Processing Parameter Influence on Texture and Microstructural Evolution in Cu-Nb Multilayer Composites Fabricated via Accumulative Roll Bonding JOHN S. CARPENTER, RODNEY J. MCCABE, SHIJIAN J. ZHENG, THOMAS A. WYNN, NATHAN A. MARA, and IRENE J. BEYERLEIN A combination of accumulative roll bonding and rolling is used to fabricate bulk sheets of multilayer Cu-Nb bimetallic composites. Alterations in the processing sequence are made in comparison with prior studies in order to expand the processing window available for bimetallic multilayer composites. Cu-Nb composites with layer thicknesses ranging from 45 lm to 10 nm with accompanying total strains of 3.8 to 12.21 are characterized via neutron diffraction, electron back scatter diffraction, and transmission electron microscopy. These characterization methods provide microstructural information such as layer morphology and grain morphology as well as orientation information such as texture and interface plane normal distribution. The evolution of these microstructural characteristics is collected as a function of increasing strain. These results can provide guidance, inputs, and validation for multiscale predictive models that are being developed on materials with interfacially-driven properties. Finally, synthesis pathways are presented that allow the fabrication of nanoscale multilayer composites with predominant interfacial structures. These fabricated materials are ideal for exploring the relative importance between inter-phase interfacial density and atomic interfacial structure in determining material properties. DOI: 10.1007/s11661-013-2162-4 The Minerals, Metals & Materials Society and ASM International (outside the USA) 2014
I.
INTRODUCTION
PHYSICAL vapor deposited (PVD) copper (Cu)niobium (Nb) nanolamellar composites[1–11] have exhibited a variety of desirable properties such as high strength,[1,2,9,10] good ductility,[9,10] high thermal stability,[4,5] enhanced resistance to radiation damage,[6,7] and resistance to shock damage.[11] The desirable properties of PVD Cu-Nb are attributed to the high density of inter-phase boundaries coupled with the particular/ special arrangement of atoms at the interface providing the material with a unique ability to self-heal in the presence of large defect fluxes.[12–16] PVD Cu-Nb composites having equal volume fractions of Cu and Nb display orders-of-magnitude improvements in these properties compared with bulk Cu and bulk Nb.[1,2,4–7,9–11]
JOHN S. CARPENTER and RODNEY J. MCCABE, Staff Members, are with the MST-6: Metallurgy Group, Los Alamos National Laboratory, Los Alamos, NM. SHIJIAN J. ZHENG, Postdoctoral Researcher, is with the Center for Integrated Nanotechnologies, Los Alamos National Laboratory. THOMAS A. WYNN, formerly Research Technologist, with MST-6: Metallurgy Group, Los Alamos National Laboratory, is now a Graduate Student with the Materials Science and Engineering Department, University of California-Davis, Davis, CA. NATHAN A. MARA, Staff Member, is with the MST-6: Metallurgy Group, Los Alamos National Laboratory, also with the Center for Integrated Nanotechnologies, Los Alamos National Laboratory. IRENE J. BEYERLEIN, Staff Member, is with the Theoretical Division, Los Alamos National Laboratory. Contact e-mail:
[email protected] Manuscript submitted June 19, 2013. Article published online January 10, 2014 2192—VOLUME 45A, APRIL 2014
PVD Cu-Nb exhibits a Kurdjumov–Sachs (KS) orientation relationship with the close packed planes {111}Cu|| {011}Nb joined at the interface and h10-1iCu||h11-1iNb.[7,12–15] This particular arrangement has been found via modeling to provide excellent sites for both dislocation nucleation as well as the recombination and annihilation of defects instigated by radiation bombardment and high strain rates.[12–16] The displayed properties make the fabrication of Cu-Nb nanolaminar composites in a bulk form desirable for applications as varied as transportation, defense, and nuclear energy. Early research into bulk nanoscale Cu-Nb composites, predating the PVD studies, focused on creating conductor materials for high field magnets and superconductors.[17–19] The fabrication methods in this early research utilized casting or powder metallurgy with an end product of thin Nb ribbons (10 to 20 pct Nb by volume) within a majority Cu matrix.[17–25] Coldrolling of these Cu-Nb composites was performed and mechanical behavior, microstructure, texture evolution, and electrical properties were investigated.[17–25] Other early research into bulk fabrication made use of bundling and drawing to fabricate Cu-Nb nanocomposite wires suitable for magnet-based applications.[26–30] The severe plastic deformation method of accumulative roll bonding (ARB) has recently been explored as a fabrication method to create equal volume fraction Cu-Nb composites with layer geometry similar to the PVD layer morphology.[31–38] ARB is a forming technique where stacks of material in sheet form are bonded using significant rolling reductions. Unlike casting or powder metallurgy, this method can be used to create lamellar composites with controlled METALLURGICAL AND MATERIALS TRANSACTIONS A
average individual layer thicknesses (h). With regular, planar inter-phase interfaces, ARB materials are ideal model systems for comparison with PVD and the investigation of the role of interfaces in optimizing the desirable material properties described earlier. Texture evolution,[31,34] hardness,[37] thermal stability,[37,39] radiation resistance,[35] interfacial structure,[33–37] and strength[40] have been explored for these ARB Cu-Nb composites. Results indicated a similar hardness,[37] layered morphology,[31–38] strength,[40] and thermal stability[37,39] in comparison with PVD Cu-Nb. The texture evolution and interfacial structure, however, differed both from initial PVD Cu-Nb,[31,34] rolled PVD CuNb,[3,8] and the prior published results for cold-rolled Cu-Nb.[19,22] In addition, while twinning was not observed in PVD Cu-Nb,[31–34] rolled PVD Cu-Nb,[3,8] and cold-rolled Cu-Nb,[19,22] deformation twinning, associated with the development of the {552}h115i texture component in Cu, was observed when h < 100 nm in ARB Cu-Nb.[33,34,41] The effect of processing parameters during the ARB fabrication method, however, was not described qualitatively or quantitatively with regards to texture evolution and interfacial structure. In this study, ARB processing parameters such as annealing and changes in rolling direction are explored in terms of their effects on morphology and texture in the resulting Cu-Nb nanolamellar composites. The overarching goal is to explore whether processing parameters can be manipulated as a means for precisely varying grain morphology and texture in multilayer materials. Texture is studied through a combination of neutron diffraction and electron backscatter diffraction (EBSD) for layer thicknesses ranging from 45 lm to 30 nm. The use of EBSD at small length scales ( 200 nm were prepared for EBSD by cross-sectioning, mounting, grinding, and polishing with the final two polishing steps consisting of vibratory polishing using MasterPrep (Buehler) on a MasterTex cloth (Buehler) for 4 hours and Colloidal Silica (Allied) on a MasterTex cloth (Buehler) for 15 minutes. The mounting method chosen was a low temperature epoxy mount in order to prevent the temperature associated with hot mounting from affecting grain morphology. This point was especially critical when h ‡ 714 nm where the driving force for recrystallization due to retained plastic strain was high. The EBSD scans were used to obtain texture information, grain morphology, and layer morphology. Steps sizes of 0.2 lm were used for h ‡ 45 lm and steps sizes of 0.1 lm were used for 200 nm < h < 45 lm. Step sizes of 50 nm were utilized for scans on samples with h £ 200 nm. Areas scanned varied widely but two examples are 70 9 70 lm for h = 200 nm and 24 9 140 lm for h = 45 lm. Grain morphology was captured via EBSD through the use of standard functions within OIM (TSL, Draper, UT) software. Grains were defined through a minimum number of pixels (six) and through a minimum misorientation from neighboring grains (5 deg). Fitted ellipses to the remaining grains were generated utilizing OIM and major and minor axes for the fitted ellipses were calculated. Finally, it was assumed that grain shapes were elongated in the rolling direction. Grains with major axes >30 deg from parallel with the heterophase interfaces were excluded so as not to introduce confusion between the rubric of major axis = rolling direction and minor axis = normal direction. Less than 5 pct of grains for each length scale had fitted ellipse major axes that were further than 30 deg from the rolling direction. Normally, the practical spatial resolution of EBSD inhibits its use for studies of microstructures with feature sizes less than 100 nm.[36] Due to the geometry of our layered materials, we were able to utilize EBSD for the lower length scales regime of Cu-Nb composites by employing a new mounting technique. The new technique, called the WM technique, involves mounting our multilayer at an angle on a Cu wedge.[42] For a 10 deg wedge this has the effect of increasing the apparent layer thickness by a factor of 5.8 compared to the actual layer thickness. For example a 58 nm layer appears to be 334 nm. Using the wedge technique, we are able to examine texture information and layer morphology using EBSD at 30 nm < h < 100 nm. A full description of the method used can be found in Reference 42. High resolution and diffraction contrast transmission electron microscopy (HRTEM and TEM) were used to investigate the atomic structure of the interfaces, the presence of twinning, and grain orientation. Preparation of the samples was performed through grinding, polishing, dimpling, and ion milling and the investigation occurred in a Tecnai (FEI) TEM operating at 300 keV. The characterization techniques utilized in this study, to this point, provide details of the texture and microstructure evolution with increasing strain. In order to METALLURGICAL AND MATERIALS TRANSACTIONS A
understand the deformation mechanisms responsible for the observed texture evolution at very fine length scales and during the strain path change, viscoplastic self-consistent (VPSC) modeling is used. VPSC is a polycrystalline scheme with broad applicability for predicting grain morphology and texture evolution with increasing plastic strain for polycrystal materials.[48] The application of this methodology to a rolled bimetallic composite has been described in an earlier publication.[34]
IV.
RESULTS AND DISCUSSION
A. Texture Results for CC Cu-Nb ODFs collected from neutron diffraction data are shown in Figure 1 for both Cu and Nb phases for several length scales. Particular component intensities in the Cu phase are observed to strengthen with decreasing h for h ‡ 50 nm. A particular orientation, {112}h111i, with orthotropically symmetric orientations at (/1, F, /2)(39.2, 65.9, 26.6 deg) and (90, 35.3, 45 deg) is observed to strengthen. This component is a common component in rolled FCC material and is called the C component.[49,50] This orientation has been observed to become dominant in both rolled single-phase Cu (pct reduction = 97 to 99 pct) and Cu in cast and rolled Cu/ 20 pct volume fraction Nb composite (pct reduction = 88 to 99.5 pct).[22] The rotation of crystals towards the C orientation in our CCARB multilayers is also consistent with what is observed in the aARB material in Reference 34. This indicates that the plastic slip that governs the texture evolution during ARB at for h ‡ 50 nm is largely unaffected by annealing treatments or maintenance of RD during processing. As h is reduced from 50 to 20 nm in the Cu phase, the C orientation decreases in intensity while orientations near (27.0, 57.7, 18.4 deg) and (0, 67.5, 25.5 deg) increase in intensity. The orientation near (27.0, 57.7, 18.4 deg) is a common orientation in rolled FCC materials and is called the S or {123}h634i component. The other orientation near (0, 67.5, 25.5 deg) is not a common orientation in rolled FCC materials suggesting that it is unstable. With further rolling to h = 10 nm, this non-common orientation weakens while the S orientation maintains its intensity and a new orientation strengthens near (90, 90, 45 deg) and (0, 45, 0 deg) corresponding to the Goss {110}h100i orientation. The Goss component[50] is a known recrystallization orientation associated with nucleation near shear bands.[51,52] However, during recrystallization, the Goss component is usually paired with a Q orientation, {013}h231i at (45, 15, 10 deg), since both share a 40 degh111i orientation relationship to the initial rolling texture.[52] The lack of Q orientation indicates that the rearrangement towards the Goss component likely does not arise from recrystallization. Finally, for the aARB material,[34] the dominant orientation never departed significantly from the {112}h111i orientation at h = 10 or 18 nm, although an increase in the Goss component was noted at h = 18 nm as seen in the fiber plot in Figure 2. The VOLUME 45A, APRIL 2014—2195
Fig. 1—Neutron-based ODFs for CC Cu/Nb at length scales h = 45 lm, 1.3 lm, 50 nm, 20 nm, and 10 nm for both Cu and Nb. A continuous evolution of increased intensity near the {225} component in Nb is visible with decreasing h. The Cu phase shows an increase in {112} intensity down to h = 50 nm before a separate evolution towards a strong brass component begins. (Color figure online).
Fig. 2—Fiber plots presented for Cu and Nb that compare evolution of specific components for both CC and aARB Cu/Nb.
behavior in aARB material deviates significantly from CCARB material as seen in Figure 2. It is noted that neither the Goss nor the S component were observed as significant populations in the investigations on rolled cast Cu/20 pct volume fraction Nb composite.[17–25] The evolution below h = 50 nm in the CC material is, therefore, unique when compared to aARB, rolled single-phase Cu, and rolled cast Cu/20 pct volume fraction Nb composite. This result indicates that, at small length scales in Cu, the texture evolution is processing dependent. In the Nb phase, a continuous evolution in texture is observed throughout the length scales. A gradual strengthening with decreased h of (0, 30, 45 deg) or the {225}h110i component is seen in Figure 1. Evolution is also noted towards a single component system, i.e., all grains within the Nb phase evolve towards two orthotropically symmetric orientations. This is different from 2196—VOLUME 45A, APRIL 2014
Cu where various orientations were present at all length scales. In rolled single-phase Nb, rolling past 99 pct leads towards a predominant {112}h110i component.[22] In rolled cast Cu/20 pct volume fraction Nb composite, further rolling past 99 pct leads towards a relatively equal distribution of intensity between {001}h110i and {112}h110i.[22] For aARB material[34] a strong Nb texture did not develop until h = 48 nm, while strong texture began to develop in Nb between h = 714 nm and h = 200 nm in CC material. Strong texture is defined as having a single dominant texture component with an ODF intensity surpassing 10. Prior to h = 48 nm, the texture evolution of Nb in the aARB material resembled that of the rolled cast Cu/20 pct volume fraction Nb composite. In addition, after developing the strong texture at h = 48 nm, the aARB material shows continued strengthening of the {112}h110i component until h = 18 nm before a reversal in evolution takes place at METALLURGICAL AND MATERIALS TRANSACTIONS A
Fig. 3—EBSD and WMEBSD-based unique grain maps for h = 8 lm and 58 nm for CCARB Cu/Nb. (Color figure online).
h = 10 nm towards a weaker texture as seen in Figure 2. The CC material shows continued strengthening with a continuous and steady evolution towards a {225}h110i as h is reduced from 200 to 10 nm. The 38 pct decrease in intensity of the {112}h110i component in Nb observed between h = 18 and 10 nm in aARB material[34] is inconsistent with the notion that the annealing treatments delayed the evolution of Nb as this reversal in texture evolution is not observed in CCARB. The texture results for Cu and Nb indicate the need in predictive processing models to include processing parameter variables and their accompanying potential effects on texture.[45,53,54]
B. Microstructural Observations for CCARB Cu-Nb Example microstructures of CCARB Cu-Nb composites for h = 8 lm and h = 58 nm are presented in Figure 3 via unique grain maps that represent a view down the TD. It should be noted that h = 58 nm data were collected via the WMEBSD technique and the image presented in Figure 3 has had the vertical dimension decreased in order to present a view as would be seen down the TD. Layer morphology is observed with both Cu and Nb showing relatively equiaxed grains at h = 8 lm and more elongated grains at h = 58 nm. Our EBSD-based microstructure data set is unique in that it captures statistically significant amounts of grain information on a system rolled from h = 45 lm to h = 30 nm. It should be noted that little difference was noted between the prior neutron diffraction-based and EBSD-based textures. Shear banding was observed at all length scales for h £ 45 lm. Shear banding becomes more profuse with diminishing h, with shear band spacing decreasing from around 200 lm (h = 8 lm) to 10 lm (h = 200 nm, not shown) to 6 lm (h = 58 nm). Despite the presence of significant shear banding, layers were observed to maintain continuity through these features.
METALLURGICAL AND MATERIALS TRANSACTIONS A
Figures 4 through 6 provide area fraction distributions of RD, ND, and TD grain lengths measured using EBSD. Data for h = 135 and 58 nm was obtained through use of the WMEBSD technique. Cu exhibits similar RD grain size distributions for length scales h = 20 lm to h = 58 nm. The Nb phase, on the other hand, shows an evolution in RD grain length. At h = 20 lm several long grains are present which stretch across the entire scan area in EBSD. This small number of grains (~2 pct) takes up a disproportionate area (~27 pct) for the Nb phase. As h is further reduced, however, Nb’s RD grain length approaches a distribution similar to that for Cu. At h = 58 nm, however, it can be seen in Figure 4 that Nb still contains a greater fraction of longer grains as shown by the 15 pct area associated with grains with RD between 10 and 32 lm. The ND, as seen in Figure 5, shows a continuous evolution of grain size for both the Cu and Nb phases, largely correlating with the reduction in average layer thickness. For both h = 135 nm and 58 nm, the ND distribution suggests that each layer is one to two grains thick. At these length scales, all grains are associated with a bimetal interface on either one or both sides. Figure 5 also suggests that at h = 20 and 8 lm, layers have many grains through the thickness. For example, utilizing the peak values for Cu at h = 20 lm, there are 20 to 40 grains through the thickness. This indicates that a majority of the grains in the Cu and Nb phases at high length scales are not associated with an interface. Finally, Figure 6 shows the TD grain lengths at low length scales in both Cu and Nb. TD grain lengths are of interest in this study in order to understand the unique texture evolution that takes place at low length scales. Having TD grain lengths allows a 3D grain size to serve as an input for plasticity simulations. Figure 6, unlike Figures 4 and 5, does not show data with regards to h = 20 or 8 lm. Either an RD cross-section or RD wedge (wedge section at 10 deg to ND and 80 deg to RD) were used to measure TD grain dimensions. The h = 714 nm data is for a non-wedged specimen and
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Fig. 4—Plot showing relative RD grain lengths for CC Cu/Nb with h = 20 lm, 8 lm, 135 nm, and 58 nm. The 135 and 58 nm data was collected using the wedge technique. All measurements shown are EBSD based.
Fig. 5—Plot showing relative ND grain lengths for CC Cu/Nb with h = 20 lm, 8 lm, 135 nm, and 58 nm. The 135 and 58 nm data was collected using the wedge technique. All measurements shown are EBSD based.
shows peak values in Cu and Nb of between 1 and 1.8 lm for TD grain length. As length scale diminishes, the average TD grain length decreases indicating that grain refinement occurs in the TD direction. This is an unusual result as rolling provides a plane-strain compression loading state with no constraints placed on the TD during rolling. Small increases in the total TD dimension were measured over the course of this study after rolling reductions. The driving force for this grain refinement, therefore, is not related to the loading state during rolling which would increase the TD dimension of grain sizes. If grain volume is conserved during rolling in plane-strain compression, it is expected that RD grain dimensions would increase with an offsetting decrease in ND. Given the fact that ND diminishes by a large degree, TD grain length by a smaller degree, and RD by almost no degree, the grain evolution during rolling must consist of grain refinement in the RD and TD. The shear stress placed on the sample by the rolling process creates instabilities within grains larger than a critical 2198—VOLUME 45A, APRIL 2014
Fig. 6—Plot showing relative TD grain lengths for CC Cu/Nb with h = 714, 135, and 58 nm collected using the wedge technique. All measurements shown are EBSD based.
size in the RD and TD.[55] This causes the grains to break apart in the RD and TD. Given the static value of grain size in the RD, the critical value of grain size for this process in the RD has been found to be between 1 and 3.2 lm. In the TD, the slow continuous refinement indicates that the minimum grain size has not been found and that it is £560 nm. The difference between the critical grain size in the TD and RD is indicative of the difference in shear state and response mechanism of the grains in these two dimensions. With a 3D grain size measurement now possible on the same population of grains through the wedge-mount technique, the peak data suggests that at h = 58 nm, RD:ND:TD for Cu and Nb are 33:1:7 and 83:1:6, respectively. Finally, it is noted that the removal of annealing for the entire process neither did not lead to an increase in edge cracking, discontinuous layers, nor other common rolling defects such as center-splitting. This result serves to expand the potential processing window of ARB for developing nanolamellar composites. A specific example of the effect of this expanded processing window is duplex steels. Due to unwanted precipitates, duplex steels cannot utilize heat treatments to relieve differences in strength arising from variations in work hardening rates. C. Layer Thickness Study Backscattered electron imaging of TD cross-section samples at large length scales (h ‡ 714 nm) and WM-section samples at low length scales (h £ 714 nm) were utilized to measure average layer thicknesses. Specific methodology for extracting layer thicknesses can be found in Reference 42. This study was conducted in order to understand the variability in layer thickness induced by the ARB process. Table II provides average measured layer thicknesses for each sample along with the level of experimental uncertainty (indicative of pixel size) and a standard deviation as a measure of the layer thickness variability. In a prior work,[34] post-rolling annealing was utilized to equilibrate the strength METALLURGICAL AND MATERIALS TRANSACTIONS A
Table II. h (lm) 45 20.6 7.8 1.3 0.200 0.135
Measured Values of Cu and Nb Layers Thickness Utilizing BSE Results
N
hNb (lm)
sNb (lm)
sNb/hNb
hCu (lm)
sCu (lm)
sCu/hCu
u (lm)
50 306 724 838 304 658
44.567 19.829 7.625 1.239 0.217 0.132
5.918 4.217 3.557 1.071 0.270 0.179
0.13 0.21 0.47 0.86 1.24 1.36
48.186 21.304 7.918 1.333 0.247 0.121
6.819 3.184 2.956 0.638 0.157 0.089
0.14 0.15 0.37 0.48 0.64 0.74
±0.486 ±0.423 ±0.182 ±0.029 ±0.009 ±0.013
Variables used in table represent nominal layer thickness (h), numbers of layers measured (N), measured average layer thickness for layer x (hx), standard deviation for layer x (sx), and experimental uncertainty (u).
between Cu and Nb due to their difference in work hardening rate. In this study where annealing is not used, a difference in strength between Cu and Nb is believed to have developed early in the rolling process. However, the measured average thicknesses for Cu and Nb are more or less equal, and the dichotomy in strength that develops does not induce differences in amounts of co-deformation to any noticeable degree. This conclusion is supported by the equal area fraction under the layer thickness distribution which indicates the material is still 50/50 Cu/Nb. This conclusion is also consistent with prior literature.[45] This result would be of critical consequence in the design of multiscale, predictive processing models. Table II shows that as nominal layer thickness decreases, the standard deviation, normalized by the measured layer thickness, increases. This observation is likely indicative of the effect of the constraints placed on plastic deformation within the layers at low length scales. When layers are thicker, tens to hundreds of grains are observed through a single layer thickness. This provides a multitude of different options for local grain reorientation and slip to accommodate the strain associated with roll bonding. When layers are progressively thinned only one or two grains are present across each layer’s thickness. This constrains each layer’s ability to accommodate the existing loading conditions. This causes a larger variation in Cu and Nb layer thicknesses at low length scales as layers with grains with favorable orientations for deformation decrease in size more than those with less favorable orientations.
D. Interfacial Character Interfacial character is of high importance in this study due to the observed interface-specific defect mitigation behavior observed in PVD Cu-Nb.[12–15] Regardless of length scale, spatially correlated orientation data is necessary to separate the character of the grains associated with the interface and those that are not. EBSD presents an ideal solution to this problem by providing large data sets that can enable statistically significant measures of local texture. WMEBSD allows the use of EBSD across lower length scales to provide statistically significant measures of texture in nanograined and ultrafine grained material.[42] In this study, successful WMEBSD scans were performed for h ‡ 30 nm. METALLURGICAL AND MATERIALS TRANSACTIONS A
Figure 7(a) shows a full phase map and 7(b) shows a partitioned phase map containing only grains associated with the interphase interfaces for a h = 20 lm sample. Eight interfaces are visible and, consistent with the grain morphology results discussed earlier, the Nb grains appear to be larger and longer than their Cu counterparts. A majority of the Cu grains are small and equiaxed consistent with grain refinement expected from severe plastic deformation.[45] The Nb phase displays a more varied distribution of grain sizes with refinement having not fully broken up the larger, elongated grains. In addition, it was found that grain shape and size distributions in both phases were insensitive when OIM grain definition misorientation was varied between 2 and 5 deg. Figure 8 presents EBSD-based ODFs for Cu and Nb at h = 45 lm and 30 nm. Figures 8(a), (c), (e), and (f) were confirmed to be representative of the bulk through a comparison with neutron diffraction-based ODFs. Intensities observed via EBSD are roughly double those observed in neutrons as OIM software normalizes the intensity for each individual phase. For the Nb phase, Figure 8(c) shows that the Nb texture is strongly reminiscent of rolled BCC metal behavior with a and c fibers both clearly present.[49,50] Figure 8(d), however, shows that the interfacial grains are considerably more textured with a near-single component texture observed with an orientation near {225}h110i. This is consistent with the observed texture at very fine length scales, h = 30 nm, seen in Figure 8(f), and indicates that the grains constrained by the heterophase interface on one side at high length scales respond to the applied rolling stress in a similar way as the severely constrained grains at low length scales. The full population of Cu grains seen in Figure 8(a) shows distinct differences when compared with the interfacial grains displayed in Figure 8(b). Instead of showing a consistent b fiber intensity consistent with rolling of FCC metal as in Figure 8(a),[49,50] a particular component, the copper component or {112}h111i orientation, is slightly stronger in grains associated with the interface. This is consistent with the texture result when length scale is reduced to h = 30 nm. Figure 8(e) shows that at h = 30 nm, the general form of the b fiber remains but with a slightly stronger copper component. The results imply that the texture at the interface evolves more rapidly than does in the bulk. The bulk grains are gradually winnowed down leaving a texture at small length scales that is equivalent to interfacial grains at larger length scales. The result at high length scales VOLUME 45A, APRIL 2014—2199
Fig. 7—(a) Phase map obtained via EBSD for a h = 20 lm sample. Red corresponds to Cu and green to Nb. (b) Partitioned phase map showing only Cu and Nb grains that are associated with interphase interfaces. (Color figure online).
could also hint at what the stable interface character will look like at low length scales (i.e., a near-single component texture in Nb paired with a predominant {112}h111i population in Cu). It should be noted that the results observed at h = 45 lm were consistent with those observed at h = 20 and 8 lm, as well. Measured textures indicate that interface-associated grains in both Cu and Nb exhibit different textures than those displayed by grains within the ‘bulk’ of the layers. Modeling studies[45,53] undertaken at high length scales likely will need modifications such that each modeled layer has an interface-affected region. Further characterization that looks at the texture of grains one, two, three, etc. grains away from the interface would be useful in determining whether a graded function approach is needed or if a binary model is more appropriate. E. Effect of Strain Path Change at Low Length Scales In order to develop new interfaces and to study the effect of a strain path change on microstructure and texture, a section of CC plate with h = 58 nm was rotated 90 deg and then rolled to various reductions hT = 50, 30, and 20 nm where hT refers to the layer thickness for the strain path change samples. In the context of these strain path change samples, RD and TD refer to the original RD and TD directions prior to the 90 deg about the ND sample rotation. Figure 9 shows representative RD section TEM micrographs and SAD patterns for the three samples that underwent the strain path change. Little change is noted in the layered morphology at hT = 50 nm as the layers still appear to be largely flat and planar. The SAD pattern for hT = 50 nm indicates that the texture in the TD is spread around the {110}Nb and the {111}Cu. With continued rolling to hT = 30 nm, the SAD pattern indicates that a larger spread in TD orientations develop. In addition, the representative TEM micrograph 2200—VOLUME 45A, APRIL 2014
in Figure 9 shows that large amounts of shear banding, but no layer pinch-off, have developed in the microstructure. A TEM investigation of the entire cross-section (total sample thickness = 155 lm) of the sample viewed along the RD indicated that the image in Figure 9 was representative. The representative micrograph for hT = 20 nm shows that the large amounts of shear banding present in the hT = 30 nm have largely disappeared leaving a more planar interface structure. The SAD pattern exhibited at hT = 20 nm is consistent with a highly textured material with Nb exhibiting a single component texture with {110}RD and {200}ND while Cu showing strong intensity associated with {111}RD but less concentrated intensity along the ND. The macrotexture measured for h = 58 nm (shown in Figure 10) indicates that the predominant interfacial planes were {113}Nb and {112}Cu/{110}Cu. Although interfacial sliding would not occur as easily as interfaces with {111}Cu paired with {110}Nb, the orientations in the TD at h = 58 nm could exhibit interfacial sliding.[56,57] During rolling, the amount of tension along the RD is larger than the amount of lateral tension along the TD. This difference in stress state leads to a strong intensity in the RD associated with {111} in Cu and {110} in Nb. It is proposed that the 90 deg strain path change leads to the shear instabilities along the TD and that these result in shear banding. The initial texture present in the material at h = 58 nm was not well-aligned to accommodate the increase in TD tensile stresses after the strain path change. A large reorientation was necessary in order to achieve the orientations preferred for the TDs in both phases after the strain path change. This large reorientation was accomplished through shear banding. Once the grains are reoriented with respect to the TD, shear banding is no longer required in order to perform the reorientations needed to create the stable texture observed at h = 58 nm. The material then exhibits flatter, more planar interfaces as observed in Figure 9 METALLURGICAL AND MATERIALS TRANSACTIONS A
Fig. 8—EBSD-based ODFs representing texture for (a) all Cu grains seen for h = 45 lm, (b) only those Cu grains associated with the interphase interface in (a), (c) all Nb grains for h = 45 lm, (d) only those Nb grains associated with the interphase interface in (a), (e) all Cu grains for h = 30 nm sample, and (f) all Nb grains for a h = 30 nm sample. (Color figure online).
for hT = 20 nm. It should be noted that the lamellar structure shows remarkable continuity in layer structure despite the significant amount of grain reorientation. F. Twinning The twinning deformation mechanism was observed in the Cu layers of individual layer thickness (h) 50 nm and lower in aARB Cu-Nb material.[33,34,36] These twins were observed to originate at near-{112}Cu||near-{112}Nb interfaces that displayed a KS orientation relationship[33] and shifted the interfacial plane from a {112} in Cu to a {552}. The accompanying alteration in texture was observed via neutron diffraction at h = 18 nm.[34] Mechanisms for the twinning behavior in ARB Cu-Nb that have been discussed in literature[32,33,38,41] suggest layers exhibiting a single grain through the thickness with the special interfacial arrangement is necessary for METALLURGICAL AND MATERIALS TRANSACTIONS A
deformation twinning in the Cu layer. Although a maximal value of h to observe twinning for this ARB Cu-Nb was discussed,[34] the effect of altering processing conditions was not explored. Twinning was observed at low length scales in CC material with significant populations when h = 20 nm. Figure 11 presents a representative HRTEM view of the types of twins observed in CCARB material at h = 20 nm. Twins were only observed in Cu and presented angles with the interface (near 19 deg) consistent with studies of aARB material[33,36,38] and CC material.[39] In addition, Reference 33 indicated that twins arose from {112}Cu||{112}Nb interfacial structures in aARB material and this was also observed in CC material. Given that CC Cu-Nb at h = 20 nm presents significantly more {112}Nb interfacial planes than does aARB Cu-Nb at h = 18 nm, it would be reasonable to expect a larger population of twins. TEM does not lend VOLUME 45A, APRIL 2014—2201
Fig. 9—Diffraction contrast micrographs for hT = 50, 30, and 20 nm showing maintenance of lamellar structure despite cross-rolling history. Accompanying SAD patterns show that at hT = 20 nm strong texture develops unlike at 30 and 50 nm.
Fig. 10—Neutron-based IPFs for h = 58 and 10 nm as well as hT = 20 nm. The IPFs indicate that near-single component textures are obtained in Cu and Nb after cross-rolling and that a characteristic interface plane is obtained with longitudinal rolling. At low length scales, EBSD indicated that layers are one or two grains through the thickness allowing the ND IPFs to be utilized as interface plane normal distributions. (Color figure online).
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METALLURGICAL AND MATERIALS TRANSACTIONS A
G. Achieving Two Unique Interfacial Structures Through Processing Parameter Control
Fig. 11—HRTEM image of h = 20 nm.
twins present in Cu phase at
Fig. 12—Twin observed in hT = 30 nm sample. The twin forms a near 20 deg angle with the interface and the Cu phase has a near{112} character consistent with Ref. [33]. Nb, however, does not have a near-{112} character inconsistent with Ref. [33].
itself to a quantitative measurement of the twinning phenomenon and future precession electron diffraction or WMEBSD studies will address this.[38] The length scale at the onset of twinning, however, is consistent with that reported for the aARB material[34] indicating that twinning arises independently of processing conditions. Twinning is sparsely observed via TEM for hT = 50 nm with an increasing number visible at hT = 30 and 20 nm. A representative twin is shown from hT =30 nm in Figure 12. Consistent with Han et al.,[33] it is noted that the twin makes a near 20 deg angle and is associated with a {338} or a near-{112} orientation in the Cu phase. The Nb orientation present, however, is near (001)[110] which differs markedly from the orientation observed by Han et al.[33] The twin observed might have been inherited from the prior longitudinal rolling or, alternatively, it might have nucleated during cross-rolling from the {338}Cu||{001}Nb interface. Given the stochastic nature of twinning coupled with the uncertain history of this twin’s nucleation, this result is reported in the hope of instigating a twinning study of the {338}Cu||{001}Nb interface utilizing recent twinning models.[33,38,39]
METALLURGICAL AND MATERIALS TRANSACTIONS A
Through manipulation of processing parameters, two unique interfacial structures were fabricated for Cu/Nb multilayers that were geometrically and chemically similar. Figure 10 presents the IPFs for h = 10 nm and hT = 20 nm, where near-single component textures in both Cu and Nb are present. HRTEM confirmed the results of the neutron-based IPFs as seen in Figure 13. In the case of h = 10 nm, the texture present in the Cu is believed to arise largely as a result of twin growth fully removing the stable {112}ND orientation present at h = 58 nm and replacing it with a {552}ND. Little change in texture is noted in the Nb phase as the near {112}ND remains stable and the texture, as a whole, maintains a single component nature. Maintaining a consistent RD in the absence of any annealing treatments coupled with the advent of deformation twinning in Cu at low length scales leads to a consistent interfacial structure of {551}h1 1 10iCu||{112}h110iNb.[39] For hT = 20 nm, it is noted that the Cu phase, through massive reorientation, has moved from a combination of {112} and {110} in the ND to a {110}ND. The Nb phase has also exhibited a massive reorientation such that the {001} is now aligned with the ND. This new interfacial structure, {110}h111iCu|| {001}h110iNb is unique with regards to the interfacial structure obtained at h = 10 nm and is unique from that obtained in PVD Cu/Nb of {111}h110iCu|| {110}h111iNb. A recent analytical model predicted the defect structure of this interface does not support twinning.[41] The fabrication of Cu/Nb multilayers with similar chemical and geometrical natures with regards to PVD allows a study to take place that delineates the effect of interfacial structure vs interfacial density in determining mechanical behavior and resistance to damage (shock, irradiation, etc.). When length scales in Cu/Nb metallic multilayers are reduced to the nanoscale, orders-ofmagnitude increases in strength, thermal stability, resistance to radiation damage, and resistance to shock damage are observed.[1–11] Future studies will provide direct one-to-one comparisons between interfacial structures such that the role of atomistic structure at the interface in determining material behavior can be determined. H. Manipulating Grain Morphology Through Processing Parameter Control Figure 14 presents RD and TD grain lengths measured via WMEBSD for sub 100 nm layer thickness multilayers fabricated via three different processing conditions. ND grain diameters are not shown as the layers are one grain thick at these length scales, and so equal the layer thickness. With regards to the RD grain lengths it is observed that the h = 86 nm sample fabricated with annealing achieves the longest grains, on average. Approximately 85 pct of the grains for
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Fig. 13—Diffraction contrast and high resolution TEM micrographs supporting texture based arguments for unique and characteristic interfacial arrangements for h = 10 nm and hT = 20 nm. The {551}Cu||{112}Nb for h = 10 nm and {110}Cu||{001}Nb for hT = 20 nm. High resolution images also provide evidence of chemically distinct interfaces and regular faceting.
h = 86 nm aARB are between 10 and 30 lm in length. When the annealing step is removed, this percentage sharply declines with the Nb phase at h = 58 nm having 55 pct of its grains with diameters between 10 and 30 lm and Cu with ~30 pct. With the addition of crossrolling, very little difference in RD average grain lengths is noted for either Cu or Nb. This indicates that RD grain lengths at the nanoscale have been greatly affected by the annealing treatments and, to a much lesser extent, by the maintenance of RD during fabrication. Heat treatments likely lead to short-range diffusion and the consolidation of more volume with grains that are of a stable orientation rather than those with an unstable orientation. Figure 14 also shows results related to the TD grain diameters for the same three samples. A comparison between the aARB material (h = 86 nm) and the CC material (h = 58 nm) with no cross-rolling shows a similar relationship to that seen in the RD. CC material fabricated without annealing tends to have smaller diameters in the TD than does aARB material. This is evidence that heat treatments encourage grain growth in both the TD and RD dimension. While cross-rolling did little to affect the RD direction grain diameter, the effect on TD grain diameter was marked. TD grain diameters
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in the CC material without cross-rolling lead to ~90 pct of grains with diameters between 0.1 and 1 lm. Crossrolling the CC material to a reduction of less than 50 pct leads to ~90 pct of the Cu grains between 1 and 3 lm and ~80 pct of Nb grains over 3 lm in diameter. These results indicate that through a combination of heating and cross-rolling, grain morphology can be manipulated to increase grain size in the RD and TD dimensions.
V.
EXPLANATION OF TEXTURE RESULTS THROUGH VPSC
As part of the analysis, polycrystal calculations were carried out using the visco-plastic self-consistent (VPSC) scheme[48] to determine the reorientation of the individual phases under rolling. Over the past decade, the VPSC polycrystal model has been advanced for predicting the texture evolution of various single-phase and multi-phase polycrystalline materials under high levels of strain and strain path changes.[43,58–60] VPSC is a mean-field approach and, therefore, does not directly account for the influence of nearest neighbor grains or interfaces. In prior work, such effects have been added by enforcing co-rotation, special hardening laws, and
METALLURGICAL AND MATERIALS TRANSACTIONS A
Fig. 15—ND IPF schematic showing dual pathway for copper component to migrate towards a stable orientation. Red arrows signify the pathway for non-twinned material while black indicates pathway for twinned material. These pathways are consistent with the experimentally measured orientations at h = 10 nm. (1) Copper, (2) twinned copper, (3) twinned and rotated copper, (4) S, and (5) Goss, brass. (Color figure online).
Fig. 14—Plots showing grain lengths in both the original RD and the original TD for three sets of processing conditions. It is observed that through processing parameters, grain size can be controlled. (Color figure online).
stress or strain fluctuations.[43,61–63] In this work, VPSC calculations are used to understand texture transitions that takes place (1) in Cu and Nb, when the layers refine below 50 nm down to the finest scales (at 10 nm) and (2) in Cu and Nb during a strain path change. An obvious disadvantage to the model is that interfaces are not directly modeled. This being the case, these calculations are intended to give first insight on how the constituent phases would deform alone without the influence of the interface. That is, comparing the predictions with measurement can indicate whether or not the interface had an influence on the texture. Recently, models that enforce traction and displacement continuity at the interface have been developed and used to study this kinematic constraint effect on texture evolution.[53,54] Extensions to account for more interface effects are in development. The initial textures used in the present VPSC calculations were representations of the measured texture discretized into several thousand discrete orientations. We begin with initial textures corresponding to material that has undergone several thousands of percent strain and initial grain shapes with a very high aspect ratio of 20:2:1 corresponding to RD:TD:ND. This value adequately captures the effect of elongated grains since the model does not produce appreciable differences for more extreme initial grain shape aspect ratios such as those observed experimentally in this study. The grain shapes and orientations are allowed to change individually with further deformation. In what follows, VPSC is used to calculate the texture evolution of the polycrystal accounting for the hardening, slip activity, reorientations, and shape changes of the individual crystals. In all cases, plane-strain compression is applied to the polycrystal. The constitutive model for Cu and Nb was described previously in METALLURGICAL AND MATERIALS TRANSACTIONS A
Reference 34, and includes a rate dependent flow rule in the form of a power law and Voce law hardening. Cu deforms by {111}h110i slip and {111}h112i twinning. For Nb, both {112}h111i and {110}h111i slip are made available, as the combination has been shown necessary for modeling single crystal bcc plastic anisotropy.[64] The predominant reorientation scheme[43,58,65] is used to capture the reorientation of crystals that have twinned. A. Texture Transition from 58 to 10 nm Under Monotonic Rolling We first consider the texture transition that takes place in the Cu phase due to the onset of deformation twinning. When twinning in the Cu phase of the ARB material begins as h reduces below 100 nm, the deformation in the Cu phase changes from homogeneous to heterogeneous. The Cu layers develop regions of twinned and untwinned crystal (Figure 11). Under further deformation, two events happen: the twinned and matrix regions will take reorientation pathways different than when deformation was mediated by slip only and the twin volume fraction increases while the matrix volume fraction decreases. In what follows, we focus on the Cu phase and not Nb, since it did not undergo twinning or an abrupt texture transition. It is observed that the twins form consistently along a {111} plane that lies 19.47 deg from the {112} interface plane as seen in Figure 11 and in published literature.[33,38] When twins develop along this special twin plane (so called 20 deg twin plane), they will reorient the parent at the C orientation to the {552}h115i orientation. The twinned orientations are unstable in rolling and thus with further rolling, the twinned regions reorient to {551}h1 1 10i and Goss {110}h001i. A schematic showing the reorientation pathway for a C component that has twinned on this particular plane is shown in Figure 15. This demonstrates that twinning provides one possible pathway by which the C (and Taylor {4 4 11}h8 8 11i component) have disappeared and the Goss component has strengthened. This pathway incidentally as been proposed by Hirsch et al.[65] to explain how deformation twinning leads to a brass-type texture in low stacking fault energy fcc alloys and has been discussed by Zheng et al.[39] VOLUME 45A, APRIL 2014—2205
The matrix regions in between adjacent twins can undergo coplanar slip along the {111} plane that lies parallel to the twin plane (Figure 11), since activating slip systems non-planar to the twin boundaries will become difficult. VPSC is used to calculate the reorientation pathway of representative matrix orientations that undergoes coplanar slip solely on this particular slip plane. Coplanar slip is modeled by increasing the CRSS fivefold on the planes that do not parallel the favored 20 deg twin plane. Figure 15 shows the resulting reorientation pathway taken by an untwinned matrix crystal with a starting orientation near the C component which is being deformed by several hundreds of percent strain. This reorientation path agrees well with the texture evolution of the Cu phase from h = 200 nm to h = 20 nm and explains the development of the S and Brass components. This suggests that the observed texture transition results primarily from the large fraction of matrix material (70 to 95 pct) being forced to deform by coplanar slip on planes that parallel the twin planes. The model is likely successful because it captures the texture evolution that results from most of the grains, but there are likely to be some processes the model overlooks. For instance, the model enforces coplanar slip over the entire deformation path. It is possible that at any point along the coplanar pathway, the crystal may attain an orientation in which coplanar slip no longer operates and the crystal will no longer take the pathway shown. Also, as deformation proceeds, the twins are increasing while the matrix is shrinking in volume fraction. Consequently, the matrix crystalline regions at any point along the ‘coplanar’ reorientation path shown in Figure 15 could stop reorienting by coplanar slip and reorient by twinning. Crystals in the vicinity of the C or Taylor orientation twin to near {552}h115i or {551}h1 1 10i, respectively, while those in the neighborhood of S twin to orientations that are near S in orientation space.[65] Unlike the Cu phase, Nb’s texture was not captured via VPSC modeling. Various adjustments to the relative slip activity of {112}h111i and {110}h111i slip were made, consistent with the attempt in Reference 34. In all cases, VPSC predicts that the interface character reorients towards a {001}h110i component, which is in fact consistent with published experimental results in systems of pure Nb.[22] VPSC, however, is unable to capture the stability of the {112}h110i Nb component observed in this study. More sophisticated crystal plasticity finite element modeling (CPEFM) on Cu-Nb bicrystals has recently been undertaken.[54] The plastic stability of many interfaces was explored and the near{112}h111iCu||near-{112}h110iNb interface structure was found to be stable. As in VPSC modeling, a combination of activated {112}h111i and {110}h111i slip systems was needed in the CPFEM model for the Nb phase. CPFEM added the additional constraint, however, of forcing the two crystals to co-deform. The difference in predictions between single-phase mean-field method (VPSC) and the spatially resolved bicrystal calculations (CPFEM) suggests that the kinematic constraint imposed by the interface plays a role in stabilizing the {112}h110i Nb component. 2206—VOLUME 45A, APRIL 2014
B. Texture Transition During Cross-Rolling from 58 to 20 nm Next we consider texture evolution during the strain path change. The Cu phase, at hT = 20 nm, presents a near-single component texture for the copper phase with a majority of the grains oriented such that {110}ND, {112}New-TD, and near-{112}New-RD. The VPSC model and constitutive laws utilized in Reference 34 were used to simulate this strain path change test. The starting texture was the measured deformation texture at h = 58 nm. The model grain morphology was 20:2:1 in the RD:TD:ND directions. To simulate cross-rolling, the polycrystal Cu or Nb was deformed in plane-strain compression in the ND and along the TD. The resulting deformation textures at a cross-rolled strain of (b) 0.4 and (c) 0.86 are shown in Figure 16 for Cu and Figure 17 for Nb. It is observed in Figure 16(a) that VPSC predictions fit the experimental data in Figure 10 extremely well with a largely single component texture developing of {110}ND, {112}RD, and near-{112}TD at a strain of 0.4. With continued cross-rolling to a strain of 0.86, however, the calculated deformation texture changes further, exhibiting a band of intensity between the {110} and {112} in the ND and between {110} and {447} in the TD. In fact, it appears to be evolving towards a texture similar to that observed in Figure 10 for h = 58 nm. This result, indicates that a strain path change leads to a unique near-single component texture in Cu under small strain only (
