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Study of rapid grain boundary migration in a nanocrystalline Ni thin film
Materials Science and Engineering A 528 (2011) 1628–1635
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Study of rapid grain boundary migration in a nanocrystalline Ni thin film Josh Kacher a,∗ , I.M. Robertson a , Matt Nowell b , J. Knapp c , Khalid Hattar c a
Department of Materials Science and Engineering, University of Illinois, 1304 West Green Street, Urbana, IL 61801, USA EDAX/TSL 392 East 12300 South, Draper, UT 84020, USA c Sandia National Laboratories, PO Box 5800, Albuquerque, NM 87185, USA b
a r t i c l e
i n f o
Article history: Received 8 September 2010 Received in revised form 22 October 2010 Accepted 29 October 2010
Keywords: Abnormal grain growth Pulsed laser deposited Ni Electron microscopy Annealing
a b s t r a c t Grain boundary migration associated with abnormal grain growth in pulsed-laser deposited Ni was studied in real time by annealing electron transparent films in situ in the transmission electron microscope. The resulting texture evolution and grain boundary types produced were evaluated by ex situ electron backscatter diffraction of interrupted anneals. The combination of these two techniques allowed for the investigation of grain growth rates, grain morphologies, and the evolution of the orientation and grain boundary distributions. Grain boundaries were found to progress in a sporadic, start/stop fashion with no evidence of a characteristic grain growth rate. The orientations of the abnormally growing grains were found to be predominately 1 1 1//ND throughout the annealing process. A high fraction of twin boundaries developed during the annealing process. The intermittent growth from different locations of the grain boundary is discussed in terms of a vacancy diffusion model for grain growth. © 2010 Elsevier B.V. All rights reserved.
1. Introduction Nanocrystalline metals have high strength, but the high volume fraction and degree of perfection of grain boundaries makes them unstable at relatively low temperature. Examples of this instability include the observation of grain growth, including abnormal grain growth, at temperatures below 600 K in Au [1–3], Cu [2], Pd [4], and Ni [5–12]. The observation of abnormal grain growth has been attributed to the initial microstructure, including the distribution of the grain size, misorientation, grain boundary energy and mobility, triple junction mobility, impurities, microstrain, topological features such as curvature and number of boundary faces, etc. [13–19]. Transmission electron microscope (TEM) annealing experiments (both in situ and ex situ) of electrodeposited nanocrystalline Ni have shown that the grain growth is driven by a few grains at the expense of the surrounding nanograins, which is consistent with abnormal grain growth. This growth proceeds in a sporadic start/stop fashion and lacks a characteristic growth rate. Past in situ TEM experiments that focused on growth front velocities have suggested that this start/stop growth results from the accumulation of impurity elements at the grain boundaries [9]. Experiments on the texture development in annealed electrodeposited nanocrystalline Ni revealed the development of a 1 1 1//ND texture via a twinning mechanism [10] with some studies also reporting an initial 1 1 4//ND texture [11].
Computational models have suggested a number of explanations for abnormal grain growth behavior but lack a consensus on the relative importance of the various potential factors. Work by Rollett et al. suggests that grain boundary characteristics and texture play a large role in encouraging abnormal grain growth [20]. Simulations by Lee et al. suggest that initial grain size distribution is the dominant factor in determining the likelihood of abnormal grain growth [21]. Other models have pointed to the importance of vacancy build up in the grain interiors from the absorption of grain boundaries [22], solute particle pinning effects [23], mobility of grain boundary triple junctions [13,24] and surface effects [1] in determining grain growth characteristics. Both computational and experimental studies have highlighted the importance of grain boundary characteristics as a driving factor in grain growth with, however, a seeming lack of consensus on which characteristics are most important [25–27]. In this paper, abnormal grain growth in 99.997% pure pulsedlaser deposited (PLD) Ni is studied using in situ TEM annealing to gather information on the dynamics of grain growth, specifically focusing on grain boundary migration rates and evolving grain morphologies. Ex situ electron backscatter diffraction (EBSD) scans in the scanning electron microscope (SEM) were used to provide information on the texture evolution at different stages of the annealing process. 2. Experiment
∗ Corresponding author. Tel.: +1 801 318 9158; fax: +1 801 422 0516. E-mail address: [email protected] (J. Kacher). 0921-5093/$ – see front matter © 2010 Elsevier B.V. All rights reserved. doi:10.1016/j.msea.2010.10.109
The synthesis conditions of the PLD Ni are as follows: nanograined Ni thin films were grown on 1 0 0 rock salt sub-
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Fig. 1. Bright field images captured from video of the change in morphology of a deposited PLD Ni film during an in situ TEM annealing experiment. Annealing time is given in seconds in each frame.
strates to a nominal thickness of 70 nm by using the pulsed-laser deposition technique. The growth chamber achieved a base pressure of 2.7 × 105 Pa, with the pressure rising to about five times higher during the laser pulses. The KrF laser pulse operated at a wavelength of 248 nm with a pulse width of 34 ns full-width halfmaximum, a pulse rate of 35 Hz, and an energy density at the target of 1–2 J/cm2 . These conditions resulted in a film growth rate of 0.25 nm/s. A particle filter was used to minimize the number of molten droplets reaching the film. Analysis of the thin films showed the initial microstructure to have a random orientation with grain sizes ranging from 2 to 16 nm with an average size of 6 nm [8]. To prepare the sample for observation in the TEM, a 2 mm × 2 mm square grid was scored on the surface of the rock salt substrate. The rock salt was then dissolved in distilled water, allowing the Ni squares to float to the surface. These squares were scooped from the water using 3 mm clamping copper grids and allowed to dry. The samples were mounted on a Gatan double-tilt heating stage and observed in a JEOL 2010 LaB6 TEM operating at 200 kV. The in situ annealing experiments were performed by rapid heating to 623 K and holding at that temperature for the reported time; this temperature was selected based on previous work [28]. While the majority of the samples were heated until the nanocrystalline matrix was completely consumed, some were removed at earlier stages of the annealing process for ex situ characterization in the SEM. The temperature ramping in the TEM stage occurred rapidly, heating from room temperature to 623 K in approximately 45 s. It should be noted that local beam heating can raise the temperature by as much as 20 K, so the reported temperature should be taken as nominal [29]. Video of the grain growth process was captured using a CCD camera at a frame rate of 33 fps. Analysis of the videos was carried out using VirtualDub, Adobe Photoshop, and ImageJ software. Individual frames were selected from the video and grains were hand traced in order to measure the grain size (surface area) and morphology at discrete time steps. EBSD scans were taken from four different samples, each one having been annealed at nominally the same temperature but for different lengths of time. The analysis was performed using EDAX-
TSL OIM data collection software at an accelerating voltage of 20 kV. 10 m × 10 m scans at a step size of 10 nm were collected from each sample. The steps were set on a hexagonal grid to better resolve grain morphology. No additional sample preparation beyond that used for the TEM analysis was needed. The orientation data were used to provide additional information on the texture evolution of the Ni samples during the annealing process including grain size distributions, grain boundary misorientation angle distributions, pole figures, and inverse pole figure (IPF) maps. Parallel electron energy loss spectrometry (PEELS) analysis was done on the annealed material to detect possible buildups of trace solute elements at the grain boundaries. PEELS spectra were collected on a JEOL 2010F (S)TEM operating in STEM mode at 10 nm increments stepping across the grain boundary. Several grain boundaries were inspected in this manner with no detection of any solute elements either at the grain boundaries or in the grain interiors. This result confirms previous compositional analysis on PLD Ni films [30]. 3. Results Analysis of the videos shows that the grain growth is abnormal with a few grains growing quickly at the expense of the surrounding nanocrystalline matrix. This growth characteristic is evident in the series of bright-field images captured from video and presented in Fig. 1. Several grains that are well-separated spatially grow and consume the other grains. The trigger initiating grain growth does not occur simultaneously for all grains, as is evident from comparison of the images in Fig. 1. These images show grain growth initiating at different times across the field-of-view. The nature of the actual initiator could not be determined. Also evident in this series of images is that the grains do not grow uniformly in all directions. On closer inspection it is found that even in directions of significant growth, the growth fronts progress in a sporadic fashion, starting at one location, stopping, and proceeding again from a different one. Consequently, the grain shapes evolve in a random fashion with no obvious inclination towards curvature driven
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Fig. 2. The change in the morphology of an individual grain as a function of time. The capture times are 41:94, 45:65, 47:71, 53:89, 55:36, 55:40, 60:06, and 79:38 s, respectively.
growth. These growth features are captured in Fig. 2, in which an individual grain has been extracted from a video to show its evolution; changing imaging conditions due to elastic distortions during the growth process makes it difficult to discern the position of the grain boundary exactly, especially when the growth is small and this causes minor discrepancies in selecting the position. The arrows indicate the regions of the grain undergoing the most growth from one image to the next. It is only towards the latter stages of growth that a more uniform expansion of the grain occurs. In addition to the growth occurring at different locations in an apparent uncorrelated fashion, the velocity of a particular front was not uniform. The fastest recorded occurred for the grain shown in Fig. 2 between images d–f and was 282 nm/s. To quantify the grain growth rate, the perimeter of several growing grains was traced manually using ImageJ software. Two time intervals, 5 and 30 s, were selected to determine if a characteristic grain growth rate existed. The change in the projected grain area as a function of annealing time for the two time intervals are plotted in Fig. 3. The plots show the growth rates vary strongly between grains and there is no characteristic growth rate, confirming the qualitative observation of Fig. 2. However, certain trends emerge which suggest that the growth rates should be delineated into stages. The first stage, occurring at the onset of grain growth, appears to be a period of normal-like grain growth. In this stage, the grains remain under ∼25 nm in size and grow slowly (Figs. 1a and b). Grain size estimates from video frames after 13.6 s and 24.7 s of annealing showed that the average grain size increased from 17 nm to 23 nm during this time, giving an average growth rate of 0.54 nm/s. Since the film thickness was ∼70 nm, multiple through thickness grains were observed, making it difficult to delineate the grain boundaries and perform a detailed analysis. The second stage is dominated by abnormal grain growth in which a small fraction of predominantly through-thickness grains grew quickly to consume the matrix. The abnormal growth was sporadic and progression of individual growth fronts followed a start/stop pattern with velocity measurements of progressing growth fronts yielding maximum values approaching 300 nm/s. These velocity measurements were taken between individual frames, or at a time change of 1/33rd of a second, making them almost instantaneous in nature. However, averaging the velocities over the length of abnormal growth leads to a measured velocity of 3.5 nm/s. Also during this stage, it was observed that the grain areas of individual grains (treating the grains in a two dimensional sense) seem to grow in a quadratic relationship with time (Fig. 4).
Multiple grains were observed to follow a similar quadratic growth pattern as that shown in Fig. 4. This stage occurs over a very short time period (<15 s generally) and only occurs for grains completely surrounded by a nanograin matrix; for example, Figs. 1c and d. As abnormally growing grains can begin growth bordering other large grains or impinge on other abnormally growing grains immediately following the onset of growth, this second stage tends to be restricted to those grains in which abnormal growth initiates first. The third stage of grain growth has its onset when abnormally growing grains significantly impinge on each other (Fig. 1e). Grain boundary interfaces between abnormal grains tend to be strongly pinned, limiting available growth to a reduced fraction of the grain boundary. During this stage, the progression of individual growth fronts follows a similar pattern to second stage growth; that is a sporadic, start/stop motion. However, due to the reduction in available growth fronts, the increase in area of individual grains no longer follows a quadratic growth pattern but instead becomes as sporadic and start/stop in nature as the progression of the individual growth fronts. Evidence of third stage growth is shown in Fig. 3 where the instance of first impingement of the abnormally growing grain with another abnormally growing grain is marked. As can be seen in the figure, the quadratic growth of the grain area is disrupted and the growth switches to a start/stop motion. For completeness, it should also be mentioned that normal growth has been reported to dominate once the nanocrystalline matrix is consumed [28,31]. This should then be considered a fourth stage of grain growth. The duration of these annealing experiments was insufficient to observe significant normal grain growth. At all stages of growth, significant twinning occurs. Twinning was observed to occur at initiation of abnormal growth, in which case the twins are narrow, grow with the growth front but do not widen during the process. Twinning initiated later in the annealing process can cause a significant portion of the grain to rotate; this result is described in detail later. Fig. 5 illustrates an example of twin nucleation at the initiation of abnormal grain growth. Still frames from a video show a twin forming in a nano-grain early in the growth process while the surrounding matrix is still nanocrystalline. As can be seen from the sequence of images, as the grain grows, the twin maintains constant width and grows at the same rate as the grain boundary progresses. The EBSD results provide additional insight into which grains are most prone to growth, how this growth serves to reduce the energy of the system, and what path the texture evolution follows to reach the final annealed microstructure. Current EBSD capabilities allow
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Fig. 3. (a) Grain size vs. time at 30-s intervals. (b) Grain size vs. time for select grains at 5-s intervals. Each line corresponds to a different grain although the same symbol identifies the same grain in both plots. The asterisk denotes the instance of first impingement with another abnormally growing grain (only shown for one of the grains).
for the reliable characterization of grains on the order of ∼40 nm. For this reason, only the abnormally growing grains and not the nanocrystalline matrix could be analyzed using EBSD. In Fig. 6 IPF maps of the PLD Ni microstructure at different annealing times are presented (note that local bending of the Ni film causes elongation of the grains). No data cleanup was used and points that could not be indexed reliably were left black. From each data set, the (1 0 0) pole figure, grain size distribution, and grain boundary misorientation distribution charts were generated (also shown in Fig. 6). Statistical data from the microstructure were limited to those grains sufficiently large for reliable characterization using EBSD. The EBSD data show that there is very little texture variation of the abnormally growing grains over the course of the annealing process. The microstructure of the abnormally growing grains quickly develops a 1 1 1//ND texture and remains 1 1 1//ND throughout the annealing process. It is unclear from the EBSD data whether 1 1 1//ND grains are more likely to grow, or if nanograins early in the process rotate to the 1 1 1//ND orientations and then begin to grow. Both paths lead to identical microstructures. The 1 1 1//ND orientations for FCC materials are the lowest surface energy orientations, suggesting that the surface energy influences texture development in the material [1]. Also, the grain boundary misorientation distribution heavily favors a 60◦ rotation. When rotation axis information is included, it is seen that this comes from a high fraction of the boundaries being 3 twin boundaries (60◦ rotation about the 1 1 1 axis). A portion of the IPF map of the fully annealed microstructure is enlarged with the twin boundaries highlighted in Fig. 7 as an example of the distribution of twin boundaries in the material. Klement et al. reported in similar studies on texture evolution in nanocrystalline Ni the development of a 1 1 4//ND texture in the early stages of grain growth [11]. This difference in texture development may be due to a number of fac-
tors, including differences in sample composition (due to different sample synthesis methods), initial grain size, thickness, or annealing temperature. It is also possible that such a texture did exist but had already evolved into a 1 1 1//ND texture before any EBSD data were collected. It is also interesting to note from the EBSD data that there seems to be little variation in the peak grain size with time. This suggests that that the grains, independent of the surrounding grain size, reach a maximum diameter quickly after initiation of abnormal growth. That is, the grains do not grow uniformly over the sample. Rather, individual grains begin growth at varying times and quickly increase in size until a plateau is reached, at which point the growth velocity decreases dramatically. This is similar to the growth patterns seen during the in situ TEM annealing experiments (Fig. 3). For this reason, there is very little shift in the peak size distribution after a certain point in the annealing process. Earlier studies on abnormal grain growth in PLD Ni using EBSD reported an HCP phase developing at temperatures lower than those used in the current study [7]. The Hough peaks from the EBSD data were reanalyzed to probe for evidence of an HCP phase in both the fully and partially annealed samples. Only scattered EBSD patterns with low confidence were indexed as HCP phase, suggesting that no HCP was present. This is consistent with results from a previous study [7].
4. Discussion This dynamic in situ TEM study of abnormal grain growth in PLD Ni at a nominal temperature of 623 K showed that: The grains that exhibit abnormal growth are distributed randomly throughout the foil and initiation occurs at different times.
Fig. 4. (a) Grain size vs. time for a single grain with quadratic line fit. (b) Grain boundary migration distance vs. time for the same grain.
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Fig. 5. Bright-field image of a growing twin. Annealing times are 50:85, 55:79, 70:54, and 360 s. Twin is indicated by arrow.
During abnormal grain growth, the growth occurs from seemingly uncorrelated regions of the grain boundary at different times. That is, growth occurs in a start/stop fashion from different regions of the grain boundary and the location appears not to be determined by the curvature. Growth twins are created during all stages of abnormal grain growth. Twins initiated in the later stages of growth are responsible for causing large rotations of a significant fraction of a grain. Grain growth falls into groups with normal-like growth occurring first, followed by abnormal growth transitioning via impingement of abnormally growing grains to a stage of normal growth [32,3]. Similar stages have been identified in other studies, with the final stage being identified as growth towards thermal equilibrium. Although the initiation of abnormal grain growth at random locations throughout the film is obvious from the images presented in Fig. 1, there was no evidence differentiating those grains that experienced abnormal grain growth from ones that did not. Earlier studies of abnormal grain growth of PLD Ni reported a large number of stacking-fault tetrahedra formed in the grain interiors. This result was interpreted as evidence for vacancies being ejected from the consumed grain boundaries by the advancing grain boundary [8]. This result, along with the start/stop motion of the growth fronts, is qualitatively consistent with a model proposed by Estrin et al. [22]. They postulated that as growth fronts consume nanograin boundaries the lower density of grain boundaries in
comparison with the matrix would lead to excess free volume in the form of vacancies being expelled into the grain interiors. This generation of vacancies in the grain interiors leads to competing energy minimization mechanisms. An initial driving force for grain growth is energy minimization through reduction of grain boundary area. This reduction, however, leads to supersaturations of vacancies in the material which contributes to an increase in system energy. Thus, as long as the energy reduction through reduction in grain boundary area exceeds the energy increase through vacancy saturation, grain growth will continue. As the growth front progresses, more vacancies are ejected into the grains, leading to the supersaturation. This non-equilibrium condition leads to an unfavorable energy state that offsets any energy reductions gained by the elimination of grain boundaries, effectively locking further grain growth. The grain growth is stopped temporarily until the excess vacancies can diffuse into sinks such as other grain boundaries or dislocations or form lower energy configurations such as stacking-fault tetrahedra or dislocation loops. When the free energy from excess vacancies is reduced sufficiently, grain growth can resume, thus leading to the overall start/stop motion of grain growth. While Estrin et al.’s model best fits the data, it does not preclude other proposed models such as the influence of microstrains [1] or grain boundary diffusion [33]. While these models, as well as others, have been considered, the data obtained in this study are inconclusive in determining their fidelity.
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Fig. 6. IPF maps, pole figures, misorientation, and grain size distribution from four EBSD scans progressing from shortest time to longest time at temperature. Legend for the IPF scans is given above.
Variations in grain boundary energy can account for the irregular shaped morphology of grains as well as the tendency for growth to sometimes progress from one location, stop, and begin again at a different location. Grain boundary velocities and energies have been theoretically and experimentally shown to be dependent on the five degrees of freedom of the grain boundary [34,35]. Uniform grain boundary energies would most likely lead to grain boundary curvature as the driving force for grain growth, which would result in uniform expansion along all portions of the grain boundary. When variability of grain boundary energy is factored in, the growth follows a more irregular path as the best means of minimization of the total system energy. This would result in formation of both concave and convex grains. Also, the grain boundaries are not constant in character, defining character as both the misorientation across the boundary as well as the grain boundary plane orientation. As each nanograin is consumed, the growth front is presented with a new misorientation across the boundary which can lead to sig-
nificant variations in the grain boundary energy. Again, assuming the growth front follows the path of greatest energy reduction of the total system, it is possible to rationalize the start/stop behavior from different locations. If the energy reduction from changing a high energy boundary to a lower energy boundary through twinning is greater than the energy increase from introducing a new twin boundary, twinning is likely to occur [36,37]. With this mechanism in mind, it is appropriate to consider why the annealing twins observed in the early stages of grain growth are narrower than those observed in the IPF maps. Grain boundary energies are at least partially dependent on the misorientation across the grain boundary. The twin changes this misorientation, so it is expected that the twin width will be the same as the length of the high energy grain boundary. At the early stages of the annealing experiments, a misorientation angle is maintained locally as the grain sizes are still around 10 nm, and consequentially twin widths tend to be narrow. Later, after
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Fig. 7. IPF map from fully annealed sample with recrystallization twin boundaries highlighted in white.
the grains have grown and abnormally growing grains begin to impinge on each other, the grains maintain a constant misorientation angle over longer lengths of their boundaries, resulting in significant portions of the grains rotating through twinning. Wu et al. considered the ability to use the Generalized Planar Fault Energy as a predictor for the generation of stacking faults and twins in nanocrystalline Ni [38]. They concluded that while the Generalized Planar Fault Energy could predict the propensity for forming stacking faults it did not for twins. It was suggested that this was attributable to the mechanism of forming a twin if it was produced by the emission of the same type of partial dislocation on successive planes. However, this finding is in contrast to Kibey et al. who found that the critical stress for twinning in FCC metals was linearly dependent on the unstable twinning energy [39]. In the present work, the twins are annealing twins which may be generated in response to a growth accident, a reduction in the total free energy of the system, or via stacking faults or fault packets [37]. Mahajan et al. proposed annealing twins would form if the driving force for grain boundary migration was curvature driven and that the curvature was in fact comprised of steps and facets with some residing on {1 1 1} planes [37]. If the migration rate of the boundary is high, the probability for a growth accident increases and this generates the partial dislocations that ultimately form the twin. The resolution of the current experiments was insufficient to determine if grain boundary defects or ledges played a role in determining the twin nucleation site. Similar studies on abnormal grain growth in Ni using in situ TEM have been carried out, notably by Hibbard et al. [9,40]. The results from that study were similar in that the sporadic and unpredictable behavior of grain growth was observed. However, in Hibbard et al.’s study, no stage of smooth grain growth was observed and the growth appeared to be more sporadic. Growth front velocities could not be directly compared due to the experiments being performed at different annealing temperatures. The main difference between the two studies was the higher concentration of impurity elements, notably the level of sulfur, in Hibbard et al.’s samples that was introduced during the pulse current electrodeposition. Hibbard et al. suggested the most likely cause of the start/stop behavior of the growth fronts was accumulation of sulfur at the progressing boundaries. Sulfur buildup in the grain boundaries reduces the grain boundary energy, which has a pinning effect on the grain boundaries. Subsequent precipitation of the sulfur lowers the pinning force, allowing the growth front to resume progression [41]. However, it should be noted that Abraham et al. showed using 3D atom probe analysis that in nanograined Ni, produced by electrodeposition in nickel sulfate baths, there was no evidence for segregation in the as-deposited condition, and the segregated level after an anneal
at 573 K for 60 min was low and unlikely to be sufficient to retard grain growth [5]. Impurities are unlikely to be important in explaining the present results as PEELS analysis, not shown, of the PLD Ni showed no solute element accumulation at the grain boundaries, which eliminates impurities as the responsible element in this case. However, it is interesting to note that while the growth front velocities may have been altered significantly by the presence of sulfur in the previous study, the general characteristics of the abnormal grain growth persisted. The influence of the proximity of the free surfaces in the electron transparent foil on the general characteristics of the growth front can be discounted as similar characteristics emerge for growth of micrometer-sized grains in the interior of the material. The 3D Xray diffraction time-resolved observations of normal grain growth in micrograined Al samples agree well both with the irregular growth morphologies of grains as well as a lack of any characteristic growth rates [42]. Also, the systems observed using 3D X-ray diffraction tended to have large grains (>5 m) and were observed over long time periods (on the order of hours), confirming that at least some of the observed phenomena in the PLD nickel samples are not limited to thin film or nanocrystalline materials. Several other potentially influential factors have been suggested by computational methods such as molecular dynamics and Monte Carlo simulations and these focused mainly on the structural details of grain boundaries. These include the importance of rough vs. smooth boundaries [27,43], grain boundary mobilities [26], grain boundary grooving [44], and local vs. global grain boundary energy states [21]. The described experiments did not provide the needed information to confirm or refute reported simulation results. 5. Conclusions In situ annealing experiments on PLD Ni show that grain growth proceeds in an abnormal fashion. The grain growth is sporadic, and grain boundary velocities vary significantly from grain to grain. Individual growth fronts progress in a start/stop motion and this results in the development of both convex and concave grain shapes. Nano-twins tend to form early in the annealing process with broader twins forming and reorienting grains later in the process. Supporting EBSD data show that a 1 1 1//ND texture quickly develops in the annealed Ni and that a significant fraction of the grain boundaries are 3 low energy twin boundaries. The growth motion is at least qualitatively consistent with that described by a vacancy diffusion model, and the final microstructure suggests that the energy of the system is minimized through the formation of low energy grain boundaries and favorable grain orientation for surface energy minimization. Acknowledgement The work at the University of Illinois (JK and IMR) was supported by the US Department of Energy Office of Basic Energy Sciences, Division of Materials Science, under award No. DEFG02-07ER46443. The microscopy was carried out in the Center for Microanalysis of Materials, University of Illinois as well as at EDAX-TSL facilities in Draper, Utah. Sandia National Laboratories is a multi-program laboratory operated by Sandia Corporation, a wholly owned subsidiary of Lockheed Martin company, for the U.S. Department of Energy’s National Nuclear Security Administration under contract DE-AC04-94AL85000. KH acknowledges support from the Division of Materials Science and Engineering, Office of Basic Energy Sciences, U.S. Department of Energy at Sandia. References [1] C.V. Thompson, Interface Sci. 6 (1998) 85–93.
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Materials Science and Engineering A journal homepage: www.elsevier.com/locate/msea
Study of rapid grain boundary migration in a nanocrystalline Ni thin film Josh Kacher a,∗ , I.M. Robertson a , Matt Nowell b , J. Knapp c , Khalid Hattar c a
Department of Materials Science and Engineering, University of Illinois, 1304 West Green Street, Urbana, IL 61801, USA EDAX/TSL 392 East 12300 South, Draper, UT 84020, USA c Sandia National Laboratories, PO Box 5800, Albuquerque, NM 87185, USA b
a r t i c l e
i n f o
Article history: Received 8 September 2010 Received in revised form 22 October 2010 Accepted 29 October 2010
Keywords: Abnormal grain growth Pulsed laser deposited Ni Electron microscopy Annealing
a b s t r a c t Grain boundary migration associated with abnormal grain growth in pulsed-laser deposited Ni was studied in real time by annealing electron transparent films in situ in the transmission electron microscope. The resulting texture evolution and grain boundary types produced were evaluated by ex situ electron backscatter diffraction of interrupted anneals. The combination of these two techniques allowed for the investigation of grain growth rates, grain morphologies, and the evolution of the orientation and grain boundary distributions. Grain boundaries were found to progress in a sporadic, start/stop fashion with no evidence of a characteristic grain growth rate. The orientations of the abnormally growing grains were found to be predominately 1 1 1//ND throughout the annealing process. A high fraction of twin boundaries developed during the annealing process. The intermittent growth from different locations of the grain boundary is discussed in terms of a vacancy diffusion model for grain growth. © 2010 Elsevier B.V. All rights reserved.
1. Introduction Nanocrystalline metals have high strength, but the high volume fraction and degree of perfection of grain boundaries makes them unstable at relatively low temperature. Examples of this instability include the observation of grain growth, including abnormal grain growth, at temperatures below 600 K in Au [1–3], Cu [2], Pd [4], and Ni [5–12]. The observation of abnormal grain growth has been attributed to the initial microstructure, including the distribution of the grain size, misorientation, grain boundary energy and mobility, triple junction mobility, impurities, microstrain, topological features such as curvature and number of boundary faces, etc. [13–19]. Transmission electron microscope (TEM) annealing experiments (both in situ and ex situ) of electrodeposited nanocrystalline Ni have shown that the grain growth is driven by a few grains at the expense of the surrounding nanograins, which is consistent with abnormal grain growth. This growth proceeds in a sporadic start/stop fashion and lacks a characteristic growth rate. Past in situ TEM experiments that focused on growth front velocities have suggested that this start/stop growth results from the accumulation of impurity elements at the grain boundaries [9]. Experiments on the texture development in annealed electrodeposited nanocrystalline Ni revealed the development of a 1 1 1//ND texture via a twinning mechanism [10] with some studies also reporting an initial 1 1 4//ND texture [11].
Computational models have suggested a number of explanations for abnormal grain growth behavior but lack a consensus on the relative importance of the various potential factors. Work by Rollett et al. suggests that grain boundary characteristics and texture play a large role in encouraging abnormal grain growth [20]. Simulations by Lee et al. suggest that initial grain size distribution is the dominant factor in determining the likelihood of abnormal grain growth [21]. Other models have pointed to the importance of vacancy build up in the grain interiors from the absorption of grain boundaries [22], solute particle pinning effects [23], mobility of grain boundary triple junctions [13,24] and surface effects [1] in determining grain growth characteristics. Both computational and experimental studies have highlighted the importance of grain boundary characteristics as a driving factor in grain growth with, however, a seeming lack of consensus on which characteristics are most important [25–27]. In this paper, abnormal grain growth in 99.997% pure pulsedlaser deposited (PLD) Ni is studied using in situ TEM annealing to gather information on the dynamics of grain growth, specifically focusing on grain boundary migration rates and evolving grain morphologies. Ex situ electron backscatter diffraction (EBSD) scans in the scanning electron microscope (SEM) were used to provide information on the texture evolution at different stages of the annealing process. 2. Experiment
∗ Corresponding author. Tel.: +1 801 318 9158; fax: +1 801 422 0516. E-mail address: [email protected] (J. Kacher). 0921-5093/$ – see front matter © 2010 Elsevier B.V. All rights reserved. doi:10.1016/j.msea.2010.10.109
The synthesis conditions of the PLD Ni are as follows: nanograined Ni thin films were grown on 1 0 0 rock salt sub-
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Fig. 1. Bright field images captured from video of the change in morphology of a deposited PLD Ni film during an in situ TEM annealing experiment. Annealing time is given in seconds in each frame.
strates to a nominal thickness of 70 nm by using the pulsed-laser deposition technique. The growth chamber achieved a base pressure of 2.7 × 105 Pa, with the pressure rising to about five times higher during the laser pulses. The KrF laser pulse operated at a wavelength of 248 nm with a pulse width of 34 ns full-width halfmaximum, a pulse rate of 35 Hz, and an energy density at the target of 1–2 J/cm2 . These conditions resulted in a film growth rate of 0.25 nm/s. A particle filter was used to minimize the number of molten droplets reaching the film. Analysis of the thin films showed the initial microstructure to have a random orientation with grain sizes ranging from 2 to 16 nm with an average size of 6 nm [8]. To prepare the sample for observation in the TEM, a 2 mm × 2 mm square grid was scored on the surface of the rock salt substrate. The rock salt was then dissolved in distilled water, allowing the Ni squares to float to the surface. These squares were scooped from the water using 3 mm clamping copper grids and allowed to dry. The samples were mounted on a Gatan double-tilt heating stage and observed in a JEOL 2010 LaB6 TEM operating at 200 kV. The in situ annealing experiments were performed by rapid heating to 623 K and holding at that temperature for the reported time; this temperature was selected based on previous work [28]. While the majority of the samples were heated until the nanocrystalline matrix was completely consumed, some were removed at earlier stages of the annealing process for ex situ characterization in the SEM. The temperature ramping in the TEM stage occurred rapidly, heating from room temperature to 623 K in approximately 45 s. It should be noted that local beam heating can raise the temperature by as much as 20 K, so the reported temperature should be taken as nominal [29]. Video of the grain growth process was captured using a CCD camera at a frame rate of 33 fps. Analysis of the videos was carried out using VirtualDub, Adobe Photoshop, and ImageJ software. Individual frames were selected from the video and grains were hand traced in order to measure the grain size (surface area) and morphology at discrete time steps. EBSD scans were taken from four different samples, each one having been annealed at nominally the same temperature but for different lengths of time. The analysis was performed using EDAX-
TSL OIM data collection software at an accelerating voltage of 20 kV. 10 m × 10 m scans at a step size of 10 nm were collected from each sample. The steps were set on a hexagonal grid to better resolve grain morphology. No additional sample preparation beyond that used for the TEM analysis was needed. The orientation data were used to provide additional information on the texture evolution of the Ni samples during the annealing process including grain size distributions, grain boundary misorientation angle distributions, pole figures, and inverse pole figure (IPF) maps. Parallel electron energy loss spectrometry (PEELS) analysis was done on the annealed material to detect possible buildups of trace solute elements at the grain boundaries. PEELS spectra were collected on a JEOL 2010F (S)TEM operating in STEM mode at 10 nm increments stepping across the grain boundary. Several grain boundaries were inspected in this manner with no detection of any solute elements either at the grain boundaries or in the grain interiors. This result confirms previous compositional analysis on PLD Ni films [30]. 3. Results Analysis of the videos shows that the grain growth is abnormal with a few grains growing quickly at the expense of the surrounding nanocrystalline matrix. This growth characteristic is evident in the series of bright-field images captured from video and presented in Fig. 1. Several grains that are well-separated spatially grow and consume the other grains. The trigger initiating grain growth does not occur simultaneously for all grains, as is evident from comparison of the images in Fig. 1. These images show grain growth initiating at different times across the field-of-view. The nature of the actual initiator could not be determined. Also evident in this series of images is that the grains do not grow uniformly in all directions. On closer inspection it is found that even in directions of significant growth, the growth fronts progress in a sporadic fashion, starting at one location, stopping, and proceeding again from a different one. Consequently, the grain shapes evolve in a random fashion with no obvious inclination towards curvature driven
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Fig. 2. The change in the morphology of an individual grain as a function of time. The capture times are 41:94, 45:65, 47:71, 53:89, 55:36, 55:40, 60:06, and 79:38 s, respectively.
growth. These growth features are captured in Fig. 2, in which an individual grain has been extracted from a video to show its evolution; changing imaging conditions due to elastic distortions during the growth process makes it difficult to discern the position of the grain boundary exactly, especially when the growth is small and this causes minor discrepancies in selecting the position. The arrows indicate the regions of the grain undergoing the most growth from one image to the next. It is only towards the latter stages of growth that a more uniform expansion of the grain occurs. In addition to the growth occurring at different locations in an apparent uncorrelated fashion, the velocity of a particular front was not uniform. The fastest recorded occurred for the grain shown in Fig. 2 between images d–f and was 282 nm/s. To quantify the grain growth rate, the perimeter of several growing grains was traced manually using ImageJ software. Two time intervals, 5 and 30 s, were selected to determine if a characteristic grain growth rate existed. The change in the projected grain area as a function of annealing time for the two time intervals are plotted in Fig. 3. The plots show the growth rates vary strongly between grains and there is no characteristic growth rate, confirming the qualitative observation of Fig. 2. However, certain trends emerge which suggest that the growth rates should be delineated into stages. The first stage, occurring at the onset of grain growth, appears to be a period of normal-like grain growth. In this stage, the grains remain under ∼25 nm in size and grow slowly (Figs. 1a and b). Grain size estimates from video frames after 13.6 s and 24.7 s of annealing showed that the average grain size increased from 17 nm to 23 nm during this time, giving an average growth rate of 0.54 nm/s. Since the film thickness was ∼70 nm, multiple through thickness grains were observed, making it difficult to delineate the grain boundaries and perform a detailed analysis. The second stage is dominated by abnormal grain growth in which a small fraction of predominantly through-thickness grains grew quickly to consume the matrix. The abnormal growth was sporadic and progression of individual growth fronts followed a start/stop pattern with velocity measurements of progressing growth fronts yielding maximum values approaching 300 nm/s. These velocity measurements were taken between individual frames, or at a time change of 1/33rd of a second, making them almost instantaneous in nature. However, averaging the velocities over the length of abnormal growth leads to a measured velocity of 3.5 nm/s. Also during this stage, it was observed that the grain areas of individual grains (treating the grains in a two dimensional sense) seem to grow in a quadratic relationship with time (Fig. 4).
Multiple grains were observed to follow a similar quadratic growth pattern as that shown in Fig. 4. This stage occurs over a very short time period (<15 s generally) and only occurs for grains completely surrounded by a nanograin matrix; for example, Figs. 1c and d. As abnormally growing grains can begin growth bordering other large grains or impinge on other abnormally growing grains immediately following the onset of growth, this second stage tends to be restricted to those grains in which abnormal growth initiates first. The third stage of grain growth has its onset when abnormally growing grains significantly impinge on each other (Fig. 1e). Grain boundary interfaces between abnormal grains tend to be strongly pinned, limiting available growth to a reduced fraction of the grain boundary. During this stage, the progression of individual growth fronts follows a similar pattern to second stage growth; that is a sporadic, start/stop motion. However, due to the reduction in available growth fronts, the increase in area of individual grains no longer follows a quadratic growth pattern but instead becomes as sporadic and start/stop in nature as the progression of the individual growth fronts. Evidence of third stage growth is shown in Fig. 3 where the instance of first impingement of the abnormally growing grain with another abnormally growing grain is marked. As can be seen in the figure, the quadratic growth of the grain area is disrupted and the growth switches to a start/stop motion. For completeness, it should also be mentioned that normal growth has been reported to dominate once the nanocrystalline matrix is consumed [28,31]. This should then be considered a fourth stage of grain growth. The duration of these annealing experiments was insufficient to observe significant normal grain growth. At all stages of growth, significant twinning occurs. Twinning was observed to occur at initiation of abnormal growth, in which case the twins are narrow, grow with the growth front but do not widen during the process. Twinning initiated later in the annealing process can cause a significant portion of the grain to rotate; this result is described in detail later. Fig. 5 illustrates an example of twin nucleation at the initiation of abnormal grain growth. Still frames from a video show a twin forming in a nano-grain early in the growth process while the surrounding matrix is still nanocrystalline. As can be seen from the sequence of images, as the grain grows, the twin maintains constant width and grows at the same rate as the grain boundary progresses. The EBSD results provide additional insight into which grains are most prone to growth, how this growth serves to reduce the energy of the system, and what path the texture evolution follows to reach the final annealed microstructure. Current EBSD capabilities allow
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Fig. 3. (a) Grain size vs. time at 30-s intervals. (b) Grain size vs. time for select grains at 5-s intervals. Each line corresponds to a different grain although the same symbol identifies the same grain in both plots. The asterisk denotes the instance of first impingement with another abnormally growing grain (only shown for one of the grains).
for the reliable characterization of grains on the order of ∼40 nm. For this reason, only the abnormally growing grains and not the nanocrystalline matrix could be analyzed using EBSD. In Fig. 6 IPF maps of the PLD Ni microstructure at different annealing times are presented (note that local bending of the Ni film causes elongation of the grains). No data cleanup was used and points that could not be indexed reliably were left black. From each data set, the (1 0 0) pole figure, grain size distribution, and grain boundary misorientation distribution charts were generated (also shown in Fig. 6). Statistical data from the microstructure were limited to those grains sufficiently large for reliable characterization using EBSD. The EBSD data show that there is very little texture variation of the abnormally growing grains over the course of the annealing process. The microstructure of the abnormally growing grains quickly develops a 1 1 1//ND texture and remains 1 1 1//ND throughout the annealing process. It is unclear from the EBSD data whether 1 1 1//ND grains are more likely to grow, or if nanograins early in the process rotate to the 1 1 1//ND orientations and then begin to grow. Both paths lead to identical microstructures. The 1 1 1//ND orientations for FCC materials are the lowest surface energy orientations, suggesting that the surface energy influences texture development in the material [1]. Also, the grain boundary misorientation distribution heavily favors a 60◦ rotation. When rotation axis information is included, it is seen that this comes from a high fraction of the boundaries being 3 twin boundaries (60◦ rotation about the 1 1 1 axis). A portion of the IPF map of the fully annealed microstructure is enlarged with the twin boundaries highlighted in Fig. 7 as an example of the distribution of twin boundaries in the material. Klement et al. reported in similar studies on texture evolution in nanocrystalline Ni the development of a 1 1 4//ND texture in the early stages of grain growth [11]. This difference in texture development may be due to a number of fac-
tors, including differences in sample composition (due to different sample synthesis methods), initial grain size, thickness, or annealing temperature. It is also possible that such a texture did exist but had already evolved into a 1 1 1//ND texture before any EBSD data were collected. It is also interesting to note from the EBSD data that there seems to be little variation in the peak grain size with time. This suggests that that the grains, independent of the surrounding grain size, reach a maximum diameter quickly after initiation of abnormal growth. That is, the grains do not grow uniformly over the sample. Rather, individual grains begin growth at varying times and quickly increase in size until a plateau is reached, at which point the growth velocity decreases dramatically. This is similar to the growth patterns seen during the in situ TEM annealing experiments (Fig. 3). For this reason, there is very little shift in the peak size distribution after a certain point in the annealing process. Earlier studies on abnormal grain growth in PLD Ni using EBSD reported an HCP phase developing at temperatures lower than those used in the current study [7]. The Hough peaks from the EBSD data were reanalyzed to probe for evidence of an HCP phase in both the fully and partially annealed samples. Only scattered EBSD patterns with low confidence were indexed as HCP phase, suggesting that no HCP was present. This is consistent with results from a previous study [7].
4. Discussion This dynamic in situ TEM study of abnormal grain growth in PLD Ni at a nominal temperature of 623 K showed that: The grains that exhibit abnormal growth are distributed randomly throughout the foil and initiation occurs at different times.
Fig. 4. (a) Grain size vs. time for a single grain with quadratic line fit. (b) Grain boundary migration distance vs. time for the same grain.
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Fig. 5. Bright-field image of a growing twin. Annealing times are 50:85, 55:79, 70:54, and 360 s. Twin is indicated by arrow.
During abnormal grain growth, the growth occurs from seemingly uncorrelated regions of the grain boundary at different times. That is, growth occurs in a start/stop fashion from different regions of the grain boundary and the location appears not to be determined by the curvature. Growth twins are created during all stages of abnormal grain growth. Twins initiated in the later stages of growth are responsible for causing large rotations of a significant fraction of a grain. Grain growth falls into groups with normal-like growth occurring first, followed by abnormal growth transitioning via impingement of abnormally growing grains to a stage of normal growth [32,3]. Similar stages have been identified in other studies, with the final stage being identified as growth towards thermal equilibrium. Although the initiation of abnormal grain growth at random locations throughout the film is obvious from the images presented in Fig. 1, there was no evidence differentiating those grains that experienced abnormal grain growth from ones that did not. Earlier studies of abnormal grain growth of PLD Ni reported a large number of stacking-fault tetrahedra formed in the grain interiors. This result was interpreted as evidence for vacancies being ejected from the consumed grain boundaries by the advancing grain boundary [8]. This result, along with the start/stop motion of the growth fronts, is qualitatively consistent with a model proposed by Estrin et al. [22]. They postulated that as growth fronts consume nanograin boundaries the lower density of grain boundaries in
comparison with the matrix would lead to excess free volume in the form of vacancies being expelled into the grain interiors. This generation of vacancies in the grain interiors leads to competing energy minimization mechanisms. An initial driving force for grain growth is energy minimization through reduction of grain boundary area. This reduction, however, leads to supersaturations of vacancies in the material which contributes to an increase in system energy. Thus, as long as the energy reduction through reduction in grain boundary area exceeds the energy increase through vacancy saturation, grain growth will continue. As the growth front progresses, more vacancies are ejected into the grains, leading to the supersaturation. This non-equilibrium condition leads to an unfavorable energy state that offsets any energy reductions gained by the elimination of grain boundaries, effectively locking further grain growth. The grain growth is stopped temporarily until the excess vacancies can diffuse into sinks such as other grain boundaries or dislocations or form lower energy configurations such as stacking-fault tetrahedra or dislocation loops. When the free energy from excess vacancies is reduced sufficiently, grain growth can resume, thus leading to the overall start/stop motion of grain growth. While Estrin et al.’s model best fits the data, it does not preclude other proposed models such as the influence of microstrains [1] or grain boundary diffusion [33]. While these models, as well as others, have been considered, the data obtained in this study are inconclusive in determining their fidelity.
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Fig. 6. IPF maps, pole figures, misorientation, and grain size distribution from four EBSD scans progressing from shortest time to longest time at temperature. Legend for the IPF scans is given above.
Variations in grain boundary energy can account for the irregular shaped morphology of grains as well as the tendency for growth to sometimes progress from one location, stop, and begin again at a different location. Grain boundary velocities and energies have been theoretically and experimentally shown to be dependent on the five degrees of freedom of the grain boundary [34,35]. Uniform grain boundary energies would most likely lead to grain boundary curvature as the driving force for grain growth, which would result in uniform expansion along all portions of the grain boundary. When variability of grain boundary energy is factored in, the growth follows a more irregular path as the best means of minimization of the total system energy. This would result in formation of both concave and convex grains. Also, the grain boundaries are not constant in character, defining character as both the misorientation across the boundary as well as the grain boundary plane orientation. As each nanograin is consumed, the growth front is presented with a new misorientation across the boundary which can lead to sig-
nificant variations in the grain boundary energy. Again, assuming the growth front follows the path of greatest energy reduction of the total system, it is possible to rationalize the start/stop behavior from different locations. If the energy reduction from changing a high energy boundary to a lower energy boundary through twinning is greater than the energy increase from introducing a new twin boundary, twinning is likely to occur [36,37]. With this mechanism in mind, it is appropriate to consider why the annealing twins observed in the early stages of grain growth are narrower than those observed in the IPF maps. Grain boundary energies are at least partially dependent on the misorientation across the grain boundary. The twin changes this misorientation, so it is expected that the twin width will be the same as the length of the high energy grain boundary. At the early stages of the annealing experiments, a misorientation angle is maintained locally as the grain sizes are still around 10 nm, and consequentially twin widths tend to be narrow. Later, after
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Fig. 7. IPF map from fully annealed sample with recrystallization twin boundaries highlighted in white.
the grains have grown and abnormally growing grains begin to impinge on each other, the grains maintain a constant misorientation angle over longer lengths of their boundaries, resulting in significant portions of the grains rotating through twinning. Wu et al. considered the ability to use the Generalized Planar Fault Energy as a predictor for the generation of stacking faults and twins in nanocrystalline Ni [38]. They concluded that while the Generalized Planar Fault Energy could predict the propensity for forming stacking faults it did not for twins. It was suggested that this was attributable to the mechanism of forming a twin if it was produced by the emission of the same type of partial dislocation on successive planes. However, this finding is in contrast to Kibey et al. who found that the critical stress for twinning in FCC metals was linearly dependent on the unstable twinning energy [39]. In the present work, the twins are annealing twins which may be generated in response to a growth accident, a reduction in the total free energy of the system, or via stacking faults or fault packets [37]. Mahajan et al. proposed annealing twins would form if the driving force for grain boundary migration was curvature driven and that the curvature was in fact comprised of steps and facets with some residing on {1 1 1} planes [37]. If the migration rate of the boundary is high, the probability for a growth accident increases and this generates the partial dislocations that ultimately form the twin. The resolution of the current experiments was insufficient to determine if grain boundary defects or ledges played a role in determining the twin nucleation site. Similar studies on abnormal grain growth in Ni using in situ TEM have been carried out, notably by Hibbard et al. [9,40]. The results from that study were similar in that the sporadic and unpredictable behavior of grain growth was observed. However, in Hibbard et al.’s study, no stage of smooth grain growth was observed and the growth appeared to be more sporadic. Growth front velocities could not be directly compared due to the experiments being performed at different annealing temperatures. The main difference between the two studies was the higher concentration of impurity elements, notably the level of sulfur, in Hibbard et al.’s samples that was introduced during the pulse current electrodeposition. Hibbard et al. suggested the most likely cause of the start/stop behavior of the growth fronts was accumulation of sulfur at the progressing boundaries. Sulfur buildup in the grain boundaries reduces the grain boundary energy, which has a pinning effect on the grain boundaries. Subsequent precipitation of the sulfur lowers the pinning force, allowing the growth front to resume progression [41]. However, it should be noted that Abraham et al. showed using 3D atom probe analysis that in nanograined Ni, produced by electrodeposition in nickel sulfate baths, there was no evidence for segregation in the as-deposited condition, and the segregated level after an anneal
at 573 K for 60 min was low and unlikely to be sufficient to retard grain growth [5]. Impurities are unlikely to be important in explaining the present results as PEELS analysis, not shown, of the PLD Ni showed no solute element accumulation at the grain boundaries, which eliminates impurities as the responsible element in this case. However, it is interesting to note that while the growth front velocities may have been altered significantly by the presence of sulfur in the previous study, the general characteristics of the abnormal grain growth persisted. The influence of the proximity of the free surfaces in the electron transparent foil on the general characteristics of the growth front can be discounted as similar characteristics emerge for growth of micrometer-sized grains in the interior of the material. The 3D Xray diffraction time-resolved observations of normal grain growth in micrograined Al samples agree well both with the irregular growth morphologies of grains as well as a lack of any characteristic growth rates [42]. Also, the systems observed using 3D X-ray diffraction tended to have large grains (>5 m) and were observed over long time periods (on the order of hours), confirming that at least some of the observed phenomena in the PLD nickel samples are not limited to thin film or nanocrystalline materials. Several other potentially influential factors have been suggested by computational methods such as molecular dynamics and Monte Carlo simulations and these focused mainly on the structural details of grain boundaries. These include the importance of rough vs. smooth boundaries [27,43], grain boundary mobilities [26], grain boundary grooving [44], and local vs. global grain boundary energy states [21]. The described experiments did not provide the needed information to confirm or refute reported simulation results. 5. Conclusions In situ annealing experiments on PLD Ni show that grain growth proceeds in an abnormal fashion. The grain growth is sporadic, and grain boundary velocities vary significantly from grain to grain. Individual growth fronts progress in a start/stop motion and this results in the development of both convex and concave grain shapes. Nano-twins tend to form early in the annealing process with broader twins forming and reorienting grains later in the process. Supporting EBSD data show that a 1 1 1//ND texture quickly develops in the annealed Ni and that a significant fraction of the grain boundaries are 3 low energy twin boundaries. The growth motion is at least qualitatively consistent with that described by a vacancy diffusion model, and the final microstructure suggests that the energy of the system is minimized through the formation of low energy grain boundaries and favorable grain orientation for surface energy minimization. Acknowledgement The work at the University of Illinois (JK and IMR) was supported by the US Department of Energy Office of Basic Energy Sciences, Division of Materials Science, under award No. DEFG02-07ER46443. The microscopy was carried out in the Center for Microanalysis of Materials, University of Illinois as well as at EDAX-TSL facilities in Draper, Utah. Sandia National Laboratories is a multi-program laboratory operated by Sandia Corporation, a wholly owned subsidiary of Lockheed Martin company, for the U.S. Department of Energy’s National Nuclear Security Administration under contract DE-AC04-94AL85000. KH acknowledges support from the Division of Materials Science and Engineering, Office of Basic Energy Sciences, U.S. Department of Energy at Sandia. References [1] C.V. Thompson, Interface Sci. 6 (1998) 85–93.
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