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Acta Materialia 60 (2012) 4311–4318

Observation of recovery and recrystallization in high-purityaluminum measured with forward modeling analysis of

high-energy diffraction microscopy

Christopher M. Hefferan a, Jonathan Lind a, Shiu Fai Li a, Ulrich Lienert b,Anthony D. Rollett c, Robert M. Suter a,⇑

a Department of Physics, Carnegie Mellon University, Pittsburgh, PA 15213, USAb Advanced Photon Source, Argonne National Laboratory, 9700 South Cass Avenue, Argonne, IL 60439, USAc Department of Materials Science and Engineering, Carnegie Mellon University, Pittsburgh, PA 15213, USA

Received 3 February 2012; received in revised form 12 April 2012; accepted 15 April 2012

Abstract

High-energy X-ray diffraction microscopy is a non-destructive materials characterization technique that is capable of tracking the evo-lution of three-dimensional microstructures as they respond to external stimuli. We present measurements of the annealing response ofhigh-purity aluminum using the near-field variant of this technique. The data are analyzed with the forward modeling method whichproduces orientation maps that exhibit complex intragranular structures. Analysis and verification of results use both reconstructed sam-ple space maps and detector space intensity patterns. Sensitivity to the ordering of the microstructure through both recovery and recrys-tallization is demonstrated. Sharpening of diffraction peaks and a corresponding reduction in intragranular orientation variations signalrecovery processes. The emergence of a new bulk grain (recrystallization) is observed in a disordered region of the microstructure; thenew grain has an orientation with no obvious relation to those of grains surrounding the disordered region.� 2012 Acta Materialia Inc. Published by Elsevier Ltd. All rights reserved.

Keywords: Microstructure; Recovery; Recrystallization; X-ray diffraction; Forward modeling

1. Introduction

The annealing response of deformed microstructures canbe separated into three classes of process, namely recovery,recrystallization and grain growth [1]. Distinct in their spe-cific mechanisms and final state characteristics, all threestages use thermal energy to evolve a polycrystal through asequence of increasingly ordered states. The available energyallows local changes in dislocation density during recovery,the nucleation of new grains in recrystallization, and themotion of grain boundaries to reduce total interfacial areaduring grain growth or coarsening. In practice, all threeresponses can be present concurrently. Study of these stagesof annealing is a mature field and has been exhaustively

1359-6454/$36.00 � 2012 Acta Materialia Inc. Published by Elsevier Ltd. All

http://dx.doi.org/10.1016/j.actamat.2012.04.020

⇑ Corresponding author. Tel.: +1 412 268 2982; fax: +1 412 681 0648.E-mail address: suter@andrew.cmu.edu (R.M. Suter).

analyzed from both theoretical and experimental perspec-tives. Nevertheless, the ability to directly observe theseresponses within bulk material has not been possible dueto the limitations of available instrumentation. Now, withthe advent of non-destructive synchrotron-based micro-structure characterization techniques, observations of inter-nal polycrystal responses to external stimuli are possible [2–9].

While all three classes of annealing response have beenstudied with synchrotron X-ray techniques, we focus hereon recovery and recrystallization. Existing observations[10–16] clearly establish the importance of detailed, non-destructive probes. These studies, however, are limited tothe tracking of small numbers of Bragg peaks or lack infor-mation regarding the microstructural neighborhood of themonitored grains. We present a three-state, in situ annealingexperiment using a polycrystalline wire of high-purity(99.999%) aluminum. The measurement, conducted at the

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4312 C.M. Hefferan et al. / Acta Materialia 60 (2012) 4311–4318

Advanced Photon Source (APS) beamline 1-ID, uses near-field high-energy X-ray diffraction microscopy (nf-HEDM)[2–4] to map crystallographic orientations over a volume ofmaterial containing thousands of distinguishable orienta-tions. Through subtle, low-temperature annealing, weobserve recovery as ordering within existing domains of sim-ilar orientation (“grains”) and recrystallization as the emer-gence of new measurable orientations from regions ofrelatively high disorder.

nf-HEDM [2–4] is used to generate 2-D maps of crystal-lographic orientations with �5 lm spatial and 0.1� orienta-tion resolution. The maps can be extended to threedimensions by measuring consecutive cross-sections in aprocess analogous to 3-D electron backscatter diffractionmicroscopy (EBSD) surface measurements [17]. Unlike 3-D EBSD, nf-HEDM is non-destructive and therefore allowsfor the repeated measurement of the same volume of micro-structure as it evolves when subjected to external stimuli.The measurement is based on high-resolution imaging ofBragg scattering from individual crystals in a monochro-matic, rotating sample scheme. Images are collected intransmission as the sample is rotated about the normal toa line-focused high-energy beam. The rotation is parameter-ized as x and images are integrated over an interval dx. Thenear-field distinction refers to the distance between sampleand X-ray detector where the diameter of the specimen isof the same order of magnitude as the sample-to-detectordistance. At this length scale, the position of diffractionpeaks on the detector is determined by the combination ofdiffracted beam propagation direction and the position oforigin in the sample.

The conversion of thousands of detector images to realspace maps of the crystallographic orientation field isaccomplished using the forward modeling method [3,18–20]. This method uses a computer simulation to optimizethe agreement between experimental data and simulatedscattering from crystallographic orientations assigned toa finely gridded tessellation of the sample space. This dis-cretization of sample space, on a length scale much smallerthan the grain size, leads to measurement sensitivity tointragranular orientation variations. Specifically, the for-ward modeling reconstruction connects each sample spacevolume element (reconstruction mesh element or voxel)with many Bragg reflections in the 4-D detector space,parameterized by pixel addresses on the detector images,j, k, the rotation axis-to-detector distance, L, and the rota-tion interval position, x. In the present work, the sample,with its complex intragranular orientation fields, generatesdetector images of diffracted beams that are similarly com-plex and difficult to interpret on a one-by-one basis. Weillustrate the efficacy of correlating sample space and detec-tor space representations.

2. Experimental procedure

The aluminum specimen was a cylindrical wire mea-suring 1 mm in diameter and 30 mm in length. Since

the wire was formed by drawing, the microstructurehad appreciable defect content. In situ annealing wasaccomplished using a fabricated boron nitride (BN)receptacle, with a 10 mm length of the sample embedded.The BN was resistively heated. The measured volume ofmicrostructure was in air, approximately 19 mm from theBN, and therefore subject to a thermal gradient, thoughthe sampled region was small enough to be approxi-mated by a uniform temperature. The initial state wasmeasured as received, with no thermal treatment. Thetwo annealing treatments achieved temperatures of �50and �70 �C, respectively, after 10 �C incremental stepsover �2.25 h. These temperatures are approximate sincecalibration of temperature at the measurement positionwas performed after the measurement. The HEDM datacollection consisted of the measurement of 11 2-D cross-sections with 20 lm spacings along the wire axis. Consis-tent mapping of the same 11 layers was achieved bymonitoring the intersection of the line-focused X-raybeam with the sample tip, using a diode detector situated�2 m downstream of the sample. The sample returned towithin a few microns of its initial position after bothanneals, permitting repeatable measurement of the samevolume of microstructure. In-plane alignment of data setsby a few microns optimized overlap of the differentstates.

Diffraction data were collected at an operating energy of50.1 keV. The line-focused beam is <6 lm high and1.3 mm wide, and therefore an entire cross-section of thecylindrical sample is illuminated. Diffraction spots areimaged at three rotation axis-to-detector distances ofL = 5.1, 7.1 and 9.1 mm. The detector system was comprisedof a lens-coupled CCD camera with a 4.2 mm square field ofview of a thin Ce-doped surface scintillation layer on a YAGsingle crystal. The effective pixel size is 4.11 lm. At each L

distance, a series of 100 diffraction images were collected asthe sample rotated over dx = 1� intervals. These imagesare split between two contiguous 50� regions, with the cen-ters offset by 90�. This set of intervals permits sufficient sam-pling of spot geometries for spatial reconstruction and spotlocations for orientation determination. Good signal-to-noise was observed for scattering up to the {440} family ofBragg peaks.

The reconstruction of the nf-HEDM diffraction data isaccomplished with the forward modeling method, whichassigns an orientation field to each X-ray-defined samplelayer. The forward modeling software [3,18–20] tessellatesthe sample space with a grid of equilateral triangles. Simu-lated scattering from each voxel is compared to observedscattering as orientation space is searched. Thus, eachvoxel is optimized against the complete data set. Broadcoverage of x leads to isotropic orientation and spatial res-olution as well as signal averaging of noisy and weak scat-tering. Over the measured range of x, arbitraryorientations generate �35 distinct Bragg peaks and theoptimization in each voxel is based on comparison withthe complete set.

C.M. Hefferan et al. / Acta Materialia 60 (2012) 4311–4318 4313

3. Results

Reconstructions of a single microstructure layer (thecenter layer of the measured volume) in each of the threeanneal states are shown in Fig. 1. The reconstructions usea uniform grid of 5.1 lm-sided equilateral triangles witheach triangle assigned an orientation independently of itsneighbors. This mesh size has been selected to approximatethe size of detector pixels, which limits the spatial resolu-tion of the measurement. Moderate- to high-angle grainboundaries are visible as boundaries between regions ofvisibly distinct color. Such boundaries with misorientationangle, hb P 2�, are decorated by the black lines in the fig-ure. Since these lines are determined locally (from near-est-neighbor voxel misorientations), the fact that narrowand continuous lines enclose uniformly colored regionsindicates that the noise level or uncertainty in determiningthe boundary locations is at or below the grid resolution ofthe reconstruction. Since most uniformly colored regionscontain no black lines, very few internal nearest-neighborvoxels are determined to have orientations that differ byas much as 2�. Note, however, that regions that appearto be uniform in color are not necessarily uniform in orien-tation since the color scale is too coarse to visualize smallvariations.

Also shown in Fig. 1 are maps of the confidence metric,C. Confidence is defined for each voxel as the fraction ofsimulated Bragg peaks generated by the optimized orienta-tion that overlap observed diffraction at more than onedetector distance, L [3]. Only voxels with C P 0.25 are

(a) (b)

(e)(d)

Fig. 1. Maps of the center layer of the measured volume. (a–c) Show orientatiomesh triangle, or voxel, is colored by mapping orientation components (Rodriglines are drawn between adjacent triangles that have misorientations P2�. (d–fmaps in (a)-(c) with the same black lines included. The confidence maps are plopeaks overlap observed scattering. The arrow in (c) indicates the nucleated gr

considered to have converged to a (reasonably) reliable ori-entation and are included in Fig. 1.

Examination of Fig. 1a–c demonstrates that mostboundaries with hb P 2� exhibit little motion duringannealing and that the measurement (both sample align-ment and reconstruction) is reproducible to the resolutionof the grid. Careful examination does reveal some subtleboundary motions and the emergence of a few new orien-tations. We conclude that only minimal grain growth orcoarsening occurs during this measurement. On the otherhand, Fig. 1d–f exhibits increasing average values of C withnumerous regions of low confidence disappearing in thefinal state. The low C regions are predominantly alonggrain boundaries and in the vicinity of triple lines. Manyof the regions extend to neighboring layers and the sameordering trends are visible in all measured layers. Regionswith low C appear to be regions of high disorder in whichthe orientation within our reconstruction voxel length scalevaries appreciably, yielding broad and weak scattering. Thereconstruction algorithm tends to extend the domain ofnearby orientations but with significantly reduced overlapbetween simulated and experimental peaks. As annealingproceeds, these regions become more ordered, the scatter-ing becomes stronger, and better orientation informationis obtained. Below, we discuss details of the orientationand confidence maps in order to extract information aboutrecovery and recrystallization processes that underlie theobserved evolution. We compare simulated scatteringdirectly to measured signals in order to corroborate subtlefeatures and trends.

(c)

(f)

n maps in the initial, �50 �C and �70 �C anneal states, respectively. Eachues vector representation) to the red–green–blue (RGB) color space. Black) Show maps of the confidence metric, C, corresponding to the orientationtted on the same color scale. C = 1 (red) means that �35 simulated Bragg

ain discussed in Section 3.2.

(a) (b) (c)

Fig. 2. Kernel-averaged misorientation (KAM) maps of the same three layers of microstructure shown in Fig. 1. The gray scale saturates at KAM = 0.1(white) and KAM = 0.5 (black) degrees. All three maps use the same gray scale.

4314 C.M. Hefferan et al. / Acta Materialia 60 (2012) 4311–4318

3.1. Recovery

The recovery process is characterized by changes in dis-location density and distributions that result in the reduc-tion of local orientation variations. Crystal orientationvariations correspond to a distribution in the directionsof local reciprocal lattice vectors, and therefore recoveryis observable as a sharpening of intensity distributions.Similar orientation distributions are conventionally mea-sured in rocking curve measurements in single crystals[21]; in effect, we measure such curves for many crystalsusing many Bragg peaks for each. In our spatially resolvedmeasurements, as grains move towards more uniform ori-entations, diffraction peaks appear at shorter x intervals,have reduced arc lengths on a given detector image, andhave more uniform intensity across peak profiles. By usingforward modeling’s voxel-by-voxel (rather than grain-by-grain) reconstruction, we are able to follow this evolution.Again, this capability relies heavily on the inclusion ofmany diffraction spots over a wide range of x and theresultant high spatial and orientational resolution.

Fig. 2 shows a measure of local orientation variation,the kernel-averaged misorientation or KAM [22]. Here,KAM is calculated as the average misorientation betweena voxel and its nearest neighbors. Since the in-plane mapspatial resolution (5.1 lm) is finer than the layer-to-layerspacing (20 lm), only in-plane contributions are included.In the figures, points with KAM > 2� are suppressed andshown as the grain boundary lines in red.1 Note the highlyexpanded gray scale in Fig. 2. All voxels with sub-orienta-tion resolution values (KAM < 0.1�) are shown as white,while 0.5� < KAM < 2� are black. This scale allows oneto pick out regions of relatively high disorder (clusters ofdark points) as well as seeing light regions with littleshort-range disorder. Generally, KAM plots do not neces-sarily show low gradients in orientation ([0.1� over 5 lm).In all three sample states, one can pick out noisy regions aswell as ordered 1-D structures that correspond to low-angleboundaries. Some of these boundaries have ends that aredisconnected from other boundary lines, implying asmooth orientation gradient around the end point. Many

1 For interpretation of color in Figs. 2–4, 6, the reader is referred to theweb version of this article.

of these features can be correlated across the three annealstates of the sample, thus reinforcing the reliability of thereconstructions. One also sees regions in which orderedstructure grows rather dramatically. The near-surfaceregion appears to be particularly active in this behavior,indicating the limitations of recovery studies using sur-face-based techniques.

While the KAM maps in different anneal states are qual-itatively similar, there is an overall lightening of the imagesas annealing progresses, implying decreased average KAM

within the sample. This is illustrated in Fig. 3 as histogramsobtained from the entire set of 11 measured layers in eachstate. The population in the sub-resolution (<0.1�) bingrows significantly, while large KAM values decrease inpopulation. This growth of orientational order at shortlength scales is consistent with expected recovery behavioras noted above.

To more directly tie the KAM trends to the experimentalsignals, we accumulate statistics on peaks widths in the xdirection (“rocking curves”). To accomplish this we usethe scattering generated by the optimized orientation fieldto identify Bragg peak locations and intensity patterns cor-responding to real-space regions defined by the P2� bound-aries shown in Figs. 1 and 2. Defining these regions to be“grains,” we count the number of measured x intervals(each 1� wide) in which each peak occurs in the experimentalimages (not in the simulation). The 96 grain cross-sectionsthat maintain roughly constant shapes throughout anneal-ing are chosen and 3022 distinct Bragg peaks are studied.The cross-sections are from three distinct layers separatedby 100 lm. The average widths are 3.6, 3.4, 3.1 degrees inthe three states. Fig. 4 shows the peak width evolution asseen in histograms of changes in widths between annealstates. The higher population of negative values indicatesa statistical sharpening of the scattering. The actual valuesof the average widths reflect orientation variations but areincreased by the inefficiency of the x rotation in moving reci-procal lattice vectors that are close to that axis through theBragg condition (as reflected in the Lorentz factor [2]).

3.2. Recrystallization

While recovery is a process defined by subtle, continu-ous changes, recrystallization is a dramatic (critical)

0 0.1 0.2 0.3 0.4 0.5 0.60

0.1

0.2

0.3

0.4

0.5

Kernel Averaged Misorientation (degrees)

Frac

tion

of V

oxel

s

Anneal 0Anneal 1Anneal 2

Fig. 3. KAM histograms for the three anneal states. Horizontal bin widthsare 0.1�. Points are average values of histograms computed for each of the11 layers and error bars are corresponding standard deviations.

−4 −3 −2 −1 0 1 2 30

500

1000

1500

2000

Change in Peak Width (degrees)

Num

ber

of P

eaks

Anneal 1 − InitialAnneal 2 − Initial

Fig. 4. Histogram of changes in peak widths in experimental detectorspace as a function of anneal. Each pair of bars refers to changes in theinteger numbers of x intervals in which above background intensity isobserved.

C.M. Hefferan et al. / Acta Materialia 60 (2012) 4311–4318 4315

microstructural event characterized by the emergence of anew grain or orientation. New grains are detectable bythe appearance of new sets of crystallographically consis-tent Bragg peaks in the 4-D detector space. The size of thisspace, coupled with its complex population of polycrystal-line diffraction peaks, makes identification of new crystalschallenging. Using a voxel-based reconstruction, small fea-tures can be identified. Post-reconstruction, identificationof new grains can be automatically accomplished withpoint-by-point comparisons of the microstructure in multi-ple sample states, assuming only subtle motions of grainboundaries. As above, examining scattering that is consis-tent with the new orientation and its location can link thenew grain to its source signal (experimental diffraction) inorder to validate its presence and to verify its absence inearlier sample states.

The probed sample volume had a significant number ofnucleation events at the surface (see Fig. 1a–c). However,we limit detailed discussion to a single internal grain whichis fully contained in our measured volume. The grain isindicated by the arrow in Fig. 1c. Fig. 5 shows the nucle-ation site in each of the three anneal states. The eventoccurs in the middle layer of the measured volume. Thenew orientation is seen in two neighboring layers in the firstanneal state, and extends into three adjacent layers(±20 lm displaced) in the final state. Note that the grainemerges from a region with low C in the initial state(Fig. 5d), implying that it nucleated in a highly defectedregion. In fact, even in the final state, essentially the entiregrain occupies this initially defective region. Growth isalmost entirely within the deformed region rather than intothe neighboring, pre-existing grains. While the new graingrows, the surrounding ones become more ordered as indi-cated by the more uniformly red C map. Fig. 5g–i showsthat the new grain grows with only small internal KAM.Evidently, the new grain replaces high-energy–density, dis-ordered material with a low stored energy, ordered crystalstructure.

The images of three example peaks shown in Fig. 6 illus-trate the evidence for the new grain in the first and finalanneal states. The peaks are well defined in x space, cover-ing only a single (or, in rare cases, a pair of adjacent)dx = 1� bins. The simulated scattering generated by thevoxel-by-voxel optimized orientation is outlined over theexperimental intensities. These x intervals are 3 of 31where simulated scattering is generated. Of these, 28 arefound to overlap experimental scattering, which leads toC = 0.9. The initial state is not shown since there is no dis-cernible intensity in the corresponding images. In the twoadjacent layers in the final state, weak intensity is observedat each of these locations and is overlapped by simulatedscattering. With only three measured cross-sections of thisgrain, due to the 20 lm layer spacings, a meaningful 3-Dgrain shape cannot be generated.

We see no clear origin for the orientation of the newgrain; this somewhat surprising result [23] is consistent withearlier work by Larsen et al. [15]. No special orientationrelations are discernible between the nucleated grain andits neighborhood. Table 1 lists the misorientations inaxis-angle form relative to the crystal axes of the new grain.The first six entries are indicated in Fig. 5, while grain 7 isdetected in one of the adjacent layers. Grains 1–6 are sep-arated by more than 3� from the nearest coincidence sitelattice (CSL) relation; this statement includes CSL num-bers as high as 73. Grain 7 has rotation angle close tothe R23 CSL but the axis is far from the appropriateh311i direction. Further, we measure no diffraction signalsin the initial state at the detector locations struck by thenew grain in later states (such as those shown in Fig. 6).Manual examination of raw detector images, without anybackground treatment, fails to locate intensity peaks atthese and other relevant detector positions. Thus, the disor-dered region where the grain nucleates does not possess

Fig. 5. Maps illustrating the nucleation of a new grain during annealing. Columns from left to right show initial, first and final anneal states. (a–c) Showorientations as in Fig. 1 along with outlines of the new grain as observed in the second and final states. (d–f) Show maps of C in the same region. (g–i)Show KAM values. Misorientations between the nucleated grain and its neighbors are discussed in Section 3.2.

4316 C.M. Hefferan et al. / Acta Materialia 60 (2012) 4311–4318

sufficient orientation content at the nucleated grain’s valueto generate observable experimental signal.

4. Discussion

Simultaneous recovery and recrystallization phenomenaare observed in situ in low-temperature (<100 � C) annealingof a high-purity aluminum wire using nf-HEDM. Recoveryis evidenced by a statistical sharpening of diffraction peakrocking curves that is reflected in subtle reductions of orien-tation gradients in real space reconstructions of the orienta-tion field. With orientation resolution of � 0.1�, subtlechanges are observed. Recrystallization is observed by theemergence of new orientations in regions of prior disorder(both in the bulk of the sample and at numerous surface-connected regions). A spatial resolution of several micronsallows the analysis of nucleating regions prior to the eventin order to search for mechanisms associated with the emer-gence of specific orientations. Spatial resolution also boundsthe average motion of grain boundaries, which in the currentwork is comparable to the spatial resolution of �5 lm.

The ability to observe subtle intragranular orientationstructure, such as in Fig. 2, originates from the measure-ment of many Bragg peaks from arbitrary orientations.

An orientation places these peaks in specific x bins (and,therefore, specific detector images) where the Bragg condi-tion is satisfied. This occurs when the x rotation orients thereciprocal lattice vector, G, such that ki � G = �0.5G2,where ki is the incident X-ray wavevector which is fixedin the laboratory coordinate system, as is the rotation axis.The reciprocal lattice vectors rotate with the sample. Bin-ning �30 distinct peaks strongly constrains the orientation.If the lattice within a nominal grain rotates by even a frac-tion of the integration interval, dx = 1�, it is likely thatsome subset of the scattering will be shifted to different xintervals. In contrast to simply shifting intensity distribu-tions within the nominal grain’s projected images, thisshifting to distinct images is easily followed by the voxel-based reconstruction code.

A caveat to literal interpretation of images such as inFig. 2 is that structure exists on length scales much smallerthan the resolution of 5 lm in this work. The orientationsdeduced from the present diffraction data are, presumably,averages of orientations within the sample space voxels.When sub-resolution orientation variations become toobroad, the scattering from a voxel becomes both broadand weak. Lattice rotations about different axes generateanisotropic peak broadening that results in some peaks

Fig. 6. Bragg scattering generated by the recrystallized grain of Fig. 5.Rows show three distinct peaks using identically sized detector regions,while columns correspond to different anneal states. All images use thesame intensity gray scale. Raw experimental data have been filtered andbackground subtracted. The red outlines define the regions covered bysimulated scattering from the reconstructed grain.

Table 1Misorientation relations of grains in the neighborhood of the nucleatedgrain. See the text for definitions of grain identifier numbers.

Number Angle (�) Axis

1 14.4 �0.559 �0.822 0.1082 12.3 �0.585 0.054 �0.8093 34.3 �0.263 �0.057 �0.9634 36.7 �0.239 0.068 0.9695 33.9 �0.554 0.293 0.7796 24.9 0.012 0.160 0.9877 40.9 �0.093 0.235 �0.968

C.M. Hefferan et al. / Acta Materialia 60 (2012) 4311–4318 4317

falling below the detection limit. In the present data, theconfidence measure is seen to drop to low values in partic-ular sample regions. High KAM and low C can qualita-tively indicate “disordered” regions, as we have describedhere. Similarly, calculations of dislocation content (not per-formed here) will need to be taken as lower bounds, sincethe orientation field has been smoothed over the resolutionlength scale. Future data sets with spatial resolution of theorder of at least some cell structures will make such calcu-lations relevant to placing bounds on actual dislocationcontent and evolution.

With regard to recrystallization, the degree to whichgeneral conclusions can be drawn from a single event is cer-tainly limited. Nevertheless, we note the counterintuitivenature of the appearance of new orientations in recrystalli-zation [23]. As discussed above, no evidence for scattering

from the orientation of the new internal grain (and from itslocation) could be identified in the initial state data evenwhen the raw diffraction images were examined. Thisobservation is consistent with a single nucleating grainobserved by Larsen et al. [15]. That grain, in a copper sam-ple, is located near a triple junction and has an orientationwith no simple relation to the three adjacent grains. Ourobservation, made possible through more detailed spatialresolution, is that the new grain presented here grows outof a highly disordered region which is surrounded by sevendistinct grains. This location possibly avoids the highlyunfavorable nucleation threshold for a random orientationin contact with surrounding crystals [23]. Our measurementcannot, of course, say anything about the prior existence ofa subcell of appropriate orientation and of a size below ourdetection threshold. The low C region is likely to contain ahigh concentration of small-scale features; future higher-resolution measurements may be able to address suchstructures. At the annealing temperatures used here, high-angle (P2�) boundaries between existing grains elsewherein the microstructure move very little; therefore, it is notsurprising that the nucleated grain does not grow discern-ibly into existing well-ordered grains. Additional event sta-tistics are clearly needed.

5. Conclusion

An application of forward modeling analysis tonf-HEDM has been presented, demonstrating its ability tocharacterize subgrain evolution in the annealing of defectedpolycrystals. Through the combined analysis of recon-structed orientation maps in sample space and diffractionpeak characterization in detector space, we are able toidentify and validate salient features associated with bothmicrostructure recovery and recrystallization. Intragranularorientation variations that reflect geometrically necessarydislocation structures are observed to evolve toward moreuniform orientations both below and above the 5 lm spatialresolution limit of this measurement. The growth of new,well-ordered grains from disordered regions of the sampleis observed both in bulk and at the sample surface. All obser-vations are based on 3-D structures deduced from extensivedata sets that include many Bragg peaks from each sampleregion and crystallographic orientation. The non-destruc-tive nature of the measurement allows for the possibility oflooking at regions in which critical events or slowly evolvingstructural changes occur in sample treatment states prior tothe events. This retrospective ability eliminates the “lost evi-dence problem” associated with critical events, such as nucle-ation [24]. The examination of specific microstructuralregions of interest as seen in the reconstructions is comple-mented by examination of detector regions that are high-lighted by the reconstruction software.

The measurements reported here were carried out in2007. Since that time, many improvements in beamlinehardware at the APS have been implemented. Data collec-tion has been accelerated to the point where typical volu-

4318 C.M. Hefferan et al. / Acta Materialia 60 (2012) 4311–4318

metric data sets of 100 layer measurements are collected inless than 24 h (actual speed depends on signal strengthinherent to sample material and its condition). The stan-dard detector system now has 1.5 lm resolution insteadof the 5 lm pixels and resulting grid resolution used here.With the development of continuous scanning, integrationintervals dx < 1� can be collected without prohibitivelyincreasing data collection time. On the reconstruction side,development is underway to implement explicit matchingof simulated intensity distributions to those observedexperimentally. Extension of the current work, as well asstudies of microstructure evolution due to other drivingforces, are in progress.

Note added in proof

On further investigation, we have detected, in the initial,unannealed state, a very weak diffraction signal associatedwith the orientation and location of the emerging grain ofFig. 5; the signal exists only in a single layer which is adja-cent to that shown. With improved signal extraction fromthe raw images of diffraction peaks and with refined recon-struction parameters, our search algorithm converges tothis orientation in a set of disconnected but vicinal voxelslocated within the low confidence region. The diffractionand reconstruction are consistent with a barely resolved,not yet fully ordered domain within the highly deformedregion. Details will be elaborated in the PhD Thesis byC. M. Hefferan, Physics Dept., Carnegie MellonUniversity.

Acknowledgements

The authors thank D. Hennessy for assistance with datacollection. This work was supported primarily by the Met-als and Nanostructures and the Materials Research Scienceand Engineering Center programs of the National ScienceFoundation under Award Numbers DMR-1105173 andDMR-0520425, respectively. The research was also sup-ported in part by the National Science Foundation throughTeraGrid resources provided by Texas Advanced Comput-ing Center under Grant No. DMR080072. Use of the Ad-vanced Photon Source was supported by the USDepartment of Energy, Office of Science, Office of Basic

Energy Sciences, under Contract No. DE-AC02-06CH11357.

References

[1] Humphreys FJ, Hatherly M. Recrystallization and related annealingphenomena. Oxford: Elsevier; 1995.

[2] Poulsen HF. Three-dimensional X-ray diffraction microscopy. Ber-lin: Springer; 2004.

[3] Suter RM, Hennessy D, Xiao C, Lienert U. Rev Sci Instrum2006;77:1–12.

[4] Lienert U, Li SF, Hefferan CM, et al. J Miner Metals Mater Soc2011;63:70–7.

[5] Liu W, Ice GE, Larson BC, Yang W, Tischler JZ. Ultramicroscopy2005;103:199–204.

[6] Larson BC, Yang W, Ice GE, Budai JD, Tischler JZ. Nature2002;415:887–90.

[7] Johnson G, King A, Goncalves Honnicke M, Marrow J, Ludwig W. JAppl Crystallogr 2008;41:310–8.

[8] Ludwig W, Schmidt S, Lauridsen EM, Poulsen HF. J ApplCrystallogr 2008;41:302–9.

[9] Ludwig W, Reischig P, King A, Herbig M, Lauridsen EM, JohnsonG, et al. Rev Sci Instrum 2009;80:1–9.

[10] Gundlach C, Pantleon W, Lauridsen EM, Margulies L, Doherty RD,Poulsen HF. Scr Mater 2004;50:477–81.

[11] Lauridsen EM, Juul Jensen D, Poulsen HF, Lienert U. Scr Mater2000;43:561–6.

[12] Offerman SE, van Dijk NH, Sietsma J, et al. Science 2002;298:1003–5.[13] Lauridsen EM, Poulsen HF, Nielsen SF, Juul Jensen D. Acta Mater

2003;51:4423–35.[14] Schmidt S, Nielsen SF, Gundlach C, Margulies L, Huang X, Juul

Jensen D. Science 2004;305:229–32.[15] Larsen AW, Poulsen HF, Margulies L, Gundlach C, Xing Q, Huang

X, et al. Scr Mater 2005;53:553–7.[16] Lauridsen EM, Schmidt S, Nielsen SF, Margulies L, Poulsen HF,

Juul Jensen D. Scr Mater 2006;55:51–6.[17] Rollett AD, Lee S-B, Campman R, Rohrer GS. Annu Rev Mater Sci

2007;37:627–58.[18] Hefferan CM, Li SF, Lind J, Suter RM. Powder Diffr 2010;25:132–7.[19] Hefferan CM, Li SF, Lind J, Lienert U, Rollett AD, Wynblatt P, et al.

Comput Mater Continuat 2009;14:209–19.[20] Li SF. PhD thesis, Carnegie Mellon University, 2011.[21] Weissmann S. Characterization of lattice defects and concomitant

strain distribution. In: Barrett CS et al., editors. Advances in X-rayanalysis, vol. 35A. New York: Plenum Press; 1992.

[22] Brewer LN, Field DP, Merriman CC. Mapping and assessing plasticdeformation using EBSD. In: Schwartz AJ et al., editors. Electronbackscatter diffraction in materials science. New York: SpringerScience; 2009.

[23] Doherty RD et al. Mater Sci Eng A 1997;238:219–74.[24] Duggan B. ICOTOM 11, 1996; as quoted in [15].