The evolution of the deformed microstructure as a function of imposed plastic strain is of interest as it provides information on the material hardening characteristics and mechanism(s) by which cold work energy is stored. This has been extensively studied using transmission electron microscopy (TEM), where the high spatial and orientational resolution of the technique is used to advantage to study local phenomenon such as dislocation core structures and interactions of dislocations. With the recent emergence of scanning electron microscope (SEM) based automated electron backscatter diffraction (EBSD) techniques, it has now become possible to make mesoscale observations that are statistical in nature and complement the detailed TEM observations. Correlations of such observations will be demonstrated for the case of Ni-base alloys, which are typically non-cell forming solid solution alloys when deformed at ambient temperatures. For instance, planar slip is dominant at low strain levels but evolves into a microstructure where distinct crystallographic dislocation-rich walls form as a function of strain and grain orientation. Observations recorded using both TEM and EBSD techniques are presented and analyzed for their implication on subsequent annealing characteristics

King, W.E.

Schwartz, A.J.

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Approved for public release; further dissemination unlimitedPreprintUCRL-JC-137073Correlating Observations of DeformationMicrostructures by TEM and AutomatedEBSD TechniquesM. Kumar, A.J. Schwartz, W.E. KingThis article was submitted toMaterials Science and Engineering A Dislocations 2000Gaithersburg, MD, June 19-22, 2000June 5, 2000LawrenceLivermoreNationalLaboratoryU.S. Department of Energy DISCLAIMER  This document was prepared as an account of work sponsored by an agency of the United StatesGovernment.  Neither the United States Government nor the University of California nor any of theiremployees, makes any warranty, express or implied, or assumes any legal liability or responsibility forthe accuracy, completeness, or usefulness of any information, apparatus, product, or process disclosed, orrepresents that its use would not infringe privately owned rights. 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Box 62, Oak Ridge, TN  37831 Prices available from (423) 576-8401 http://apollo.osti.gov/bridge/  Available to the public from the National Technical Information Service U.S. Department of Commerce 5285 Port Royal Rd., Springfield, VA  22161 http://www.ntis.gov/  OR  Lawrence Livermore National Laboratory Technical Information Department’s Digital Library http://www.llnl.gov/tid/Library.html  Correlating observations of deformation microstructures by TEM andautomated EBSD techniquesMukul Kumar, Adam J. Schwartz and Wayne E. KingLawrence Livermore National Laboratory, University of California, Livermore, CA 94550 (USA)AbstractThe evolution of the deformed microstructure as a function of imposed plastic strain is of interest as it providesinformation on the material hardening characteristics and mechanism(s) by which cold work energy is stored.This has been extensively studied using transmission electron microscopy (TEM), where the high spatial andorientational resolution of the technique is used to advantage to study local phenomenon such as dislocationcore structures and interactions of dislocations.  With the recent emergence of scanning electron microscope(SEM) based automated electron backscatter diffraction (EBSD) techniques, it has now become possible tomake mesoscale observations that are statistical in nature and complement the detailed TEM observations.Correlations of such observations will be demonstrated for the case of Ni-base alloys, which are typically non-cell forming solid solution alloys when deformed at ambient temperatures.  For instance, planar slip is dominantat low strain levels but evolves into a microstructure where distinct crystallographic dislocation-rich walls formas a function of strain and grain orientation.  Observations recorded using both TEM and EBSD techniques arepresented and analyzed for their implication on subsequent annealing characteristics.1.  IntroductionEngineering of the grain boundary microstructurein low to medium stacking fault energy (SFE) FCCmaterials, as for instance Cu and Ni-base alloys iseffectively achieved by thermomechanical processing[1-3]. One approach has been a multi-cycle treatmentof moderate strain levels (5-30%) with annealingtreatments at relatively high temperatures but for veryshort times, typically on the order of 5 to 30 minutes.An important aspect is that the total formingreduction is broken up into several cycles of strainand annealing. The microstructure thus obtainedcontains a higher fraction of special boundaries ascompared with the as-received condition.This is unlike general industrial practice of largecold reductions followed by a recrystallizationanneal, or hot-working. In this case the annealedmicrostructure after single-step deformation on theorder of the cumulative strain from the multi-steptreatment does not show such dramatic differenceswhen compared with the as-received condition [3]. Itis well understood that a material after undergoingdeformations on the order of 50-70% (or higher)completely recrystallizes within a short period oftime on thermal treatments above 0.5Tm. The processof nucleation sites being randomly seeded in thedeformed microstructure as a consequence of thelarge strains is generally accepted.Current understanding of the evolution of themicrostructure and the mechanism(s) responsible forthe increase in the grain boundary characterdistribution (GBCD) after several moderate strain andannealing cycles is rather empirical in nature. Inparticular, the deformation microstructure atmoderate strain levels of 5-30% has not been wellcharacterized for non-cell forming FCC materials.Therefore, the objective of this study is to investigatethe deformation microstructure as a way of betterunderstanding the annealing behavior.2.  Experimental procedureTEM and EBSD investigations were conducted onsections perpendicular to the compression axis toenable determination of the orientation dependence ofslip in the polycrystalline samples of a Ni-16Cr-9Fealloy. These observations provided information atdifferent length scales. The typical resolutionobtained with the EBSD technique is about 1 µm,while features can be routinely resolved to about 1-2nm in the TEM. Similarly, the orientationalresolution is better in the TEM as compared to theEBSD technique, but not by orders of magnitude asin spatial resolution. Instead the comparison showsthat the resolving power of EBSD is about 1°, whichcompares with about 0.1° in the TEM.Individual dislocations are not sampled in EBSD(as in the TEM) but rather lattice rotations that ariseas a consequence of deformation are indicated asorientations that deviate from the referential point.Thus point-to-point correlations (misorientations) areobtained. This data can then be converted intoinformation on texture and the character of grainboundaries. As an extension, the effects of latticerotations (due to deformation) near grain boundarieson the character of the grain boundary, as defined bythe coincident site lattice (CSL) model [4], can beeasily studied [5,6].3.  ObservationsThe dislocation substructure as a function of roomtemperature compressive strain is shown in the TEMmicrographs of Fig. 1. After 5% strain, planar arraysof dislocations were observed (Fig. 1a). This wasindicative of negligible cross-slip in the alloy, whichis a concentrated Ni-base solid solution and in alllikelihood has a low SFE (< 70 mJ/m2) [7]. Inaddition, there was minimal evidence for interactionsbetween neighboring slip bands. Further deformationto about 15% strain revealed a well-developed Taylorlattice type dislocation substructure (Fig. 1b).Interactions between the planar arrays of dislocationsdid indicate the incipient stages of deformationbanding. This was observed to be more pronouncedafter 25% (Fig. 1c), and almost all of the grainsexhibited some evidence of deformation banding at45% strain. The widths of the bands, or domains ascalled by Hughes [8], typically were in the range of200-250 nm. No evidence for a cellular sub-structurewas seen within or outside these bands at the strainlevels of 5-45%.The deformation bands, as seen in Fig. 1c, did notparticularly form near grain boundaries. Based ondetailed observations that were made of thisdeformation microstructure it appeared that thedislocation rich walls formed on the primary slipplane (dotted line in Fig. 1c). This was also verifiedby diffraction where it was evident that themisorientations across the walls are of the order of 1°.A statistical sampling of the deformedmicrostructure revealed that deformation (after 5%strain) within the grains was pronounced in regionsnear the grain boundaries, as evidenced by therotations that are mapped by EBSD in Fig. 2b. Therotations (Fig. 2b) within the grains could be due todeformation incompatibility or formation ofdislocation walls that manifest as misorientations.These are indicated by the fine grayscale segments(1-15°) observed near the high angle grain boundariesthat are shown in bold. In comparison, the EBSDmap from the undeformed, well-annealed sample isFigure 1.  TEM micrographs of dislocation structureafter a) 5%, b) 15%, and c) 25% strain.shown in Fig. 2a.Significantly more rotation was observed byEBSD mapping after 15% strain. The EBSD maps inFig. 3 illustrate substantial slip activity in all grainsand well within the grain interiors. The maps showspecial and random high angle grain boundariesmapped in black and light grayscale color,respectively. An interesting observation was of highangle boundaries changing character as aconsequence of lattice rotations in their vicinity.Examples of have been marked as ‘A’ in Fig. 3a,where it is observed that the random boundarybecomes special in character as its deviation fromBrandon criterion [9] changes locally. Specialboundaries locally converting to random characterwere also observed, and one such example is markedas B in Fig. 3b.Low-angle boundaries between 1 and 2° mappedas fine grayscale segments (in Fig. 3a) are observedat a higher frequency when compared with Fig. 2band are uniformly distributed. The interesting point isthat at this strain level the frequency of low-angleboundaries between 2-5° (mapped in Fig. 3b)increases dramatically over that of the 5% sample,where practically none were observed. Theseobservations are in agreement with the TEM imagesof Fig. 1, which suggest that regions of local rotationappear within the grains with increasing strain. Thelength scale of EBSD mapping would suggest thatseveral of the domains (typically 250 nm in width)separated by dislocation rich walls (or domainboundary [6]) could be sampled as one singularregion of high misorientation.Formation of low-angle boundaries of angularrotation greater than 2° was correlated with anorientation dependence. It was generally observedthat such rotations were observed in grains where thestress axis was closer to the <110> or <111>directions. Grains with an axis near <100> typicallyhad higher densities of misorientations close to 1°.These observations were further investigated byTEM. An example has been taken from a sampledeformed 15%, where the stress axis is perpendicularto the plane of view. It is quite apparent that thedeformation microstructure, as seen in Fig. 4, oneither side of the grain boundary is vastly different. Inthe case of the top grain with only two active slipsystems (hard deformation direction, ~<111>)numerous dislocation rich walls (as in Fig. 1) alongthe primary slip planes could be observed. PresenceFigure 2.  EBSD maps showing lattice rotations withingrains in a) solutionized sample, and b) after 5% strain.Figure 3.  EBSD maps showing rotations within grainsof a) 1-2°, and b) 2-15° after 15% strain.of these walls lead to spot splitting in the diffractionpattern tilted almost perpendicular to the grain axisthus indicating substantial local rotations. Thedeformation microstructure appeared to be moreuniform in the bottom grain with multiple possibleslip systems (soft deformation direction, ~<001>) andno discernable rotations within the grain. In addition,deformation bands, as at position ‘A’, could also beobserved in the top grain in the vicinity of the grainboundary. From diffraction evidence taken acrossseveral points along the boundary, it was possible toascertain that the locally the segment of the boundaryhad changed character from a Σ91d (<111>/54°) toΣ37c(<111>/50.5°). This is agreement with the EBSDobservations of Fig. 3 and points to the localfluctuations in the deformed microstructure [5,6] thatneed to be investigated in order to predict theannealing behavior.Grain sub-division into regions of misorientationsover a wide range of angles was postulated even fornon-cell forming materials [10] and experimentallyverified in an Al-5.5at%Mg alloy [8]. It has beenshown in this study that a combination of TEM andEBSD is an effective way of extracting informationon the evolutionary process of such deformationmicrostructures. TEM investigations enabledidentification of the character of individualdislocations and their participation in planar slipprocesses that evolve into dislocation tangles ororganize themselves into the so-called Taylor lattices[10]. This process was also captured by EBSDobservations, albeit at a length scale close to theresolution limit of the technique. This enabled astatistical sampling of the collective behavior of anensemble of dislocations that manifests itself aslattice rotations due to the presence of geometricallynecessary dislocations.The arrangement of dislocations informs us aboutthe storage of cold work. For instance, in the presentstudy it is clear that new high angle boundaries thatcould serve as recrystallization nuclei are notintroduced. This has indicated that in the process ofgrain boundary engineering, where the material isonly deformed 25% strain in each cycle [1-3],annealing in all likelihood proceeds by boundarymigration mechanisms that rely on local variations inthe stored energy of cold work [11].AcknowledgmentsWork was performed under the auspices of the U.S.Dept. of Energy at Lawrence Livermore NationalLaboratory under contract no. W-7405-Eng-48.References1. G. Palumbo, U.S. Patent No. 5702543 (1997) and5817193 (1998).2. M. Kumar, W. E. King, and A. J. Schwartz, ActaMater., 48 (2000) 2081.3. A. J. Schwartz, M. Kumar, and W. E. King, MRSSymp. Proc., 586 (2000) in press.4. H. Grimmer, W. Bollmann, and D. H. Warrington,Acta Crystallog., A30 (1974) 197.5. V. Randle, N. Hansen, and D. Juul-Jensen, Philos.Mag. A, 73 (1996) 265.6. S. Sun, B. L. Adams, and W. E. King, Philos. Mag. A,80 (2000) 9.7. M. Kumar and V. K. Vasudevan, Acta Mater., 44(1996) 4865.8. D. A. Hughes, Acta Metall. Mater., 41 (1993) 1421.9. D. G. Brandon, Acta Metall., 14 (1966) 1479.10. D. Kuhlmann-Wilsdorf, Mater. Sci. Eng., A113 (1989)1.11. J. E. Bailey and P. B. Hirsch, Proc. Royal Soc., 267A(1962) 11.a)b) c)Figure 4.  a) TEM micrograph of a grain boundaryregion showing the formation of dislocation walls in agrain with hard orientation, b). Stereographic projectionof c) shows the bottom grain in a soft orientation.

Correlating Observations of Deformation Microstructures by TEM and Automated EBSD Techniques

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Correlating Observations of Deformation Microstructures by TEM and Automated EBSD Techniques

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