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NiO multilayers
Materials Science & Engineering A 568 (2013) 49–60
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Materials Science & Engineering A journal homepage: www.elsevier.com/locate/msea
Thermal stability of Ni/NiO multilayers Josh Kacher a,n, Patricia Elizaga a, Stephen D. House a, Khalid Hattar b, Matt Nowell c, I.M. Robertson a a b c
Materials Science and Engineering, University of Illinois at Urbana-Champaign, 1304 W Green St, Urbana, IL 61801, United States Sandia National Laboratories, Albuquerque, NM 87185, United States EDAX/TSL 392 East 12300 South, Draper, UT 84020, United States
a r t i c l e i n f o
abstract
Article history: Received 5 November 2012 Received in revised form 14 January 2013 Accepted 17 January 2013 Available online 23 January 2013
The effects of NiO multilayer additions to pulsed-laser deposited nanocrystalline Ni on the grain growth and texture evolution during annealing have been studied using a combination of in situ annealing in the transmission electron microscope and the recently developed transmission electron backscatter diffraction technique. Grain growth in pure Ni proceeds in an abnormal manner with a small number of grains growing rapidly at the expense of the surrounding nanocrystalline matrix. The addition of NiO layers in the Ni suppressed grain growth and twin formation upon annealing, with increasing oxide content resulting in decreasing extent of grain growth. In pure Ni, the post-annealing texture was found to be composed of a strong {001} texture coupled with the four {122} twin variants from the {001}. With the addition of NiO, annealing twinning, and hence the {122} orientation, was suppressed and instead the final texture was composed of the {001} and {111} orientations. & 2013 Elsevier B.V. All rights reserved.
Keywords: EBSD Electron microscopy Nanostructured materials Grain growth
1. Introduction Nanocrystalline metals have been garnering increased attention due to their unique properties such as exceptionally high strength [1], but microstructural instability at relatively low homologous temperatures limits their use. The high grain boundary energy density creates a high driving force for grain growth, both normal and abnormal, at temperatures as low as 600 K in various materials [2–9]. Abnormal grain growth has been attributed to factors such as the initial grain size distribution [10], orientation distribution [8,11–13], grain boundary characteristics [14–16], triple junction mobility [17], impurities in the form of solute atoms [18] and second phase particles [19], heterogeneous microstrains [5], and topological features such as grain boundary grooving [20]. Multilayered films and nanostructured metal-oxides are also of considerable interest due to their unique magnetic and electronic properties [21–25]. Efforts specifically targeted at the synthesis of Ni–NiO multilayers have included the periodic oxidation of metallic layers using pulsed oxygen molecular beams during deposition [26,27] and evaporation of NiO with an electron beam [26], with single layer thicknesses ranging from 2 to 10 nm. Studies on the thermal stability of Ni–NiO multilayer structures have shown that annealing can cause the multilayer structure to disintegrate, leaving dispersed spherical oxide particles in the metallic matrix. This can affect the electronic properties of the
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material, such as the loss of the angular dependence of magnetoresistance [24]. To investigate the reactivity between Ni and NiO, Symianakis et al. deposited Ni at near monolayer thickness onto a (001) oriented NiO substrate [28]. They found that the two systems were stable at room temperature, but on annealing the Ni layer oxidized, eventually causing all of the Ni to convert to NiO. The onset of this effect was at approximately 773 K with the oxygen diffusion accelerating with increasing temperature. Experimental studies of microstructure stability and grain growth in nanocrystalline and nanostructured materials have primarily utilized four characterization techniques: transmission electron microscopy (TEM) [3,4,6,7,29,30], x-ray diffraction [24,31], differential scanning calorimetry [32,33], and scanning electron microscopy (SEM) based electron backscatter diffraction (EBSD) [30,31]. By annealing nanocrystalline Ni films in situ in the TEM, the dynamics of grain growth also can be observed and the formation of an oxide layer during the annealing can be minimized. For example, studies have shown that grain growth in nanocrystalline materials often proceeds in an abnormal fashion, with a relatively small number of grains growing quickly at the expense of the surrounding nanocrystalline matrix [4,6,7,11]. This growth proceeds in a sporadic fashion; often beginning at one grain boundary, stopping, and continuing from a different location on the same boundary. The resulting grains have been observed to have an unusual morphology, with both convex and concave regions coexisting along the grain boundary. Such observations suggest that in-plane curvature driven growth is not the determining factor governing grain expansion [30]. Local velocity measurements have shown that individual growth fronts propagate
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at rates many times higher than the average growth rate would dictate [6]. Post mortem analysis following annealing showed a high density of vacancy-based defects in the grain interiors, including stacking-fault tetrahedra [4]. This was interpreted as being from the absorption and subsequent ejection of the excess free volume found in grain boundaries in PLD nanocrystalline nickel, highlighting vacancy diffusion as a potentially important mechanism dictating grain growth. Solute atoms are known to play an important role in stabilizing the microstructure, although the mechanisms for how this is accomplished are still being debated. The two prevailing theories essentially follow either a kinetic or a thermodynamic argument. The kinetic, or solute drag, approach, first developed by Cahn [34], focuses on the drag force that solute atoms can exert on grain boundaries during grain growth. Solute atoms occupy the free volume in the boundary and lower the grain boundary energy, thus reducing the driving force for grain growth. As the grains grow, the grain boundary area of individual grains proportionally increases and more solute atoms are accommodated in the boundary. In nanocrystalline materials, grain boundaries can propagate at high velocity and the grain boundary area of the individual grains increases rapidly, which along with the concurrent increase in solute content leads to complete boundary pinning and halts grain growth [35,36]. Solute drag theory, for example, has been used to explain the buildup of sulfur at grain boundaries after grain growth in nanocrystalline Ni [6]. Similarly, da Silva et al. found that additions of phosphorous to nanocrystalline cobalt improved microstructural stability, resulting in a delayed onset of grain growth and more homogeneous growth once it initiated. Using atom probe tomography, they showed that this reduced grain growth was linked to phosphorous segregation to the grain boundaries [33]. This knowledge has motivated the addition of sulfur and phosphorous during electrodeposition to stabilize the fine grain structure [37,38]. The thermodynamic approach considers the Gibbs free energy of the complete system [39–42]. Solute atom segregation to grain boundaries can lower the Gibbs free energy of the system, below that achievable through elimination of grain boundaries. Thus, there exists a solute concentration dependent metastable grain boundary area for the system, which once attained, removes the impetus for further grain growth. Thermodynamic arguments therefore predict that, as more solute atoms must be accommodated, the equilibrium grain size will decrease with increasing solute concentration. EBSD, x-ray diffraction, and selected area electron diffraction have provided useful tools in examining the texture evolution during grain growth. Most studies of FCC materials agree that a strong {111} fiber texture develops during the annealing process, but with some disagreement about the initial texture. Klement et al. reported an initial {114} texture at the onset of grain growth in both Ni and Ni-20 at % Fe, with twinning transitioning the texture to {111} [11]. Kim et al. reported an initial texture dominated by {001} in nanocrystalline Ni, with the {111} texture growing in strength at the expense of the {001} oriented grains with increasing annealing temperature [31]. However, the EBSD analysis in these studies was conducted at a minimum step size of 15 nm, excluding information about the smallest grains. Consequently, much of the grain boundary information was not obtained and the analysis was limited to only those grains that had reached a certain stage of grain growth. Recently, Rajasekhara et al. demonstrated that the initial texture of nanocrystalline PLD Ni thin films could be characterized using precession microscopy in a TEM, which allows orientation mapping with a minimum resolvable grain size of 4 nm [43]. The improved spatial resolution allowed them to identify a significant fraction of HCP-phase Ni even in the as-deposited
state. Keller and Geiss developed a technique similar to EBSD that is based on the collection of transmitted electrons in an SEM to form a diffraction pattern and is referred to as transmission EBSD or t-EBSD [44]. By reducing the thickness of the sample to a few tens to hundreds of nanometers, similar to the thickness of a TEM sample, and collecting only the transmitted electrons, the interaction volume becomes a narrow column instead of the larger teardrop shape from which backscattered electrons are collected. This can improve the spatial resolution to below 10 nm, compared to the effective resolution of traditional EBSD which, while dependent on multiple factors, is worse by at least a factor of 2 [45]. In this manner, orientation mapping can be done in an SEM at a resolution approaching that which is possible using precession microscopy in a TEM. The main limitation with using t-EBSD for nanocrystalline materials arises from the presence of multiple through-thickness grains, leading to mixed diffraction patterns and preventing reliable pattern indexing. In this study, nanocrystalline Ni–NiO multilayer thin films were deposited with varying thicknesses of the NiO layers. The thermal stability and grain growth characteristics of these films were investigated through TEM characterization, including in situ annealing, focused ion beam (FIB) lift-out for cross-sectional analysis, and post mortem t-EBSD investigations of the orientation and grain boundary distributions.
2. Experimental Nanocrystalline Ni samples were made using pulsed-laser deposition at Sandia National Laboratory [46]. The samples were deposited on a base of (001) oriented rock salt with a beam energy of 500 mJ at room temperature to a final thickness of approximately 70 nm. The growth chamber achieved a base pressure of approximately 5 10 7 Torr. Three sample compositions, Ni, nominally Ni–10 vol% NiO, and nominally Ni–20 vol% NiO, were synthesized to investigate the effects of sample composition and multilayer structure on the thermal stability. Pure Ni (99.97%) and pure NiO (99.9% ) targets were used during the deposition. Pure Ni was deposited at a nominal laser pulse frequency of 35 Hz until the desired thickness was achieved. Ni and NiO were deposited in alternating layers, with the Ni deposited at a pulse frequency of 35 Hz and the NiO at 7 Hz. For the 10 vol% NiO material, the layer thickness alternated between nominally 4.8 nm for pure Ni and 0.5 nm for the NiO. This was changed to nominally 4.5 nm for the pure Ni and 1.1 nm for the NiO in the 20 vol% NiO material. The volume percent of the NiO is a nominal value based on calibrated deposition rates. The sample deposition conditions are summarized in Table 1. Using the same conditions, Ni–NiO multilayers were also deposited on an amorphous SiO2 substrate to facilitate cross-sectional TEM examination of samples extracted using FIB machining. Plan-view samples were prepared for analysis in the TEM and SEM by scoring a 2 mm grid on the surface of the film and placing the salt crystal in deionized water. Upon dissolution of the salt, the samples floated to the surface and were collected using 3 mmdiameter clamping copper grids. Cross-section samples, to investigate the multilayer Ni–NiO structure, were prepared using an FEI Dual Table 1 Summary of materials investigated. Note that all layer thicknesses are nominal. Material
Ni layer thickness (nm)
NiO layer thickness (nm)
Pure Ni Ni–10 vol% NiO Ni–20 vol% NiO
4.8 4.5
0.5 1.1
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Beam 235 FIB and standard FIB lift-out and thinning techniques (see, for example [47]). A platinum layer was deposited prior to any ion milling to ensure the preservation of the underlying microstructure. Grain growth in the plan-view thin films was observed during in situ annealing in a Phillips CM30 LaB6 TEM. Samples were heated to 473 K and allowed to stabilize. Next they were heated to 623 K and held at that temperature for 5–7 min and, finally, the temperature was raised to 723 K. All stages of heating were done at a ramp rate of approximately 5 K/s. As no grain growth was observed at 473 K, the evolution of the microstructure was recorded only during the latter two stages of annealing. The grain growth was captured using a CCD camera at a frame rate of three frames per second for at least 5 min. This was found to be ample time for the microstructure to reach a state beyond which no noticeable changes were observed. Pre- and post-annealing characterization of the films was carried out using TEM analysis with the microscope operating at 300 kV. Images using multiple different diffraction vectors were collected to better resolve the grain structure. T-EBSD was used to study the texture and microstructure of the plan-view samples after the annealing was complete. Analysis was performed using EDAX/TSL orientation imaging microscopy (OIM) data collection software at an accelerating voltage of 30 kV.
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The samples were held at 51 from horizontal and at a working distance of approximately 4 mm using a custom built stage that allowed for transmitted electrons to be detected on the phosphor screen. Two 2 2 mm scans were collected from each sample at a step size of 5 nm and a scanning speed of 200 points per second for a total scan time of approximately 800 s. The collected data were used to construct inverse pole figure (IPF) maps and pole figures. No data interpolation was used and all points indexed with a confidence level below 0.1, as defined by the OIM software, were filtered from the data; these are displayed as black areas in the IPF maps. Both FCC- and HCPphase Ni were included in the search, but no HCP-phase Ni grains were reliably identified. The cross-sectional samples were characterized both before and after annealing using high angle annular dark-field scanning TEM (HAADF-STEM) imaging and energy-dispersive x-ray spectroscopy (EDS) line scans in a JEOL 2010 F (S)TEM operated at 200 kV. Annealing of the FIB lift-out samples was performed in situ in a JEOL 2010 LaB6 TEM operating at 200 kV after the FIB lift-out. As thermal stresses could cause the separation of the FIB lift-out sample from the copper grid, the heating rate was reduced to approximately 1 K/s.
Fig. 1. Cross-sectional characterization of the Ni–20 vol% NiO sample before (a) and after (b) annealing using HAADF-STEM imaging and post-annealing characterization using bright-field TEM imaging (c). The vacuum/SiO2 interface is delineated in (b) by a dotted line and different layers, Pt, Ni–NiO, and SiO2 are labeled for convenience. An inset with contrast levels adjusted and a dotted line added to highlight the Moire´ fringes is included in (c).
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3. Results 3.1. Pre-annealing characterization Cross-sectional examination of the Ni–20 vol% NiO material using HAADF-STEM imaging showed that the multilayer structure remained coherent during the deposition and sample preparation processes (Fig. 1a). The periodicity as well as the layer thickness
was preserved with no evidence of local rupturing in the layer structure. Post-annealing characterization showed the multilayer structure to still be largely unchanged. Annealing the films caused the separation of the Ni–NiO from the SiO2, evident by the black contrast layer between the two materials seen in Fig. 1b. Examination of the post-annealing film in bright-field mode was in agreement with the HAADF imaging; the layer structure was mostly preserved during the annealing process (Fig. 1c). Parallel
Fig. 2. Cross-sectional characterization of the Ni–10 vol% NiO sample before annealing using HAADF-STEM imaging (a) and EDS counts from a line scan taken across the Ni–NiO. The different layers, Pt, Ni–NiO, and SiO2 are labeled for convenience and the location and direction of the line scan is indicated by an arrow.
Fig. 3. Bright-field TEM image of the pre-annealed pure Ni sample with accompanying selected area diffraction patterns. Pattern (b) corresponds to the bright-field image shown in (a). Patterns (c) and (d) show other areas of the thin film which exhibit strong {001} and more randomized texture, respectively.
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Fig. 4. Bright-field TEM images of the samples containing 10% and 20% NiO with corresponding diffraction patterns (a and b, and c and d, respectively). Examples of ridges on the sample surface, caused by imperfections on the base salt crystal surface, are indicated by arrowheads in (a).
Moire´ fringes extending across the layers, however, suggest that some grain growth occurred through local rupturing of the NiO layers. No multilayer structure was evident in the cross-sectional examination of the Ni–10 vol% NiO material (Fig. 2). The light-todark contrast from the edges to the center of the Ni–NiO suggests that significant Pt and Si diffusion into the material occurred prior to annealing, but no evidence of discrete NiO layers was seen. An EDS line scan showed significant Pt penetration approximately 40 nm either from diffusion or redeposition during ion milling. Similarly, the presence of Si was detected 20 nm from the edge of the SiO2 layer. Pre-annealing characterization of the Ni films in plan view showed the initial microstructure to have a distribution of grain sizes ranging from less than 4 nm at the smallest to approximately 80 nm at the largest (Figs. 3 and 4). The average grain size in the pure Ni material was found from the micrographs to be approximately 14 nm. For the material containing NiO, the initial grain size was found to be approximately 8 nm for both the 10% NiO and 20% NiO materials. Due to multiple through-thickness grains and dynamic contrast effects, an accurate full grain size distribution of the unannealed films could not be obtained. Selected area diffraction patterns show strong variations in the initial texture between the different samples. The pure Ni samples contained areas with a dominant {001} texture (Fig. 1c), as well as areas of a few mixed orientations (Fig. 3b) and random texture as indicated by the solid rings making up the diffraction pattern (Fig. 3d). These variations in initial texture could result in strong
variations in growth behavior at different sample locations. The diffraction patterns from the samples with NiO added were found to be solid rings with a weak {001} texture apparent by the areas of stronger diffraction (Fig. 4). Multiple through-thickness grains prohibited the application of t-EBSD. 3.2. In situ annealing In the pure Ni sample annealed at 623 K, grain growth initiated in a few scattered grains (Fig. 5). These grains grew sporadically, rapidly reaching a size many times larger than the grains in the surrounding nanocrystalline matrix. This is especially apparent between Fig. 5b and c as the area of a single grain grew to many times the size of the grains in the surrounding matrix. The initial distribution of the abnormally growing grains was sparse, and they ceased growing before there was significant impingement between different growing grains. The addition of NiO influenced the grain growth, with no growth being observed for the 20% NiO sample and only limited growth for the 10% NiO sample (Fig. 6). Instead of a few grains growing large at the expense of the surrounding nanograins, the microstructure underwent a more general coarsening with several regions or pockets of grains growing to a few times their original size. This growth was somewhat more prevalent around preexisting ridges. Other regions of the sample remained unchanged. In pure Ni, further grain growth was observed at an annealing temperature of 723 K, including the expansion of grains whose
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Fig. 5. Snapshots taken from video showing the microstructure evolution in the same region of pure Ni on annealing at 623 K. Images were taken at 1 min intervals beginning from when the stage temperature was increased to 623 K.
Fig. 6. Snapshots taken from video showing the microstructure evolution in the same region of Ni–10% NiO on annealing at 623 K. Images were taken at 1 min intervals beginning from when the stage temperature was increased to 623 K.
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Fig. 7. Snapshots taken from video showing the microstructure evolution in the same region of pure Ni on annealing at 723 K. Images were taken at 1 min intervals beginning from when the stage temperature was increased to 723 K. Areas of further expansion of grains which were already growing at 623 K (A) and initiation of abnormal grain growth in a previously stable area (B) are labeled in (b). Examples of annealing twins are indicated by arrowheads in (e).
Fig. 8. Snapshots taken from video showing microstructure evolution of Ni–10% NiO on annealing at 723 K. Images were taken at 1 min intervals beginning from when the stage temperature was increased to 723 K.
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growth initiated at 623 K, as well as new growth in areas that were stable at the lower annealing temperatures. Examples of both types of growth are indicated in the time series of images presented in Fig. 7. The grain growth continued to be abnormal in nature; that is, it was driven by the rapid growth of a few grains instead of a general coarsening of the structure. The grain boundaries contain areas that are convex and others that are concave, with the growth seemingly dictated by a path of least resistance to boundary migration rather than driven by minimization of boundary area. Also noteworthy is the appearance of planar defects in the grain interiors (indicated by arrows in Fig. 7). Though not characterized in the TEM, the EBSD data presented later (Fig. 13) suggest them to be nanotwins. These nanotwins initiated at an early stage of the grain growth, extended across the grains, and grew at the same rate as the grain boundary progression. The small initial grain size prevented the identification of the precise moment of twin nucleation. At 723 K, the growth behavior of the 10% NiO material changed from a general microstructure coarsening to more abnormal growth (Fig. 8). In comparison to pure Ni, this abnormal grain growth was less pronounced and the abnormally growing grains attained a smaller final grain size. Grain growth in the 20% NiO material at 723 K resembled that in the 10% NiO material at 623 K (Fig. 9). Rather than a few grains dominating the growth, a more general coarsening of the microstructure was seen with higher growth areas concentrated around ridges in the sample. Only a small number of grains grew significantly larger than the surrounding matrix. No nanotwin formation was observed, though this may simply be due to the difficulty of resolving such features in the small grains. The growth behavior of a select grain in each of the three systems is directly compared in Fig. 10; here the same grain has
Fig. 10. Largest grain from each sample taken at 1 min intervals during annealing at 723 K showing how the abnormal grain growth varies with varying amounts of NiO. Notice that as the NiO content decreases the grains grow more rapidly and have a less equiaxed grain morphology.
Fig. 11. Average grain size as a function of annealing time and temperature measured from the TEM micrographs. Annealing time is given as time annealed at the given temperature (either 623 K or 723 K). The vertical bar indicates the transition from 623 K annealing temperature to 723 K.
Fig. 9. Snapshots taken from video showing the evolution of the microstructure in Ni–20% NiO on annealing at 723 K. Images are taken at 1 min intervals beginning from when the stage temperature was increased to 723 K.
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been extracted during its evolution with time at an annealing temperature of 723 K. As can be seen, with decreasing NiO content the grains attain a larger final size, grow more rapidly, and have more complex grain morphology. The average grain size as a function of temperature was measured from the TEM micrographs and is displayed in Fig. 11. The grain growth occurs in a relatively short period of time with the average grain size quickly reaching a plateau and not increasing further until the temperature is increased. 3.3. Post-annealing characterization The t-EBSD data show a distinct change in texture development with addition of NiO as evident by the pole figures and inverse pole figure (IPF) maps constructed from the orientation data (Fig. 12). In pure Ni, there are two dominant fiber textures,
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the {001} and the {122}. The {122} oriented grains are confined to four distinct orientations, corresponding to the four twin variants from the {001} orientation. In contrast, with the addition of NiO, the texture is composed of three dominant orientation variants: {001}, {122} (again, confined to the twin variants from the {001} direction), and {111} orientations. The {122} component of the texture, and the texture in general, becomes significantly weaker with the addition of NiO, suggesting that the grain growth transitions from a twin nucleation and propagation mode to a more homogeneous nucleation of energetically favorable grains. Grain boundary characterization similarly shows a steady decrease in the total length fraction of S3 grain boundaries from 0.37 in pure Ni to 0.26 and 0.22 with the addition of 10% and 20% NiO, respectively. This is consistent with NiO causing a decrease in the formation of annealing twins. These twin boundaries predominately exist between {001} and {122} oriented grains,
Fig. 12. Comparison of the IPF maps and corresponding pole figures after annealing at 723 K. Twin boundaries are highlighted as white. a and b, Ni; c and d, Ni–10% NiO; and e and f, Ni–20% NiO. Sample drift during the scan led to some grain elongation in the IPF maps. The orientation legend is included in the upper right corner of the image. (For interpretation of the references to color in this figure caption, the reader is referred to the web version of this article.)
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but to a lesser extent twinning also occurs between near {111} oriented grains; an example of such a twin is presented in Fig. 13, which is a select region of Fig. 12a. Besides orientation information, t-EBSD allows direct quantification of the grain size distribution in the annealed films. Fig. 14 graphically compares the grain size distributions of the three different samples after annealing at 723 K; the mean and standard deviation of the grain size are reported in Table 2. As the t-EBSD data do not include the smallest grains, the average grain size measured using the line intercept method from bright-field TEM micrographs is also reported in Table 1. The mean grain size decreases with increasing NiO content, which is consistent with the in situ annealing results. The standard deviation of the grain size also decreases with increasing NiO content, indicating more homogeneous growth. Again, this supports the observation that increasing the NiO content in the film leads to a transition from abnormal grain growth dominating the grain size increase to a more general coarsening of the microstructure. From Fig. 14 it is also seen that the maximum grain size in each sample increases with decreasing NiO content, with this effect being much more pronounced in the 20% NiO material where the largest grain in the
Table 2 Average grain size as measured by both t-EBSD and from TEM micrographs using the line intercept method after annealing at 723 K. Differences in the measured grain sizes are largely due to the smallest grains not being detected by t-EBSD. Sample
Mean grain size – t-EBSD (lm)
Mean grain size – TEM (lm)
Pure Ni 10% NiO 20% NiO
0.115 7 0.034 0.100 70.03 0.069 7 0.02
0.088 0.035 0.031
characterized area is 0.19 mm in diameter, compared to 0.35 mm and 0.39 mm in the 10% NiO and pure Ni samples, respectively. Partitioning of the orientation data into grain size regimes allows inspection of variations in texture with grain size; this information is presented in Fig. 15. As is evident from the figure, the texture shows a slight weakening with decreasing grain size; this trend is more evident in the materials containing NiO. The same trends, however, exist in all grain regimes; the {001}, {111}, and {122} orientations, are present in all regimes.
4. Discussion
Fig. 13. An enlarged area taken from the IPF map of annealed pure Ni shown in Fig. 12a. Twin boundaries are highlighted as white. As can be seen, twin boundaries exist predominately between {001} and {122} orientations (A), but also can exist between two near {111} oriented grains (B). Also indicated in the figure by an arrowhead is an approximately 10 nm-wide twin (orange twin in blue grain) that would not be resolved by using traditional EBSD. The wavy nature of the grain boundaries is an artifact of overlapping grains. (For interpretation of the references to color in this figure caption, the reader is referred to the web version of this article.)
Fig. 14. Comparison of the fraction of the area occupied by grains of different sizes on annealing at 723 K.
With increasing NiO content, the microstructure stability of the Ni was shown to increase. This is likely due to both the increased density of solute atoms as well as to the constraining effect of a multilayer structure. While the ability of solute atoms to slow or arrest grain growth is well documented [42,48], grain growth behavior in metal–metal/oxide multilayer structures has not been equally well investigated. Herweg et al. reported that at a temperature of 823 K, Ni–NiO nanolayer structures can become unstable and disintegrate into a Ni matrix with dispersed spherical NiO particles. In their study, the layer thickness varied between 2 and 7 nm for the Ni and NiO, respectively. The study presented here has shown that at temperatures 100 K lower, local rupturing of the multilayers occurred, likely as a precursor to the more general structure disintegration observed by Herweg et al. This local rupturing allowed grains to grow beyond the constraints imposed by layer spacing, but abnormal grain growth was still limited in comparison to pure Ni. The texture of the samples has been shown to depend strongly on the NiO content. In all samples there is a strong {001} texture, but as NiO content increases, the second dominant orientation transitions from a {122}, corresponding to the four twin variants of the {001}, to a {111} with the overall texture becoming more randomized. This texture development for each composition was independent of grain size regime except for a tendency for more random texture in the smallest grains. Thermally activated energy minimization processes in FCC structured thin films have been shown previously to drive the microstructure towards the {001} and the {111} orientations [49]. These texture developments are motivated by strain energy density and surface/interface energy reduction, respectively. Factors influencing which orientation develops in the texture include the elastic constants and the internal stress state of the material, and the temperature at which grain growth or recrystallization occurs. Thompson et al. showed that for metals with an anisotropy ratio greater than one (2C 44 =ðC 11 C 12 Þ 41, where C is the stiffness matrix), the strain energy and the surface/interface energy should play a role in determining the texture development during annealing [49]. As Ni has a moderate anisotropy ratio, 2.5 [50], it is not surprising that the evolved texture involves two orientations. The {001} texture develops during the deposition of the thin films, with this effect being more apparent in the pure Ni. As the films are annealed, surface/interface energy reduction is the dominant driving force for grain growth, leading to the formation
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Fig. 15. Pole figures constructed from orientations taken from different grain size regimes of the three materials. The grain size regime is given by the range in grain diameter included above each column of pole figures.
of annealing twins, which transform the texture to a {111} orientation. This suggests that the final texture seen in pure Ni is due to a favorable {001} orientation developed during the deposition of the films with annealing twin nucleation driving the transition towards a {122} orientation. There is also growth of energetically favorable {111} oriented grains during the annealing process. The observed texture development of the pure Ni is similar to that reported by Kim et al. [31], but differs significantly from that reported by Klement et al. [11]. This is most likely due to the initial texture, and highlights the important role initial conditions play in grain growth evolution. Klement et al. began the annealing experiments with a randomly oriented microstructure, which developed into a strong {111} fiber texture. The nanocrystalline Ni films annealed by Kim et al. began with a strong {001} texture and developed into a dual fiber texture of {001} and {111} oriented grains, which is consistent with the results presented herein. Chen et al. investigated the effects of substrate temperature on the texture of sputtered NiO thin films [51]. They found that films deposited at room temperature had a preferential {111} texture, while those deposited at 623 K developed a {001} texture, suggesting that the presence of the NiO layers could be an additional driving force for the {001} texture seen at higher temperatures in the multilayer films. The results clearly show that the formation of annealing twins and development of a {122} texture is suppressed by the addition of NiO. A prevalent theory on the formation of annealing twins is that they arise from growth accidents during the migration of grain boundaries [52,53]. Stacking errors in the lattice lead to the formation of twin boundaries when energetically favorable conditions arise. The frequency of these growth errors is thought to be linked to the migration velocity of the boundary, and so, at higher migration rates, the density of twin boundaries should
increase. As the NiO acts to reduce grain boundary mobility, this could explain the suppression of annealing twin formation, limiting the microstructure to growth of already favorably oriented grains instead of twin nucleation to reduce the system energy. A similar reduction in twin density was observed by Mahajan et al. [48]. In their samples, this reduction was from the addition of boron to Ni prior to annealing and grain growth, but similar mechanisms appear active in both cases. Computational and theoretical studies have suggested that solute atoms can increase the likelihood of abnormal grain growth due to preferential pinning of certain boundaries [18,54]. The results presented in this study suggest the opposite effect that the addition of an impurity element suppresses abnormal grain growth, delaying the onset and reducing the size disparity between abnormally growing grains and the surrounding nanocrystalline matrix. This reduction in abnormal grain growth was observed both in the 10% NiO film, where the NiO was likely dispersed throughout the Ni matrix, as well as in the 20% NiO film, where a clear multilayer structure was formed.
5. Conclusion Pure Ni with an initial nanocrystalline matrix develops a texture with dominant orientations of {001} and {122} upon annealing. The {122} texture develops as a result of twinning of the {001} grains. During annealing, grain growth proceeds in an abnormal manner, with a few grains growing quickly at the expense of the surrounding nanocrystalline matrix. Cross-sectional examination of Ni–NiO thin films showed that it is possible to synthesize nanolayered Ni–NiO structures with nominal alternating layer thicknesses of 5 and 1 nm. Local disruptions of the layer structure occurred at a temperature of
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723 K. When the NiO layer thickness was reduced to 0.5 nm, no coherent multilayer structure was observed. The addition of NiO to nanocrystalline Ni acts as a stabilizing influence against thermally activated grain growth. With increasing NiO content, grain growth initiates at higher temperatures and the maximum grain size reached decreases. The abnormal grain growth seen in pure Ni is somewhat suppressed as grain growth occurs more homogenously throughout with the variance in grain size decreasing with increasing NiO content. Postannealing texture is also affected by the presence of NiO as it changes from a dual fiber texture of {001} and {122} to a texture composed of {001} and {111} orientations with a somewhat weaker presence of {122} oriented grains.
Acknowledgments The assistance with FIB machining by May Martin is gratefully acknowledged. The work at the University of Illinois (JK, SH, PE, and IMR) was supported by the US Department of Energy Office of Basic Energy Sciences, Division of Materials Science, under award No. DEFG-02-07ER46443. Sandia National Laboratories is a multiprogram laboratory managed and operated by Sandia Corporation, a wholly owned subsidiary of Lockheed Martin Corporation for the U.S. Department of Energy’s National Nuclear Security Administration under contract DE-AC04-94AL85000. A portion of this research was carried out in the Frederick Seitz Materials Research Laboratory Central Facilities, University of Illinois. KH, JK, and PE acknowledge support from the Division of Materials Science and Engineering, Office of Basic Energy Sciences, U.S. Department of Energy at Sandia. IMR acknowledges the support of the National Science Foundation. References [1] M.A. Meyers, A. Mishra, D.J. Benson, Prog. Mater. Sci. 51 (2006) 427–556. [2] L.N. Brewer, D.M. Follstaedt, K. Hattar, J.A. Knapp, M.A. Rodriguez, I.M. Robertson, Adv. Mater. 22 (2010) 1161–1164. [3] D.M. Follstaedt, K. Hattar, J.A. Knapp, I.M. Robertson, Mater. Res. Soc. Symp. Proc. 907E (2006). [4] K. Hattar, D.M. Follstaedt, J.A. Knapp, I.M. Robertson, Acta Mater. 56 (2008) 794–801. [5] M. Ames, J. Markmann, R. Karos, A. Michels, A. Tschope, R. Birringer, Acta Mater. 56 (2008) 4255–4266. [6] G.D. Hibbard, K.T. Aust, U. Erb, J. Mater. Sci. 43 (2008) 6441–6452. [7] G.D. Hibbard, V. Radmilovic, K.T. Aust, U. Erb, Mater. Sci. Eng. A 494 (2008) 232–238. [8] K. Bae-Kyun, J.A. Szpunar, R. Varano, Mater. Sci. Forum. 408–412 (2002) 937–942. [9] M.C. Iordache, S.H. Whang, Z. Jiao, Z.M. Wang, Nanostruct. Mater. 11 (1999) 1343–1349. [10] V.Y. Novikov, Interface Sci. 6 (1998) 77–84. [11] U. Klement, M. Da Silva, W. Skrotzki, J. Microsc. (2008) 455–463. [12] A.D. Rollett, W.W. Mullins, Scr. Mater. 36 (1997) 975–980. [13] A.D. Rollett, D.J. Srolovitz, M.P. Anderson, Acta Metall. 37 (1989) 1227–1240.
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Contents lists available at SciVerse ScienceDirect
Materials Science & Engineering A journal homepage: www.elsevier.com/locate/msea
Thermal stability of Ni/NiO multilayers Josh Kacher a,n, Patricia Elizaga a, Stephen D. House a, Khalid Hattar b, Matt Nowell c, I.M. Robertson a a b c
Materials Science and Engineering, University of Illinois at Urbana-Champaign, 1304 W Green St, Urbana, IL 61801, United States Sandia National Laboratories, Albuquerque, NM 87185, United States EDAX/TSL 392 East 12300 South, Draper, UT 84020, United States
a r t i c l e i n f o
abstract
Article history: Received 5 November 2012 Received in revised form 14 January 2013 Accepted 17 January 2013 Available online 23 January 2013
The effects of NiO multilayer additions to pulsed-laser deposited nanocrystalline Ni on the grain growth and texture evolution during annealing have been studied using a combination of in situ annealing in the transmission electron microscope and the recently developed transmission electron backscatter diffraction technique. Grain growth in pure Ni proceeds in an abnormal manner with a small number of grains growing rapidly at the expense of the surrounding nanocrystalline matrix. The addition of NiO layers in the Ni suppressed grain growth and twin formation upon annealing, with increasing oxide content resulting in decreasing extent of grain growth. In pure Ni, the post-annealing texture was found to be composed of a strong {001} texture coupled with the four {122} twin variants from the {001}. With the addition of NiO, annealing twinning, and hence the {122} orientation, was suppressed and instead the final texture was composed of the {001} and {111} orientations. & 2013 Elsevier B.V. All rights reserved.
Keywords: EBSD Electron microscopy Nanostructured materials Grain growth
1. Introduction Nanocrystalline metals have been garnering increased attention due to their unique properties such as exceptionally high strength [1], but microstructural instability at relatively low homologous temperatures limits their use. The high grain boundary energy density creates a high driving force for grain growth, both normal and abnormal, at temperatures as low as 600 K in various materials [2–9]. Abnormal grain growth has been attributed to factors such as the initial grain size distribution [10], orientation distribution [8,11–13], grain boundary characteristics [14–16], triple junction mobility [17], impurities in the form of solute atoms [18] and second phase particles [19], heterogeneous microstrains [5], and topological features such as grain boundary grooving [20]. Multilayered films and nanostructured metal-oxides are also of considerable interest due to their unique magnetic and electronic properties [21–25]. Efforts specifically targeted at the synthesis of Ni–NiO multilayers have included the periodic oxidation of metallic layers using pulsed oxygen molecular beams during deposition [26,27] and evaporation of NiO with an electron beam [26], with single layer thicknesses ranging from 2 to 10 nm. Studies on the thermal stability of Ni–NiO multilayer structures have shown that annealing can cause the multilayer structure to disintegrate, leaving dispersed spherical oxide particles in the metallic matrix. This can affect the electronic properties of the
n
Corresponding author. Tel.: þ1 317 835 5274. E-mail addresses: [email protected], [email protected] (J. Kacher).
0921-5093/$ - see front matter & 2013 Elsevier B.V. All rights reserved. http://dx.doi.org/10.1016/j.msea.2013.01.033
material, such as the loss of the angular dependence of magnetoresistance [24]. To investigate the reactivity between Ni and NiO, Symianakis et al. deposited Ni at near monolayer thickness onto a (001) oriented NiO substrate [28]. They found that the two systems were stable at room temperature, but on annealing the Ni layer oxidized, eventually causing all of the Ni to convert to NiO. The onset of this effect was at approximately 773 K with the oxygen diffusion accelerating with increasing temperature. Experimental studies of microstructure stability and grain growth in nanocrystalline and nanostructured materials have primarily utilized four characterization techniques: transmission electron microscopy (TEM) [3,4,6,7,29,30], x-ray diffraction [24,31], differential scanning calorimetry [32,33], and scanning electron microscopy (SEM) based electron backscatter diffraction (EBSD) [30,31]. By annealing nanocrystalline Ni films in situ in the TEM, the dynamics of grain growth also can be observed and the formation of an oxide layer during the annealing can be minimized. For example, studies have shown that grain growth in nanocrystalline materials often proceeds in an abnormal fashion, with a relatively small number of grains growing quickly at the expense of the surrounding nanocrystalline matrix [4,6,7,11]. This growth proceeds in a sporadic fashion; often beginning at one grain boundary, stopping, and continuing from a different location on the same boundary. The resulting grains have been observed to have an unusual morphology, with both convex and concave regions coexisting along the grain boundary. Such observations suggest that in-plane curvature driven growth is not the determining factor governing grain expansion [30]. Local velocity measurements have shown that individual growth fronts propagate
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at rates many times higher than the average growth rate would dictate [6]. Post mortem analysis following annealing showed a high density of vacancy-based defects in the grain interiors, including stacking-fault tetrahedra [4]. This was interpreted as being from the absorption and subsequent ejection of the excess free volume found in grain boundaries in PLD nanocrystalline nickel, highlighting vacancy diffusion as a potentially important mechanism dictating grain growth. Solute atoms are known to play an important role in stabilizing the microstructure, although the mechanisms for how this is accomplished are still being debated. The two prevailing theories essentially follow either a kinetic or a thermodynamic argument. The kinetic, or solute drag, approach, first developed by Cahn [34], focuses on the drag force that solute atoms can exert on grain boundaries during grain growth. Solute atoms occupy the free volume in the boundary and lower the grain boundary energy, thus reducing the driving force for grain growth. As the grains grow, the grain boundary area of individual grains proportionally increases and more solute atoms are accommodated in the boundary. In nanocrystalline materials, grain boundaries can propagate at high velocity and the grain boundary area of the individual grains increases rapidly, which along with the concurrent increase in solute content leads to complete boundary pinning and halts grain growth [35,36]. Solute drag theory, for example, has been used to explain the buildup of sulfur at grain boundaries after grain growth in nanocrystalline Ni [6]. Similarly, da Silva et al. found that additions of phosphorous to nanocrystalline cobalt improved microstructural stability, resulting in a delayed onset of grain growth and more homogeneous growth once it initiated. Using atom probe tomography, they showed that this reduced grain growth was linked to phosphorous segregation to the grain boundaries [33]. This knowledge has motivated the addition of sulfur and phosphorous during electrodeposition to stabilize the fine grain structure [37,38]. The thermodynamic approach considers the Gibbs free energy of the complete system [39–42]. Solute atom segregation to grain boundaries can lower the Gibbs free energy of the system, below that achievable through elimination of grain boundaries. Thus, there exists a solute concentration dependent metastable grain boundary area for the system, which once attained, removes the impetus for further grain growth. Thermodynamic arguments therefore predict that, as more solute atoms must be accommodated, the equilibrium grain size will decrease with increasing solute concentration. EBSD, x-ray diffraction, and selected area electron diffraction have provided useful tools in examining the texture evolution during grain growth. Most studies of FCC materials agree that a strong {111} fiber texture develops during the annealing process, but with some disagreement about the initial texture. Klement et al. reported an initial {114} texture at the onset of grain growth in both Ni and Ni-20 at % Fe, with twinning transitioning the texture to {111} [11]. Kim et al. reported an initial texture dominated by {001} in nanocrystalline Ni, with the {111} texture growing in strength at the expense of the {001} oriented grains with increasing annealing temperature [31]. However, the EBSD analysis in these studies was conducted at a minimum step size of 15 nm, excluding information about the smallest grains. Consequently, much of the grain boundary information was not obtained and the analysis was limited to only those grains that had reached a certain stage of grain growth. Recently, Rajasekhara et al. demonstrated that the initial texture of nanocrystalline PLD Ni thin films could be characterized using precession microscopy in a TEM, which allows orientation mapping with a minimum resolvable grain size of 4 nm [43]. The improved spatial resolution allowed them to identify a significant fraction of HCP-phase Ni even in the as-deposited
state. Keller and Geiss developed a technique similar to EBSD that is based on the collection of transmitted electrons in an SEM to form a diffraction pattern and is referred to as transmission EBSD or t-EBSD [44]. By reducing the thickness of the sample to a few tens to hundreds of nanometers, similar to the thickness of a TEM sample, and collecting only the transmitted electrons, the interaction volume becomes a narrow column instead of the larger teardrop shape from which backscattered electrons are collected. This can improve the spatial resolution to below 10 nm, compared to the effective resolution of traditional EBSD which, while dependent on multiple factors, is worse by at least a factor of 2 [45]. In this manner, orientation mapping can be done in an SEM at a resolution approaching that which is possible using precession microscopy in a TEM. The main limitation with using t-EBSD for nanocrystalline materials arises from the presence of multiple through-thickness grains, leading to mixed diffraction patterns and preventing reliable pattern indexing. In this study, nanocrystalline Ni–NiO multilayer thin films were deposited with varying thicknesses of the NiO layers. The thermal stability and grain growth characteristics of these films were investigated through TEM characterization, including in situ annealing, focused ion beam (FIB) lift-out for cross-sectional analysis, and post mortem t-EBSD investigations of the orientation and grain boundary distributions.
2. Experimental Nanocrystalline Ni samples were made using pulsed-laser deposition at Sandia National Laboratory [46]. The samples were deposited on a base of (001) oriented rock salt with a beam energy of 500 mJ at room temperature to a final thickness of approximately 70 nm. The growth chamber achieved a base pressure of approximately 5 10 7 Torr. Three sample compositions, Ni, nominally Ni–10 vol% NiO, and nominally Ni–20 vol% NiO, were synthesized to investigate the effects of sample composition and multilayer structure on the thermal stability. Pure Ni (99.97%) and pure NiO (99.9% ) targets were used during the deposition. Pure Ni was deposited at a nominal laser pulse frequency of 35 Hz until the desired thickness was achieved. Ni and NiO were deposited in alternating layers, with the Ni deposited at a pulse frequency of 35 Hz and the NiO at 7 Hz. For the 10 vol% NiO material, the layer thickness alternated between nominally 4.8 nm for pure Ni and 0.5 nm for the NiO. This was changed to nominally 4.5 nm for the pure Ni and 1.1 nm for the NiO in the 20 vol% NiO material. The volume percent of the NiO is a nominal value based on calibrated deposition rates. The sample deposition conditions are summarized in Table 1. Using the same conditions, Ni–NiO multilayers were also deposited on an amorphous SiO2 substrate to facilitate cross-sectional TEM examination of samples extracted using FIB machining. Plan-view samples were prepared for analysis in the TEM and SEM by scoring a 2 mm grid on the surface of the film and placing the salt crystal in deionized water. Upon dissolution of the salt, the samples floated to the surface and were collected using 3 mmdiameter clamping copper grids. Cross-section samples, to investigate the multilayer Ni–NiO structure, were prepared using an FEI Dual Table 1 Summary of materials investigated. Note that all layer thicknesses are nominal. Material
Ni layer thickness (nm)
NiO layer thickness (nm)
Pure Ni Ni–10 vol% NiO Ni–20 vol% NiO
4.8 4.5
0.5 1.1
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Beam 235 FIB and standard FIB lift-out and thinning techniques (see, for example [47]). A platinum layer was deposited prior to any ion milling to ensure the preservation of the underlying microstructure. Grain growth in the plan-view thin films was observed during in situ annealing in a Phillips CM30 LaB6 TEM. Samples were heated to 473 K and allowed to stabilize. Next they were heated to 623 K and held at that temperature for 5–7 min and, finally, the temperature was raised to 723 K. All stages of heating were done at a ramp rate of approximately 5 K/s. As no grain growth was observed at 473 K, the evolution of the microstructure was recorded only during the latter two stages of annealing. The grain growth was captured using a CCD camera at a frame rate of three frames per second for at least 5 min. This was found to be ample time for the microstructure to reach a state beyond which no noticeable changes were observed. Pre- and post-annealing characterization of the films was carried out using TEM analysis with the microscope operating at 300 kV. Images using multiple different diffraction vectors were collected to better resolve the grain structure. T-EBSD was used to study the texture and microstructure of the plan-view samples after the annealing was complete. Analysis was performed using EDAX/TSL orientation imaging microscopy (OIM) data collection software at an accelerating voltage of 30 kV.
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The samples were held at 51 from horizontal and at a working distance of approximately 4 mm using a custom built stage that allowed for transmitted electrons to be detected on the phosphor screen. Two 2 2 mm scans were collected from each sample at a step size of 5 nm and a scanning speed of 200 points per second for a total scan time of approximately 800 s. The collected data were used to construct inverse pole figure (IPF) maps and pole figures. No data interpolation was used and all points indexed with a confidence level below 0.1, as defined by the OIM software, were filtered from the data; these are displayed as black areas in the IPF maps. Both FCC- and HCPphase Ni were included in the search, but no HCP-phase Ni grains were reliably identified. The cross-sectional samples were characterized both before and after annealing using high angle annular dark-field scanning TEM (HAADF-STEM) imaging and energy-dispersive x-ray spectroscopy (EDS) line scans in a JEOL 2010 F (S)TEM operated at 200 kV. Annealing of the FIB lift-out samples was performed in situ in a JEOL 2010 LaB6 TEM operating at 200 kV after the FIB lift-out. As thermal stresses could cause the separation of the FIB lift-out sample from the copper grid, the heating rate was reduced to approximately 1 K/s.
Fig. 1. Cross-sectional characterization of the Ni–20 vol% NiO sample before (a) and after (b) annealing using HAADF-STEM imaging and post-annealing characterization using bright-field TEM imaging (c). The vacuum/SiO2 interface is delineated in (b) by a dotted line and different layers, Pt, Ni–NiO, and SiO2 are labeled for convenience. An inset with contrast levels adjusted and a dotted line added to highlight the Moire´ fringes is included in (c).
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3. Results 3.1. Pre-annealing characterization Cross-sectional examination of the Ni–20 vol% NiO material using HAADF-STEM imaging showed that the multilayer structure remained coherent during the deposition and sample preparation processes (Fig. 1a). The periodicity as well as the layer thickness
was preserved with no evidence of local rupturing in the layer structure. Post-annealing characterization showed the multilayer structure to still be largely unchanged. Annealing the films caused the separation of the Ni–NiO from the SiO2, evident by the black contrast layer between the two materials seen in Fig. 1b. Examination of the post-annealing film in bright-field mode was in agreement with the HAADF imaging; the layer structure was mostly preserved during the annealing process (Fig. 1c). Parallel
Fig. 2. Cross-sectional characterization of the Ni–10 vol% NiO sample before annealing using HAADF-STEM imaging (a) and EDS counts from a line scan taken across the Ni–NiO. The different layers, Pt, Ni–NiO, and SiO2 are labeled for convenience and the location and direction of the line scan is indicated by an arrow.
Fig. 3. Bright-field TEM image of the pre-annealed pure Ni sample with accompanying selected area diffraction patterns. Pattern (b) corresponds to the bright-field image shown in (a). Patterns (c) and (d) show other areas of the thin film which exhibit strong {001} and more randomized texture, respectively.
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Fig. 4. Bright-field TEM images of the samples containing 10% and 20% NiO with corresponding diffraction patterns (a and b, and c and d, respectively). Examples of ridges on the sample surface, caused by imperfections on the base salt crystal surface, are indicated by arrowheads in (a).
Moire´ fringes extending across the layers, however, suggest that some grain growth occurred through local rupturing of the NiO layers. No multilayer structure was evident in the cross-sectional examination of the Ni–10 vol% NiO material (Fig. 2). The light-todark contrast from the edges to the center of the Ni–NiO suggests that significant Pt and Si diffusion into the material occurred prior to annealing, but no evidence of discrete NiO layers was seen. An EDS line scan showed significant Pt penetration approximately 40 nm either from diffusion or redeposition during ion milling. Similarly, the presence of Si was detected 20 nm from the edge of the SiO2 layer. Pre-annealing characterization of the Ni films in plan view showed the initial microstructure to have a distribution of grain sizes ranging from less than 4 nm at the smallest to approximately 80 nm at the largest (Figs. 3 and 4). The average grain size in the pure Ni material was found from the micrographs to be approximately 14 nm. For the material containing NiO, the initial grain size was found to be approximately 8 nm for both the 10% NiO and 20% NiO materials. Due to multiple through-thickness grains and dynamic contrast effects, an accurate full grain size distribution of the unannealed films could not be obtained. Selected area diffraction patterns show strong variations in the initial texture between the different samples. The pure Ni samples contained areas with a dominant {001} texture (Fig. 1c), as well as areas of a few mixed orientations (Fig. 3b) and random texture as indicated by the solid rings making up the diffraction pattern (Fig. 3d). These variations in initial texture could result in strong
variations in growth behavior at different sample locations. The diffraction patterns from the samples with NiO added were found to be solid rings with a weak {001} texture apparent by the areas of stronger diffraction (Fig. 4). Multiple through-thickness grains prohibited the application of t-EBSD. 3.2. In situ annealing In the pure Ni sample annealed at 623 K, grain growth initiated in a few scattered grains (Fig. 5). These grains grew sporadically, rapidly reaching a size many times larger than the grains in the surrounding nanocrystalline matrix. This is especially apparent between Fig. 5b and c as the area of a single grain grew to many times the size of the grains in the surrounding matrix. The initial distribution of the abnormally growing grains was sparse, and they ceased growing before there was significant impingement between different growing grains. The addition of NiO influenced the grain growth, with no growth being observed for the 20% NiO sample and only limited growth for the 10% NiO sample (Fig. 6). Instead of a few grains growing large at the expense of the surrounding nanograins, the microstructure underwent a more general coarsening with several regions or pockets of grains growing to a few times their original size. This growth was somewhat more prevalent around preexisting ridges. Other regions of the sample remained unchanged. In pure Ni, further grain growth was observed at an annealing temperature of 723 K, including the expansion of grains whose
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Fig. 5. Snapshots taken from video showing the microstructure evolution in the same region of pure Ni on annealing at 623 K. Images were taken at 1 min intervals beginning from when the stage temperature was increased to 623 K.
Fig. 6. Snapshots taken from video showing the microstructure evolution in the same region of Ni–10% NiO on annealing at 623 K. Images were taken at 1 min intervals beginning from when the stage temperature was increased to 623 K.
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Fig. 7. Snapshots taken from video showing the microstructure evolution in the same region of pure Ni on annealing at 723 K. Images were taken at 1 min intervals beginning from when the stage temperature was increased to 723 K. Areas of further expansion of grains which were already growing at 623 K (A) and initiation of abnormal grain growth in a previously stable area (B) are labeled in (b). Examples of annealing twins are indicated by arrowheads in (e).
Fig. 8. Snapshots taken from video showing microstructure evolution of Ni–10% NiO on annealing at 723 K. Images were taken at 1 min intervals beginning from when the stage temperature was increased to 723 K.
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growth initiated at 623 K, as well as new growth in areas that were stable at the lower annealing temperatures. Examples of both types of growth are indicated in the time series of images presented in Fig. 7. The grain growth continued to be abnormal in nature; that is, it was driven by the rapid growth of a few grains instead of a general coarsening of the structure. The grain boundaries contain areas that are convex and others that are concave, with the growth seemingly dictated by a path of least resistance to boundary migration rather than driven by minimization of boundary area. Also noteworthy is the appearance of planar defects in the grain interiors (indicated by arrows in Fig. 7). Though not characterized in the TEM, the EBSD data presented later (Fig. 13) suggest them to be nanotwins. These nanotwins initiated at an early stage of the grain growth, extended across the grains, and grew at the same rate as the grain boundary progression. The small initial grain size prevented the identification of the precise moment of twin nucleation. At 723 K, the growth behavior of the 10% NiO material changed from a general microstructure coarsening to more abnormal growth (Fig. 8). In comparison to pure Ni, this abnormal grain growth was less pronounced and the abnormally growing grains attained a smaller final grain size. Grain growth in the 20% NiO material at 723 K resembled that in the 10% NiO material at 623 K (Fig. 9). Rather than a few grains dominating the growth, a more general coarsening of the microstructure was seen with higher growth areas concentrated around ridges in the sample. Only a small number of grains grew significantly larger than the surrounding matrix. No nanotwin formation was observed, though this may simply be due to the difficulty of resolving such features in the small grains. The growth behavior of a select grain in each of the three systems is directly compared in Fig. 10; here the same grain has
Fig. 10. Largest grain from each sample taken at 1 min intervals during annealing at 723 K showing how the abnormal grain growth varies with varying amounts of NiO. Notice that as the NiO content decreases the grains grow more rapidly and have a less equiaxed grain morphology.
Fig. 11. Average grain size as a function of annealing time and temperature measured from the TEM micrographs. Annealing time is given as time annealed at the given temperature (either 623 K or 723 K). The vertical bar indicates the transition from 623 K annealing temperature to 723 K.
Fig. 9. Snapshots taken from video showing the evolution of the microstructure in Ni–20% NiO on annealing at 723 K. Images are taken at 1 min intervals beginning from when the stage temperature was increased to 723 K.
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been extracted during its evolution with time at an annealing temperature of 723 K. As can be seen, with decreasing NiO content the grains attain a larger final size, grow more rapidly, and have more complex grain morphology. The average grain size as a function of temperature was measured from the TEM micrographs and is displayed in Fig. 11. The grain growth occurs in a relatively short period of time with the average grain size quickly reaching a plateau and not increasing further until the temperature is increased. 3.3. Post-annealing characterization The t-EBSD data show a distinct change in texture development with addition of NiO as evident by the pole figures and inverse pole figure (IPF) maps constructed from the orientation data (Fig. 12). In pure Ni, there are two dominant fiber textures,
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the {001} and the {122}. The {122} oriented grains are confined to four distinct orientations, corresponding to the four twin variants from the {001} orientation. In contrast, with the addition of NiO, the texture is composed of three dominant orientation variants: {001}, {122} (again, confined to the twin variants from the {001} direction), and {111} orientations. The {122} component of the texture, and the texture in general, becomes significantly weaker with the addition of NiO, suggesting that the grain growth transitions from a twin nucleation and propagation mode to a more homogeneous nucleation of energetically favorable grains. Grain boundary characterization similarly shows a steady decrease in the total length fraction of S3 grain boundaries from 0.37 in pure Ni to 0.26 and 0.22 with the addition of 10% and 20% NiO, respectively. This is consistent with NiO causing a decrease in the formation of annealing twins. These twin boundaries predominately exist between {001} and {122} oriented grains,
Fig. 12. Comparison of the IPF maps and corresponding pole figures after annealing at 723 K. Twin boundaries are highlighted as white. a and b, Ni; c and d, Ni–10% NiO; and e and f, Ni–20% NiO. Sample drift during the scan led to some grain elongation in the IPF maps. The orientation legend is included in the upper right corner of the image. (For interpretation of the references to color in this figure caption, the reader is referred to the web version of this article.)
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but to a lesser extent twinning also occurs between near {111} oriented grains; an example of such a twin is presented in Fig. 13, which is a select region of Fig. 12a. Besides orientation information, t-EBSD allows direct quantification of the grain size distribution in the annealed films. Fig. 14 graphically compares the grain size distributions of the three different samples after annealing at 723 K; the mean and standard deviation of the grain size are reported in Table 2. As the t-EBSD data do not include the smallest grains, the average grain size measured using the line intercept method from bright-field TEM micrographs is also reported in Table 1. The mean grain size decreases with increasing NiO content, which is consistent with the in situ annealing results. The standard deviation of the grain size also decreases with increasing NiO content, indicating more homogeneous growth. Again, this supports the observation that increasing the NiO content in the film leads to a transition from abnormal grain growth dominating the grain size increase to a more general coarsening of the microstructure. From Fig. 14 it is also seen that the maximum grain size in each sample increases with decreasing NiO content, with this effect being much more pronounced in the 20% NiO material where the largest grain in the
Table 2 Average grain size as measured by both t-EBSD and from TEM micrographs using the line intercept method after annealing at 723 K. Differences in the measured grain sizes are largely due to the smallest grains not being detected by t-EBSD. Sample
Mean grain size – t-EBSD (lm)
Mean grain size – TEM (lm)
Pure Ni 10% NiO 20% NiO
0.115 7 0.034 0.100 70.03 0.069 7 0.02
0.088 0.035 0.031
characterized area is 0.19 mm in diameter, compared to 0.35 mm and 0.39 mm in the 10% NiO and pure Ni samples, respectively. Partitioning of the orientation data into grain size regimes allows inspection of variations in texture with grain size; this information is presented in Fig. 15. As is evident from the figure, the texture shows a slight weakening with decreasing grain size; this trend is more evident in the materials containing NiO. The same trends, however, exist in all grain regimes; the {001}, {111}, and {122} orientations, are present in all regimes.
4. Discussion
Fig. 13. An enlarged area taken from the IPF map of annealed pure Ni shown in Fig. 12a. Twin boundaries are highlighted as white. As can be seen, twin boundaries exist predominately between {001} and {122} orientations (A), but also can exist between two near {111} oriented grains (B). Also indicated in the figure by an arrowhead is an approximately 10 nm-wide twin (orange twin in blue grain) that would not be resolved by using traditional EBSD. The wavy nature of the grain boundaries is an artifact of overlapping grains. (For interpretation of the references to color in this figure caption, the reader is referred to the web version of this article.)
Fig. 14. Comparison of the fraction of the area occupied by grains of different sizes on annealing at 723 K.
With increasing NiO content, the microstructure stability of the Ni was shown to increase. This is likely due to both the increased density of solute atoms as well as to the constraining effect of a multilayer structure. While the ability of solute atoms to slow or arrest grain growth is well documented [42,48], grain growth behavior in metal–metal/oxide multilayer structures has not been equally well investigated. Herweg et al. reported that at a temperature of 823 K, Ni–NiO nanolayer structures can become unstable and disintegrate into a Ni matrix with dispersed spherical NiO particles. In their study, the layer thickness varied between 2 and 7 nm for the Ni and NiO, respectively. The study presented here has shown that at temperatures 100 K lower, local rupturing of the multilayers occurred, likely as a precursor to the more general structure disintegration observed by Herweg et al. This local rupturing allowed grains to grow beyond the constraints imposed by layer spacing, but abnormal grain growth was still limited in comparison to pure Ni. The texture of the samples has been shown to depend strongly on the NiO content. In all samples there is a strong {001} texture, but as NiO content increases, the second dominant orientation transitions from a {122}, corresponding to the four twin variants of the {001}, to a {111} with the overall texture becoming more randomized. This texture development for each composition was independent of grain size regime except for a tendency for more random texture in the smallest grains. Thermally activated energy minimization processes in FCC structured thin films have been shown previously to drive the microstructure towards the {001} and the {111} orientations [49]. These texture developments are motivated by strain energy density and surface/interface energy reduction, respectively. Factors influencing which orientation develops in the texture include the elastic constants and the internal stress state of the material, and the temperature at which grain growth or recrystallization occurs. Thompson et al. showed that for metals with an anisotropy ratio greater than one (2C 44 =ðC 11 C 12 Þ 41, where C is the stiffness matrix), the strain energy and the surface/interface energy should play a role in determining the texture development during annealing [49]. As Ni has a moderate anisotropy ratio, 2.5 [50], it is not surprising that the evolved texture involves two orientations. The {001} texture develops during the deposition of the thin films, with this effect being more apparent in the pure Ni. As the films are annealed, surface/interface energy reduction is the dominant driving force for grain growth, leading to the formation
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Fig. 15. Pole figures constructed from orientations taken from different grain size regimes of the three materials. The grain size regime is given by the range in grain diameter included above each column of pole figures.
of annealing twins, which transform the texture to a {111} orientation. This suggests that the final texture seen in pure Ni is due to a favorable {001} orientation developed during the deposition of the films with annealing twin nucleation driving the transition towards a {122} orientation. There is also growth of energetically favorable {111} oriented grains during the annealing process. The observed texture development of the pure Ni is similar to that reported by Kim et al. [31], but differs significantly from that reported by Klement et al. [11]. This is most likely due to the initial texture, and highlights the important role initial conditions play in grain growth evolution. Klement et al. began the annealing experiments with a randomly oriented microstructure, which developed into a strong {111} fiber texture. The nanocrystalline Ni films annealed by Kim et al. began with a strong {001} texture and developed into a dual fiber texture of {001} and {111} oriented grains, which is consistent with the results presented herein. Chen et al. investigated the effects of substrate temperature on the texture of sputtered NiO thin films [51]. They found that films deposited at room temperature had a preferential {111} texture, while those deposited at 623 K developed a {001} texture, suggesting that the presence of the NiO layers could be an additional driving force for the {001} texture seen at higher temperatures in the multilayer films. The results clearly show that the formation of annealing twins and development of a {122} texture is suppressed by the addition of NiO. A prevalent theory on the formation of annealing twins is that they arise from growth accidents during the migration of grain boundaries [52,53]. Stacking errors in the lattice lead to the formation of twin boundaries when energetically favorable conditions arise. The frequency of these growth errors is thought to be linked to the migration velocity of the boundary, and so, at higher migration rates, the density of twin boundaries should
increase. As the NiO acts to reduce grain boundary mobility, this could explain the suppression of annealing twin formation, limiting the microstructure to growth of already favorably oriented grains instead of twin nucleation to reduce the system energy. A similar reduction in twin density was observed by Mahajan et al. [48]. In their samples, this reduction was from the addition of boron to Ni prior to annealing and grain growth, but similar mechanisms appear active in both cases. Computational and theoretical studies have suggested that solute atoms can increase the likelihood of abnormal grain growth due to preferential pinning of certain boundaries [18,54]. The results presented in this study suggest the opposite effect that the addition of an impurity element suppresses abnormal grain growth, delaying the onset and reducing the size disparity between abnormally growing grains and the surrounding nanocrystalline matrix. This reduction in abnormal grain growth was observed both in the 10% NiO film, where the NiO was likely dispersed throughout the Ni matrix, as well as in the 20% NiO film, where a clear multilayer structure was formed.
5. Conclusion Pure Ni with an initial nanocrystalline matrix develops a texture with dominant orientations of {001} and {122} upon annealing. The {122} texture develops as a result of twinning of the {001} grains. During annealing, grain growth proceeds in an abnormal manner, with a few grains growing quickly at the expense of the surrounding nanocrystalline matrix. Cross-sectional examination of Ni–NiO thin films showed that it is possible to synthesize nanolayered Ni–NiO structures with nominal alternating layer thicknesses of 5 and 1 nm. Local disruptions of the layer structure occurred at a temperature of
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723 K. When the NiO layer thickness was reduced to 0.5 nm, no coherent multilayer structure was observed. The addition of NiO to nanocrystalline Ni acts as a stabilizing influence against thermally activated grain growth. With increasing NiO content, grain growth initiates at higher temperatures and the maximum grain size reached decreases. The abnormal grain growth seen in pure Ni is somewhat suppressed as grain growth occurs more homogenously throughout with the variance in grain size decreasing with increasing NiO content. Postannealing texture is also affected by the presence of NiO as it changes from a dual fiber texture of {001} and {122} to a texture composed of {001} and {111} orientations with a somewhat weaker presence of {122} oriented grains.
Acknowledgments The assistance with FIB machining by May Martin is gratefully acknowledged. The work at the University of Illinois (JK, SH, PE, and IMR) was supported by the US Department of Energy Office of Basic Energy Sciences, Division of Materials Science, under award No. DEFG-02-07ER46443. Sandia National Laboratories is a multiprogram laboratory managed and operated by Sandia Corporation, a wholly owned subsidiary of Lockheed Martin Corporation for the U.S. Department of Energy’s National Nuclear Security Administration under contract DE-AC04-94AL85000. A portion of this research was carried out in the Frederick Seitz Materials Research Laboratory Central Facilities, University of Illinois. KH, JK, and PE acknowledge support from the Division of Materials Science and Engineering, Office of Basic Energy Sciences, U.S. Department of Energy at Sandia. IMR acknowledges the support of the National Science Foundation. References [1] M.A. Meyers, A. Mishra, D.J. Benson, Prog. Mater. Sci. 51 (2006) 427–556. [2] L.N. Brewer, D.M. Follstaedt, K. Hattar, J.A. Knapp, M.A. Rodriguez, I.M. Robertson, Adv. Mater. 22 (2010) 1161–1164. [3] D.M. Follstaedt, K. Hattar, J.A. Knapp, I.M. Robertson, Mater. Res. Soc. Symp. Proc. 907E (2006). [4] K. Hattar, D.M. Follstaedt, J.A. Knapp, I.M. Robertson, Acta Mater. 56 (2008) 794–801. [5] M. Ames, J. Markmann, R. Karos, A. Michels, A. Tschope, R. Birringer, Acta Mater. 56 (2008) 4255–4266. [6] G.D. Hibbard, K.T. Aust, U. Erb, J. Mater. Sci. 43 (2008) 6441–6452. [7] G.D. Hibbard, V. Radmilovic, K.T. Aust, U. Erb, Mater. Sci. Eng. A 494 (2008) 232–238. [8] K. Bae-Kyun, J.A. Szpunar, R. Varano, Mater. Sci. Forum. 408–412 (2002) 937–942. [9] M.C. Iordache, S.H. Whang, Z. Jiao, Z.M. Wang, Nanostruct. Mater. 11 (1999) 1343–1349. [10] V.Y. Novikov, Interface Sci. 6 (1998) 77–84. [11] U. Klement, M. Da Silva, W. Skrotzki, J. Microsc. (2008) 455–463. [12] A.D. Rollett, W.W. Mullins, Scr. Mater. 36 (1997) 975–980. [13] A.D. Rollett, D.J. Srolovitz, M.P. Anderson, Acta Metall. 37 (1989) 1227–1240.
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