nanosized pillars can increase although the strain hardening rate will decrease. Originallybrittle covalent materials with a size of below 100 nm can accordingly become tougher [18,


Page 2 of 29

Page 3 of 29 1 2 3 4 5 6 7 8 9 10 11 12 13 14 15 16 17 18 19 20 21 22 23 24 25 26 27 28 29 30 31 32 33 34 35 36 37 38 39 40 41 42 43 44 45 46 47 48 49 50 51 52 53 54 55 56 57 58 59 60

Nano Letters

19]. However, no clear mechanism for such hardenability and plasticity of nanosized ionic materials, at room temperature without a high pressure, has been demonstrated. The nanoscale mechanical properties and deformation behavior (including plasticity) of brittle ionic materials (mainly oxides and nitrides) are of great importance as they dominate the mechanical reliability of nano-optoelectronic devices such as flexible displays. Although nanomechanical testing particularly in-situ has been intensively used in the past decade for characterizing the mechanical behavior of nanosized metallic materials [20-25], however, few in-situ studies were conducted for nanosized ionic materials [26-28]. How the mechanical response of nanosized ionic materials will differ from that of bulk samples, and whether other slip systems than conventional charge-balanced {110}〈110〉 will be activated at the nanoscale, are still of interest. Hence in this study, the hardness (strength) and plastic deformation of single-crystalline fcc ionic MgO pillars were investigated. The indentation tests of micro- to nanosized MgO pillars (Figure 1a) were carried out to investigate the small-size effect on strength and plasticity, and the in-situ compression tests of nanopillars (Figure 1b) were also conducted in a transmission electron microscope (TEM) to examine dislocation activities (see Materials and Methods). Particularly, 〈111〉 oriented pillars, for which the resolved shear stress on the {110}〈110〉 slip system was zero to inhibit dislocation gliding on a conventional charge-balanced system, were tested to determine the possibility of dislocation activation on the non-charge-balanced {111}〈110〉 and {100}〈110〉 slip systems of ionic materials.

Small-size effect on the hardness and plasticity of MgO pillars First, the elastic modulus of the single-crystalline 〈111〉 MgO wafer was measured, by nanoindentation, to be 252 GPa, very close to the value given in the literature, 250 GPa [29], which would verify the accuracy of the measurements. Two unique mechanical behaviors were observed in the present study when the diameter of MgO pillars reached the nanoscale


Nano Letters 1 2 3 4 5 6 7 8 9 10 11 12 13 14 15 16 17 18 19 20 21 22 23 24 25 26 27 28 29 30 31 32 33 34 35 36 37 38 39 40 41 42 43 44 45 46 47 48 49 50 51 52 53 54 55 56 57 58 59 60

(below 200 nm), including the hardening and the ductile deformation of the pillars. Figure 2 (a) first plots the indentation load–depth (P−h) curves of micro- to nanosized MgO pillars (20 dynamic mechanical analysis (DMA) data points at different depths for one measured curve, and five measured curves for the average fitting curve of each sample size), and Figure 2 (b) plots the contact area–depth function of indenter tip to a blanket MgO wafer ( Ac ≈ 24.56 h 2 , provided by the instrument and carefully calibrated using fused silica). From the applied load, P, and the indenter contact area, Ac, the hardness of the samples can be obtained simply as H = P / Ac . Apparently, the same range of loading curves for all the MgO pillars of different diameters suggests the same level of hardness rather than any size effect on the hardness of the pillars. However, the real loading areas of the pillars (in particular of the nanosized pillars) would differ from (be smaller than) the contact area of indenter tip to the blanket wafer, Ac , and accordingly the true hardness of the pillars are believed to change (increase) with decreasing pillar size. The contact areas of indenter tip to (i.e. the loading areas of) the differently-sized pillars were hence calibrated, referring to the King model [30, 31], as Ad = Ac I = Ac [1 − exp(− αL / a )] where L is the pillar length (~ 2.5×diameter d), a is the square root of projected contact area (~ 5h), and α is a function of indenter geometry and a/L (~ 0.9 for a triangle indenter tip at a/L ~ 1), as also presented in Figure 2 (b). To check the accuracy of the calibrated contact areas using the King model [30, 31], the geometrical contact models for large and small pillars have also been built (see Materials and Methods), and the geometrical contact areas have been calculated. The geometrical contact areas accord with the estimations using the King model, suggesting the applicability of the King model for estimating the loading areas of the pillars. Figure 3 (a) further plots the ratios of indenter contact area (given by the instrument) to calibrated contact areas (by the King model), M = Ac / Ad = 1 / I , for the micro- to nanosized


Page 4 of 29

Page 5 of 29 1 2 3 4 5 6 7 8 9 10 11 12 13 14 15 16 17 18 19 20 21 22 23 24 25 26 27 28 29 30 31 32 33 34 35 36 37 38 39 40 41 42 43 44 45 46 47 48 49 50 51 52 53 54 55 56 57 58 59 60

Nano Letters

MgO pillars. It is clear that the calibrated contact areas, Ad, for the large-size pillars (diameter ≥ 800 nm) are about the same as the indenter contact area, Ac, whereas the contact areas for the small-size pillars (diameter ≤ 200 nm) are obviously smaller than the indenter contact area (but still larger than the cross-sectional barrel areas of compressed pillars, referred to

πd 2 L = πD 2 (L − h ) where D is the average barrel diameter). Figure 3 (b) thereafter gives the more correct hardness of the micro- to nanosized MgO pillars (at different indentation depths; contact areas calibrated by the King model). Clearly, the hardness of the large-size pillars (diameter ≥ 800 nm) remains constant, at 13−14 GPa (the same as that of MgO wafer), but, when the size of the pillars reached the nanoscale (diameter ≤ 400 nm and in particular ≤ 200 nm), a small size-induced hardening effect is obvious. Although the measured hardness might change with the geometrical imperfections of the pillars (tapper angles in the pillar diameter and rounding at the pillar/substrate transition) [32] (with an aspect ratio of about 2.5 and a tapper angle of about 5°, about 10% higher than that of more geometrically perfect pillars), however, all the different-size pillars have similar geometrical imperfections, and the 10% inaccuracy is much lower than the small size-caused hardening discovered above. The other unique mechanical behavior interestingly discovered is the small size-induced plasticity, as presented in the SEM images of compressed micro- to nanosized MgO pillars in Figure 4. The microsized pillars (diameter ≥ 800 nm and in particular = 3200 nm) exhibited the typical brittleness of bulk ionic MgO, with clear cracks at the edges of pillar tips. With decreasing pillar size, by contrast, obvious plastic deformation was noticed. Especially for the nanosized pillars (diameter = 200 and 100 nm), the ductile deformation of the pillar tips, with the clear plastic flow of material rather than brittle cracking, was observed. The upper size limit for the ductile deformation might be around 200−400 nm as no clear plastic flow was observed at the tip of the pillar with a diameter of 400 nm. The unique small size-caused hardening and plasticity of the originally-brittle MgO may be attributed to several possible


Nano Letters 1 2 3 4 5 6 7 8 9 10 11 12 13 14 15 16 17 18 19 20 21 22 23 24 25 26 27 28 29 30 31 32 33 34 35 36 37 38 39 40 41 42 43 44 45 46 47 48 49 50 51 52 53 54 55 56 57 58 59 60

mechanisms that include limited dislocation sources [12], dislocation starvation [12], stress shielding (dislocation alignment) [18, 19], surface bond reconstruction [10] and/or even the change of dislocation slip systems. The more possible mechanism will be investigated below, nanoscopically, through the in-situ TEM observations of compressive deformation of singlecrystalline 〈111〉 MgO nanopillars.

Mechanical response and dislocation activities at the nanoscale Figure 5 presents the mechanical response of single-crystalline 〈111〉 MgO nanopillars under in-situ compression tests in TEM, and Figure 6 shows the TEM lattice images of the nanopillars before and after compression. The stress−strain curves (Figure 5 (a)) indicate an early elastic response of the nanopillars with a modulus of about 250 GPa at strain ≤ 5×10−3. With compression for a small strain ~ 2×10−2, the sharp tip of the nanopillars was obviously blunted (Figure 6); however, for a too large strain ≥ 4−5×10−2, the pillars would fracture. The stress−strain curves (Figure 5 (a)) and the in-situ observations of corresponding dislocation activities (Figures 5 (b) (c) and Supplemental Video 1) also suggest that the nanopillars yielded at a stress around 1.2−1.7 GPa, during which dislocations nucleated at the pillar tip and moved upwards. At the early stage of plastic deformation, the dislocation activities were not obvious, at a density of about 1010 cm−2, which caused the steep ascending of stress (near the elastic response) to 3−5 GPa. The intermittent formation, motion and starvation (sinking to the pillar surface) of dislocations resulted in the serrations of the stress−strain curves. At the late stage of plastic deformation, more dislocations at a higher density of about 1012 cm−2 were formed into aligned (parallel-arranged) clusters and moved upwards along one specific direction. The continual formation and movement of dislocation clusters in only one direction clearly contributed to the plastic strain. However, at the very late stage, dislocations began to move along two directions; the intersection of them from the different directions led to the


Page 6 of 29

Page 7 of 29 1 2 3 4 5 6 7 8 9 10 11 12 13 14 15 16 17 18 19 20 21 22 23 24 25 26 27 28 29 30 31 32 33 34 35 36 37 38 39 40 41 42 43 44 45 46 47 48 49 50 51 52 53 54 55 56 57 58 59 60

Nano Letters

initiation of cracks and the ultimate fracture of the nanopillars [33]. Additionally, in the TEM lattice images of the single-crystalline MgO nanopillar before the in-situ TEM compression test (Figures 1 (b) and 6 (a)), an almost perfect 〈111〉 MgO lattice was observed at the pillar tip. After compression (Figures 6 (b) (c)), the blunting of the pillar tip with obvious plastic flow of compressed material (piling-up of squeezed MgO lattice at the edges, similar to that in Figures 4 (e) (f)) was clearly noted. An amorphous surface layer with a thickness of about 4−5 nm was also noticed after compression (after long-time exposure to the electron beam). However, before compression, the amorphous layer showed a very small thickness below 1−2 nm because the pillars were milled at an ultralow FIB current to avoid the damage caused by ion bombardments. The influence of the ultrathin FIB-affected zone on the deformation of the pillars was thus neglected [34, 35]. To clearly understand the plastic deformation behavior and dislocation slip systems of similar brittle fcc 〈111〉 structure, single-crystalline fcc 〈111〉 covalent GaAs nanopillars (no charge-balance consideration) with an elastic modulus of about 100 GPa were also in-situ tested (Figure 5 (a) and Supplemental Video 2) for comparison. The GaAs nanopillars were more ductile, and similar (but more intensive) dislocation activities were observed. Although the strain-to-fracture of the MgO nanopillars, around 4−5×10−2, is smaller than that of the GaAs nanopillars and much smaller than those of typical metallic materials, it is still larger than the extremely low plastic strain of bulk ionic materials, which is possibly attributed to the viable dislocation activities at the nanoscale, as aforementioned. However, as it was designed for compression along the 〈111〉 direction of the 〈111〉 MgO nanopillars, dislocation gliding on the conventional charge-balanced {110}〈110〉 slip system should be inhibited since the resolved shear stress was zero (the detailed examination of slip systems and resolved shear stress will be addressed below). The observed dislocation gliding revealed the activation of other slip systems in MgO at the nanoscale. Figure 7 summarizes


Nano Letters 1 2 3 4 5 6 7 8 9 10 11 12 13 14 15 16 17 18 19 20 21 22 23 24 25 26 27 28 29 30 31 32 33 34 35 36 37 38 39 40 41 42 43 44 45 46 47 48 49 50 51 52 53 54 55 56 57 58 59 60

Page 8 of 29

the moving directions and gliding speeds of dislocations in the ionic MgO and the covalent GaAs nanopillars, measured in in-situ TEM compression tests. From the statistic plot of the moving directions of dislocations (Figure 7 (a)), indeed, nearly no dislocation moved along 0−5° (parallel to the {111} compression plane). Alternatively, most dislocation clusters in the ionic MgO moved at 85−90° (along the compression axis), while a few moved at 50−55° (inclined to the compression plane), exactly the same as the moving directions in the covalent GaAs. This observation suggests the same slip system of the MgO nanopillars as that of the GaAs nanopillars, that is the {111}〈110〉 for perfect dislocations (or the {111}〈112〉 for partial dislocations [17]) rather than the conventional charge-balanced {110}〈110〉. The fitting curves of dislocation gliding speeds, ν (Figure 7 (b)), obtained by using the

[(

Weibull accumulation function, F (v ) = 1 − exp − v / v

) ] where γ

v is the average speed and

γ is the curve shape parameter, indicate the difference of dislocation mobility in differentlybonded materials. For ionic MgO, v is determined to be 4.6 nm/s, γ is 1.6, and the average thickness of dislocation clusters is 31 nm, while for the covalent GaAs, v is 11.8 nm/s, γ is 2.0, and the thickness of dislocation clusters is about 20 nm. The lower mobility, the broader speed distribution and the more intense (thicker) clustering of dislocations all reveal a higher energy barrier for driving dislocation gliding in the ionic MgO than that in the covalent GaAs with the same level of binding energy, possibly also owing to the awkward atom motion on the non-charge-balanced {111}〈110〉 (or {111}〈112〉) slip system determined below.

Possible non-charge-balanced slip systems at the nanoscale Figures 8 (a) (b) show the fast Fourier transformed (FFT) TEM lattice images of slightly deformed 〈111〉 MgO nanopillar tips after in-situ TEM compression (for a small extent: 3 µN for 10 s). The dislocations (circled) are observed to fall on two specific orientations, 85−90° of majority and 50−55°. To determine the exact slip planes and slip vectors (Burgers vectors) 8 Environment ACS Paragon Plus

Page 9 of 29 1 2 3 4 5 6 7 8 9 10 11 12 13 14 15 16 17 18 19 20 21 22 23 24 25 26 27 28 29 30 31 32 33 34 35 36 37 38 39 40 41 42 43 44 45 46 47 48 49 50 51 52 53 54 55 56 57 58 59 60

Nano Letters

of these dislocations, the zone axis (the incident direction of electron beam) was identified first, by referring the experimentally detected SAD patterns (ex. Figure 1 (b)) to the standard pattern (Figure 8 (c)), to be 〈 112 〉 (i.e. 〈 211 〉). That is, the projections of the slip vectors on the { 211 } plane are aligned along 85−90° and 50−55°. Afterwards, the Thompson Tetrahedron [36] and corresponding cubic cell/coordinate system (Figure 8 (d)) are adopted to illustrate possible slip systems that include three series of slip planes, {110} (e.g. ACδ), {111} (e.g. ACD) and {100} (e.g. the light blue triangle), and one series of slip vectors, 〈110〉 (e.g. CA, AD, CD) [36]. The { 211 } projections of these possible slip systems (Figure 8 (e); {110}〈110〉, {111}〈110〉 and {100}〈110〉) are compared with the experimental observations of dislocation motion directions (Figures 5 (b) (c) and 8 (a) (b); TEM zone axis 〈 211 〉, i.e. { 211 } plane view) to determine the exact slip systems of the ionic MgO nanopillars. Table 1 listed the possible slip systems and the resolved shear stress on each system,

τ = ( f n / An ) cos θ cos φ = σ n cos θ cos φ [37] where the normal stress, σ n , is the normal force divided by the cross-sectional area of the nanopillars, f n / An , θ is the angle between the force and the normal direction of the slip plane, and φ is the angle between the force and the slip vector. It is accordingly realized that, as designed, for compression along the 〈111〉 direction of the 〈111〉 MgO nanopillars, the resolved shear stress on the typical {110}〈110〉 slip system was zero, and hence dislocation motion on this slip system should be inhibited. Dislocation motion was only viable on the {111}〈110〉 (ABD, ACD, BCD planes and AD, BD, CD directions) and {100}〈110〉 slip systems (three DEF-equivalent planes and AD, BD, CD directions) which received a resolved shear stress of 0.28σn and 0.49σn, respectively. As illustrated in Figure 8 (e), the { 211 } projections of the valid 〈110〉 slip vectors (AD to A’D’, and BD to BD’ or CD to CD’; red arrow lines) would fall at 90° and 51.4° to the projection line of the compression plane (ABC to BC), which is in good consistence with experimental 9 Environment ACS Paragon Plus

Nano Letters 1 2 3 4 5 6 7 8 9 10 11 12 13 14 15 16 17 18 19 20 21 22 23 24 25 26 27 28 29 30 31 32 33 34 35 36 37 38 39 40 41 42 43 44 45 46 47 48 49 50 51 52 53 54 55 56 57 58 59 60

observations, 85−90° and 50−55° (Figure 7 (a)). This finding concludes that the non-charge balanced {111}〈110〉 or {100}〈110〉, rather than the charge-balanced {110}〈110〉, are the most possible slip systems for the observed dislocation activities herein. The high-resolution TEM (HRTEM) 2D lattice image (with marked dislocations) of the slightly deformed 〈111〉 MgO nanopillar tip and the atomistic model of a defective 〈111〉 MgO crystal (with a b = 1/2 〈110〉 dislocation on a {111}〈110〉 slip system) are further presented in Figures 9 (a) (b) to more carefully determine the slip system. The atomistic model was built using the Atomsk and visualized using the Ovito software. With comparison of the 2D lattice image and the atomistic model, clearly, an identical structure/geometry of dislocations is found, which verifies the above observed dislocation gliding on the non-charge balanced slip system in the 〈111〉 MgO nanopillar, as schematically illustrated in Figure 9 (c). However, in a normal condition at room temperature without a high pressure, the energy required to activate dislocation gliding on the non-charge balanced {111}〈110〉 or {100}〈110〉 slip system is very high [7-8, 38], which should limit the plastic deformation but cause the brittleness of the ionic MgO pillars. Indeed, at the microscale, cracks were formed in the compressed microsized pillars (Figure 4 (a)) that released the applied stress and diminished work hardening. Only when the pillar size reached the nanoscale, dislocation activities were stimulated, and small size-induced plasticity was present (Figure 4 (e) (f)). A recent modeling study revealed that, for single-crystalline MgO under a very high hydrostatic pressure, the activation energy for dislocation movement on different slip systems would change [38], while the Bond Order-Length-Strength theory suggested that, on nanosized materials, the reconstruction of shortened and strengthened surface bonds would generate an extreme hydrostatic stress [10]. The stress criterions for partial dislocations on the close-packed {111} slip plane to move along 1/6 〈112〉, and for perfect dislocations on the {111} or {100} slip plane to move along 1/2 〈110〉, might become lower than that for perfect dislocations on the


Page 10 of 29

Page 11 of 29 1 2 3 4 5 6 7 8 9 10 11 12 13 14 15 16 17 18 19 20 21 22 23 24 25 26 27 28 29 30 31 32 33 34 35 36 37 38 39 40 41 42 43 44 45 46 47 48 49 50 51 52 53 54 55 56 57 58 59 60

Nano Letters

non-close-packed {110} plane to move along 1/2 〈110〉 [38]. Especially, when the chargebalanced {110}〈110〉 slip system was prohibited in the present study, the non-charge-balanced {111}〈110〉 and/or {100}〈110〉 slip systems might therefore be activated, which also accords with the observed low mobility of thick dislocation clusters (Figure 7 (b)). Although the presence of a small quantity of Fe might also change the slip system of dislocations in ionic MgO (such as Mg0.8Fe0.2O) [39], however, in this study on the plastic deformation of highpurity (99.95%) MgO, the small-size effect is believed to activate the dislocation gliding on the non-charge-balanced slip systems. Once dislocations were formed on the unusual slip systems, their parallel arrangement and aligned movement along one specific direction (mainly 90°, Figures 5 (b) (c), 7 (a), and Supplemental Video 1) would yield a stress shielding effect to lower the energy of the whole dislocation piling system [18, 19]. The shielding effect would reduce the strain hardening rate but increase the density of total accommodated dislocations [18, 19], accordingly enhancing both the hardenability and ductility of the nanosized, originally-brittle ionic MgO pillars. Also, when the pillar size decreased to the nanoscale, the number of geometrically necessary dislocations would increase to elevate hardenability [40]. Additional hardening mechanisms might include that (1) dislocation starvation (or difficult nucleation) in the nanosized pillars would cause a near-elastic strain rather than plastic deformation, i.e. a steep stress-strain response [11-12, 17, 41-45], and (2) surface bond reconstruction to lower surface energy would generate more stable shortened bonds to stiffen/strengthen nanosized materials [10, 46-49]. However, at the late stage of plastic deformation when some dislocations moving along the other specific direction (51.4°) intersected with these along 90° (Figure 7 (a) and Supplemental Video 1), cracks would then be initiated to cause the fracture of the pillars. Additionally, the brittle-ductile transition might possibly be strain rate dependent which is definitely worth further investigations. According to the Orowan equation [37], if the strain


Nano Letters 1 2 3 4 5 6 7 8 9 10 11 12 13 14 15 16 17 18 19 20 21 22 23 24 25 26 27 28 29 30 31 32 33 34 35 36 37 38 39 40 41 42 43 44 45 46 47 48 49 50 51 52 53 54 55 56 57 58 59 60

rate increases, both the density and gliding speed of dislocations will increase, causing more intense clustering of dislocations. Once the spacing of the piled dislocations (in the shielding model [18, 19]) markedly decreases, the local stress associated with dislocation activities will drastically increase to that sufficient to break bonds; an instability will occur and result in the brittle fracture of the pillars [19].

In summary, the indentation tests of micro- to nanosized single-crystalline 〈111〉 MgO pillars and the in-situ TEM compression tests of 〈111〉 MgO nanopillars suggested a smallsize effect on the strength and plasticity of ionic materials. When the conventional chargebalanced {110}〈110〉 slip system was prohibited, abnormal dislocation gliding would be activated alternatively on the non-charge-balanced {111}〈110〉 or {100}〈110〉 slip system, only at the nanoscale, probably owing to the change of energy criterion and the transition of slip systems caused by a nanosize-equivalent hydrostatic pressure effect. On the {111}〈110〉 or {100}〈110〉 slip system, the aligned movement of those activated dislocations to cause a stress shielding effect would accordingly enhance the hardenability and plasticity of the nanosized ionic MgO pillars.


Page 12 of 29

Page 13 of 29 1 2 3 4 5 6 7 8 9 10 11 12 13 14 15 16 17 18 19 20 21 22 23 24 25 26 27 28 29 30 31 32 33 34 35 36 37 38 39 40 41 42 43 44 45 46 47 48 49 50 51 52 53 54 55 56 57 58 59 60

Nano Letters

ASSOCIATED CONTENT

Supporting Information The Supporting Information includes materials and methods, and the supplementary videos of in-situ TEM observation on the nanocompression of single-crystalline 〈111〉 nanopillars. This material is available free of charge via the Internet at http://pubs.acs.org.

AUTHOR INFORMATION

Corresponding Author *(S.Y.C.) E-mail: [email protected]. Phone: +886-3-5715131 ext. 33806. Fax: +8863-5722366.

Author Contributions C.-C.Y. and S.-Y.L. contributed equally. Y.-C.H. and C.-H.T. contributed equally.

Notes The authors declare no competing financial interest.

ACKNOWLEDGMENTS

The authors gratefully thank the financial supports for this research by the Ministry of Science and Technology (MOST), Taiwan, under Grant Nos. MOST 102-2221-E-007-150MY3, MOST 106-2218-E-007-018 and MOST 107-3017-F-007-003, and in part by the “High Entropy Materials Center” of the Ministry of Education (MOE), Taiwan.


Nano Letters 1 2 3 4 5 6 7 8 9 10 11 12 13 14 15 16 17 18 19 20 21 22 23 24 25 26 27 28 29 30 31 32 33 34 35 36 37 38 39 40 41 42 43 44 45 46 47 48 49 50 51 52 53 54 55 56 57 58 59 60

REFERENCES

(1) Hirth, J. P.; Lothe. J. Theory of Dislocations, 2nd Ed. (Wiley, New York, 1982). (2) Argon, A. S.; Orowan. E. Philos. Mag. 1964, 9, 1003–1021. (3) Argon, A. S.; Orowan. E. Philos. Mag. 1964, 9, 1023–1039. (4) Korte, S.; Clegg, W. J. Philos. Mag. 2011, 91, 1150–1162. (5) Smallman, R. E.; Bishop, R. J. Modern Physical Metallurgy and Materials Engineering, 6th Ed. (Reed Educational and Professional Publishing, London, UK, 1999). (6) Brinckmann, S.; Kim, J. Y.; Greer, J. R. Phys. Rev. Lett. 2008, 100, 155502. (7) Hulse, C. O.; Copley, S. M.; Pask, J. A. J. Am. Ceramic Soc. 1963, 46, 317–323. (8) Li, L.; Weidner, D. J.; Chen, J.; Vaughan, M. T.; Davis, M.; Durham, W. B. J. Appl. Phys.

2004, 95, 8357–8365. (9) Gerbericha, W. W.; Mook, W. M.; Perrey, C. R.; Carter, C. B.; Baskes, M. I.; Mukherjee, R.; Gidwani, A.; Heberlein, J.; McMurry, P. H.; Girshick, S. L. J. Mech. Phys. Solids 2003, 51, 979–992. (10) Chang, Q. S. Prog. Mater. Sci. 2009, 54, 179–307. (11) Zhu, T.; Li, J. Prog. Mater. Sci. 2010, 55, 710–757. (12) Greer, J. R.; De Hosson, J. Th. M. Prog. Mater. Sci. 2011, 56, 654–724. (13) Kiener, D.; Hosemann, P.; Maloy, S. A.; Minor, A. M. Nat. Mater. 2011, 10, 608–613. (14) Wang, L.; Zhang, Z.; Han, X. NPG Asia Mater. 2013, 5, e40. (15) Yu, Q.; Qi, L.; Mishra, R. K.; Li, J.; Minor, A. M. PNAS 2013, 110, 13289–13293. (16) Seo, J. H.; Park, H. S.; Yoo, Y.; Seong, T. Y.; Li, J.; Ahn, J. P.; Kim, B.; Choi, I. S. Nano Lett. 2013, 13, 5112–5116. (17) O¨stlund, F.; Howie, P. R.; Ghisleni, R.; Korte, S.; Leifer, K.; Clegg, W. J.; Michler, J. Philos. Mag. 2011, 91, 1190–1199. (18) Tanguy, D.; Razafindrazaka, M.; Delafosse, D. Acta Mater. 2008, 56, 2441–2449.


Page 14 of 29

Page 15 of 29 1 2 3 4 5 6 7 8 9 10 11 12 13 14 15 16 17 18 19 20 21 22 23 24 25 26 27 28 29 30 31 32 33 34 35 36 37 38 39 40 41 42 43 44 45 46 47 48 49 50 51 52 53 54 55 56 57 58 59 60

Nano Letters

(19) Gerberich, W. W.; Michler, J.; Mook, W. M.; Ghisleni, R.; O¨stlund, F.; Stauffer, D. D.; Ballarini, R. J. Mater. Res. 2009, 24, 898–906. (20) Kumar, S.; Li, X.; Haque, A.; Gao, H. Nano Lett. 2011, 11, 2510–2516. (21) Huang, L.; Li, Q. J.; Shan, Z. W.; Li, J.; Sun, J.; Ma, E. Nat. Commun. 2011, 2, 547. (22) Yue, Y.; Chen, N.; Li, X.; Zhang, S.; Zhang, Z.; Chen, M.; Han, X. Nano Lett. 2013, 13, 3812–3816. (23) Wang, J. W.; Narayanan, S.; Huang, J. Y.; Zhang, Z.; Zhu, T.; Mao, S. X. Nat. Commun.

2013, 4, 2340. (24) Liu, B. Y.; Wang, J.; Li, B.; Lu, L.; Zhang, X. Y.; Shan, Z. W.; Li, J.; Jia, C. L.; Sun, J.; Ma, E. Nat. Commun. 2013, 4, 3297. (25) Wang, L.; Teng, J.; Liu, P.; Hirata, A.; Ma, E.; Zhang, Z.; Chen, M.; Han, X. Nat. Commun. 2014, 5, 4402. (26) Apple, F.; Bethge, H.; Messerschmidt, U. Phys. Stat. Sol. 1977, 42, 61–71. (27) Huang, J. Y.; Zheng, H.; Mao, S. X.; Li, Q.; Wang, G. T. Nano Lett. 2011, 11, 1618– 1622. (28) Issa, I.; Amodeo, J.; Joly‐Pottuz, L.; Réthoré, J.; Esnouf, C.; Garnier, V.; Morthomas, J.; Masenelli‐Varlot, K. European Microscopy Congress. Proceedings 2016, 1, 255–256. (29) Shackelford, J. F.; Alexander, W.; Park, J. S. Materials Science and Engineering, 2nd Ed. (CRC Press, Boca Raton, FL, 1994). (30) King, R. B. Int. J. Solids Structures 1987, 23, 1657–1664. (31) Sakai, M. J. Mater. Res. 2009, 24, 831–843. (32) Yang, Y.; Ye, J. C.; Lu, J.; Liu, F. X.; Liaw, P. K. Acta Mater. 2009, 57, 1613–1623. (33) Carrez, P.; Cordier, P.; Devincre, B.; Kubin, L. P. Mater. Sci. Eng. A-Struct. Mater. Prop. Microstruct. Process. 2005, 400–401, 325–328. (34) Samayoa, M. J.; Haque, M. A.; Cohen, P. H. J. Micromech. Microeng. 2008, 18, 095005


Nano Letters 1 2 3 4 5 6 7 8 9 10 11 12 13 14 15 16 17 18 19 20 21 22 23 24 25 26 27 28 29 30 31 32 33 34 35 36 37 38 39 40 41 42 43 44 45 46 47 48 49 50 51 52 53 54 55 56 57 58 59 60

(1–7). (35) Kiener, D.; Motz. C.; Rester, M.; Jenko, M.; Dehm, G. Mater. Sci. Eng. A-Struct. Mater. Prop. Microstruct. Process. 2007, 459, 262–272. (36) Zhu, Y. T.; Liao, X. Z.; Wu, X. L. Prog. Mater. Sci. 2012, 57, 1–62. (37) Abbaschian, R.; Abbaschian, L.; Reed-Hill, R. E. Physical Metallurgy Principles, 4th Ed. (Cengage Learning, New York, 2009). (38) Amodeo, J.; Carrez, P.; Cordier, P. Philos. Mag. 2012, 12, 1523–1541. (39) Stretton, I.; Heidelbach, F.; Mackwell, S.; Langenhorst, F. Earth Planetary Sci. Lett.

2001, 194, 229–240. (40) Kiener, D.; Guruprasad, P. J.; Keralavarma, S. M.; Dehm, G.; Benzerga, A. A. Acta Mater. 2011, 59, 3825–3840. (41) Zhu, T.; Li, J.; Ogata, S.; Yip, S. MRS Bulletin 2009, 34, 167–172. (42) Tse, J. S. J. Superhard Mater. 2010, 32, 177–191. (43) Greer, J. R.; Oliver, W. C.; Nix, W. D. Acta Mater. 2005, 53, 1821–1830. (44) Greer J. R.; Nix, W. D. Phys. Rev. B 2006, 73, 245410. (45) Shan, Z. W.; Mishra, R. K.; Syed Asif, S. A.; Warren, O. L.; Minor, A. M. Nat. Mater.

2008, 7, 115–119. (46) Sun, C. Q. Phys. Rev. B 2004, 69, 045105. (47) Ding, Y.; Sun, C. Q.; Zhou, Y. C. J. Appl. Phys. 2008, 103, 084317. (48) He, M. R.; Shi, Y.; Zhou, W.; Chen, J. W.; Yan, Y. J.; Zhu, J. Appl. Phys. Lett. 2009, 95, 091912. (49) Sun, Y.; Wang, Y.; Pan, J. S.; Wang, L. L.; Sun, C. Q. J. Phys. Chem. C 2009, 113, 14696–14701.


Page 16 of 29

Page 17 of 29 1 2 3 4 5 6 7 8 9 10 11 12 13 14 15 16 17 18 19 20 21 22 23 24 25 26 27 28 29 30 31 32 33 34 35 36 37 38 39 40 41 42 43 44 45 46 47 48 49 50 51 52 53 54 55 56 57 58 59 60

Nano Letters

Table 1. Possible slip planes and slip vectors of single-crystalline 〈111〉 MgO pillars, and the resolved shear stress τ on each slip system along 〈111〉 compression (corresponding to Figure 8 (d) and the equation τ = σ n cos θ cos φ ; σ n : normal stress, θ : the angle between the force and the normal direction of the slip plane, φ : the angle between the force and the slip vector).

Slip Plane {110}

{111}

{100}

cosθ

ADδ, BDδ, CDδ

cos90° (0)

ABδ, ACδ, BCδ

cos35° (0.82)

ABC

cos0° (1.00)

ABD, ACD, BCD

cos70° (0.34)

DEF (×3) *

cos53° (0.60)

Slip Vector 〈110〉

〈110〉

〈110〉

cosφ

τ

AD, BD, CD

cos35° (0.82)

0

AB, AC, BC

cos90° (0)

0

AB, AC, BC

cos90° (0)

0

AD, BD, CD

cos35° (0.82)

AB, AC, BC

cos90° (0)

AD, BD, CD

cos35° (0.82)

* DEF (×3): three planes equivalent to the DEF plane.


0.28σn 0 0.49σn

Nano Letters 1 2 3 4 5 6 7 8 9 10 11 12 13 14 15 16 17 18 19 20 21 22 23 24 25 26 27 28 29 30 31 32 33 34 35 36 37 38 39 40 41 42 43 44 45 46 47 48 49 50 51 52 53 54 55 56 57 58 59 60

FIGURE CAPTIONS

Figure 1. (a) SEM image of nanosized single-crystalline 〈111〉 MgO pillar for indentation tests; (b) TEM lattice images and SAD pattern of single-crystalline 〈111〉 MgO nanopillar for in-situ TEM compression tests.

Figure 2. (a) Average indentation load–depth curves of micro- to nanosized single-crystalline 〈111〉 MgO pillars (with diameters of 3200−100 nm; 20 DMA data points at different depths for one measurement curve and five measurement curves for each average fitting curve); (b) contact area–depth curves of indenter tip to a blanket MgO wafer (given by the instrument) and to the micro- to nanosized MgO pillars (calibrated by the King model).

Figure 3. (a) Ratios of indenter contact areas (given by the instrument) to calibrated contact areas (by the King model) for micro- to nanosized MgO pillars; (b) calibrated hardness of the micro- to nanosized single-crystalline 〈111〉 MgO pillars (at different indentation depths).

Figure 4. SEM images of micro- to nanosized single-crystalline 〈111〉 MgO pillars after indentation tests: (a) 3200, (b) 1600, (c) 800, (d) 400, (e) 200, (f) 100 nm in diameter.

Figure 5. (a) Stress−strain curves of in-situ TEM compression tests of single-crystalline 〈111〉 MgO nanopillars (purple: stress−strain curve of single-crystalline 〈111〉 GaAs nanopillar for comparison; dashed lines: elastic responses, orange circle: yielding range); (b) (c) in-situ TEM observations of dislocation activities in the MgO nanopillars under compression (white arrows: dislocation gliding).

Figure 6. TEM lattice images of single-crystalline 〈111〉 MgO nanopillar (a) before and (b) after in-situ TEM compression (dashed lines: marker lines at the same position of the pillar, referred to the dark spots at the surfaces of the pillar; regions C and D: protrusions (piles-up) of MgO lattice at the edges of the pillar tip); (c) magnified lattice image of region C.

Figure 7. Statistics of (a) moving directions and (b) gliding speeds of dislocations in single-


Page 18 of 29

Page 19 of 29 1 2 3 4 5 6 7 8 9 10 11 12 13 14 15 16 17 18 19 20 21 22 23 24 25 26 27 28 29 30 31 32 33 34 35 36 37 38 39 40 41 42 43 44 45 46 47 48 49 50 51 52 53 54 55 56 57 58 59 60

Nano Letters

crystalline 〈111〉 ionic MgO and covalent GaAs nanopillars during in-situ TEM compression (angle 90°: along the pillar/compression axis, 0°: parallel to the compression plane; columns in (a) and symbols in (b): experimental measurements, curves: fittings).

Figure 8. (a) and (b) FFT TEM lattice images of single-crystalline 〈111〉 MgO nanopillars after in-situ TEM compression (for a small extent of plastic deformation; circles: dislocations, lines: slip lines, red arrows: compression direction); (c) determination of zone axis (referred to the detected SAD pattern in Figure 1 (b)). (d) Thompson Tetrahedron and corresponding cubic cell/coordinate system (red arrow: compression direction, green arrow: zone axis); (e) projections of slip planes (yellow, green and blue triangles) and slip vectors (red arrow lines) on { 211 } plane (that includes the red triangle BCD’ in (d1)) for: (e1) {110}〈110〉, (e2) {111}〈110〉, (e3) {100}〈110〉 respectively.

Figure 9. (a) HRTEM observation (zone axis 〈 211 〉) of 〈111〉 MgO nanopillar after in-situ TEM compression (for a small extent of plastic deformation): (a1) FFT diffraction pattern, (a2) 2D lattice image marked with dislocations. (b) Atomistic model (zone axis 〈 211 〉) of defective 〈111〉 MgO crystal: (b1) atomistic lattice inserted with a b = 1/2 〈110〉 dislocation on a {111}〈110〉 slip system, (b2) FFT diffraction pattern, (b3) FFT 2D lattice image marked with the dislocation. (c) Schematic illustration of the geometry of the 1/2 〈110〉 dislocation and the {111}〈110〉 slip system in the 〈111〉 MgO nanopillar (green arrow: 〈 211 〉 zone axis, green frame: { 211 } projection plane; red arrow: 〈111〉 compression direction, red triangle: one {111} plane as the compression plane; blue triangle: another {111} plane as the slip plane of the 1/2 〈110〉 dislocation; black arrow: moving direction of the 1/2 〈110〉 dislocation, grey arrow: projection of the dislocation movement on the { 211 } plane).


Nano Letters 1 2 3 4 5 6 7 8 9 10 11 12 13 14 15 16 17 18 19 20 21 22 23 24 25 26 27 28 29 30 31 32 33 34 35 36 37 38 39 40 41 42 43 44 45 46 47 48 49 50 51 52 53 54 55 56 57 58 59 60

Figure 1. (a) SEM image of nanosized single-crystalline 〈111〉 MgO pillar for indentation tests; (b) TEM lattice images and SAD pattern of single-crystalline 〈111〉 MgO nanopillar for in-situ TEM compression tests.


Page 20 of 29

Page 21 of 29 1 2 3 4 5 6 7 8 9 10 11 12 13 14 15 16 17 18 19 20 21 22 23 24 25 26 27 28 29 30 31 32 33 34 35 36 37 38 39 40 41 42 43 44 45 46 47 48 49 50 51 52 53 54 55 56 57 58 59 60

Nano Letters

Figure 2. (a) Average indentation load–depth curves of micro- to nanosized single-crystalline 〈111〉 MgO pillars (with diameters of 3200−100 nm; 20 DMA data points at different depths for one measurement curve and five measurement curves for each average fitting curve); (b) contact area–depth curves of indenter tip to a blanket MgO wafer (given by the instrument) and to the micro- to nanosized MgO pillars (calibrated by the King model).


Nano Letters 1 2 3 4 5 6 7 8 9 10 11 12 13 14 15 16 17 18 19 20 21 22 23 24 25 26 27 28 29 30 31 32 33 34 35 36 37 38 39 40 41 42 43 44 45 46 47 48 49 50 51 52 53 54 55 56 57 58 59 60

Figure 3. (a) Ratios of indenter contact areas (given by the instrument) to calibrated contact areas (by the King model) for micro- to nanosized MgO pillars; (b) calibrated hardness of the micro- to nanosized single-crystalline 〈111〉 MgO pillars (at different indentation depths).


Page 22 of 29

Page 23 of 29 1 2 3 4 5 6 7 8 9 10 11 12 13 14 15 16 17 18 19 20 21 22 23 24 25 26 27 28 29 30 31 32 33 34 35 36 37 38 39 40 41 42 43 44 45 46 47 48 49 50 51 52 53 54 55 56 57 58 59 60

Nano Letters

Figure 4. SEM images of micro- to nanosized single-crystalline 〈111〉 MgO pillars after indentation tests: (a) 3200, (b) 1600, (c) 800, (d) 400, (e) 200, (f) 100 nm in diameter.


Nano Letters 1 2 3 4 5 6 7 8 9 10 11 12 13 14 15 16 17 18 19 20 21 22 23 24 25 26 27 28 29 30 31 32 33 34 35 36 37 38 39 40 41 42 43 44 45 46 47 48 49 50 51 52 53 54 55 56 57 58 59 60

Figure 5. (a) Stress−strain curves of in-situ TEM compression tests of single-crystalline 〈111〉 MgO nanopillars (purple: stress−strain curve of single-crystalline 〈111〉 GaAs nanopillar for comparison; dashed lines: elastic responses, orange circle: yielding range); (b) (c) in-situ TEM observations of dislocation activities in the MgO nanopillars under compression (white arrows: dislocation gliding).


Page 24 of 29

Page 25 of 29 1 2 3 4 5 6 7 8 9 10 11 12 13 14 15 16 17 18 19 20 21 22 23 24 25 26 27 28 29 30 31 32 33 34 35 36 37 38 39 40 41 42 43 44 45 46 47 48 49 50 51 52 53 54 55 56 57 58 59 60

Nano Letters

Figure 6. TEM lattice images of single-crystalline 〈111〉 MgO nanopillar (a) before and (b) after in-situ TEM compression (dashed lines: marker lines at the same position of the pillar, referred to the dark spots at the surfaces of the pillar; regions C and D: protrusions (piles-up) of MgO lattice at the edges of the pillar tip); (c) magnified lattice image of region C.


Nano Letters 1 2 3 4 5 6 7 8 9 10 11 12 13 14 15 16 17 18 19 20 21 22 23 24 25 26 27 28 29 30 31 32 33 34 35 36 37 38 39 40 41 42 43 44 45 46 47 48 49 50 51 52 53 54 55 56 57 58 59 60

Figure 7. Statistics of (a) moving directions and (b) gliding speeds of dislocations in singlecrystalline 〈111〉 ionic MgO and covalent GaAs nanopillars during in-situ TEM compression (angle 90°: along the pillar/compression axis, 0°: parallel to the compression plane; columns in (a) and symbols in (b): experimental measurements, curves: fittings).


Page 26 of 29

Page 27 of 29 1 2 3 4 5 6 7 8 9 10 11 12 13 14 15 16 17 18 19 20 21 22 23 24 25 26 27 28 29 30 31 32 33 34 35 36 37 38 39 40 41 42 43 44 45 46 47 48 49 50 51 52 53 54 55 56 57 58 59 60

Nano Letters

Figure 8. (a) and (b) FFT TEM lattice images of single-crystalline 〈111〉 MgO nanopillars after in-situ TEM compression (for a small extent of plastic deformation; circles: dislocations, lines: slip lines, red arrows: compression direction); (c) determination of zone axis (referred to the detected SAD pattern in Figure 1 (b)). (d) Thompson Tetrahedron and corresponding cubic cell/coordinate system (red arrow: compression direction, green arrow: zone axis); (e) projections of slip planes (yellow, green and blue triangles) and slip vectors (red arrow lines) on { 211 } plane (that includes the red triangle BCD’ in (d1)) for: (e1) {110}〈110〉, (e2) {111}〈110〉, (e3) {100}〈110〉 respectively.


Nano Letters 1 2 3 4 5 6 7 8 9 10 11 12 13 14 15 16 17 18 19 20 21 22 23 24 25 26 27 28 29 30 31 32 33 34 35 36 37 38 39 40 41 42 43 44 45 46 47 48 49 50 51 52 53 54 55 56 57 58 59 60

Figure 9. (a) HRTEM observation (zone axis 〈 211 〉) of 〈111〉 MgO nanopillar after in-situ TEM compression (for a small extent of plastic deformation): (a1) FFT diffraction pattern, (a2) 2D lattice image marked with dislocations. (b) Atomistic model (zone axis 〈 211 〉) of defective 〈111〉 MgO crystal: (b1) atomistic lattice inserted with a b = 1/2 〈110〉 dislocation on a {111}〈110〉 slip system, (b2) FFT diffraction pattern, (b3) FFT 2D lattice image marked with the dislocation. (c) Schematic illustration of the geometry of the 1/2 〈110〉 dislocation and the {111}〈110〉 slip system in the 〈111〉 MgO nanopillar (green arrow: 〈 211 〉 zone axis, green frame: { 211 } projection plane; red arrow: 〈111〉 compression direction, red triangle: one {111} plane as the compression plane; blue triangle: another {111} plane as the slip plane of the 1/2 〈110〉 dislocation; black arrow: moving direction of the 1/2 〈110〉 dislocation, grey arrow: projection of the dislocation movement on the { 211 } plane).


Page 28 of 29

Page 29 of 29 1 2 3 4 5 6 7 8 9 10 11 12 13 14 15 16 17 18 19 20 21 22 23 24 25 26 27 28 29 30 31 32 33 34 35 36 37 38 39 40 41 42 43 44 45 46 47 48 49 50 51 52 53 54 55 56 57 58 59 60

Nano Letters

Table of Contents Graphic


Small-Size-Induced Plasticity and Dislocation ... - ACS Publications

Recommend Documents