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Acta Materialia 61 (2013) 5685–5693

Mechanism of abnormal grain growth in ultrafine-grained nickel

Sang-Hyun Jung a, Duk Yong Yoon b, Suk-Joong L. Kang a,⇑

a Materials Interface Laboratory, Department of Materials Science and Engineering, Korea Advanced Institute of Science and Technology,

Daejeon 305-701, Republic of Koreab Pohang University of Science and Technology, Pohang 790-784, Republic of Korea

Received 4 April 2013; received in revised form 30 May 2013; accepted 5 June 2013Available online 29 June 2013

Abstract

During annealing of ultrafine nickel powder compacts, abnormal grain growth (AGG) occurred with the formation of cube-shapedgrains. The observed AGG is similar to the late-stage abnormal grain growth (LS-AGG) that is commonly observed in electrodeposited(ED) Ni sheets. The cube shape of the abnormal grains with {100} facets was identical to that of late-stage abnormal grains in ED Ni. Asthe temperature increased, the time to the appearance of abnormal grains decreased. After impingement of abnormal grains, little furthergrain growth occurred, indicating stagnant grain growth behavior. The formation of abnormal grains in the present study, however, can-not be explained by the previously suggested AGG mechanisms, i.e. the S phase wetting mechanism and initial texture mechanism,because no amorphous phase containing S was present and no texture developed in our samples. In contrast, the observed grain growthbehavior can be well explained in terms of the coupling effect of the maximum driving force for grain growth and the critical driving forcefor appreciable migration of faceted boundaries. It is concluded that the AGG in ultrafine-grained Ni occurs via mixed migration mech-anisms, i.e. diffusion and interface reaction-controlled, of faceted boundaries with respect to the driving force.� 2013 Acta Materialia Inc. Published by Elsevier Ltd. All rights reserved.

Keywords: Abnormal grain growth; Ultrafine nickel; Faceting; Microstructural evolution; Mixed control model

1. Introduction

Ultrafine-grained (UFG) Ni, unlike microcrystalline Ni,exhibits superplasticity even at low temperatures compared[1]. The technological application of UFG Ni, however, hasbeen hampered due to the difficulty in retaining the ultra-fine grain structure as abnormal grain growth (AGG) gen-erally occurs during material processing and use. (In fact,AGG is a common phenomenon in most materials with aUFG structure.) Grain growth behavior in UFG Ni hasbeen widely studied by many groups [2–22]. During anneal-ing of electrodeposited (ED) UFG Ni sheets below 550 �C,AGG usually takes place in two stages, initial and late[2–5]. When initial-stage (IS) AGG occurs, the averagegrain size increases from several tens of nanometers to overseveral hundreds of nanometers [2–22]. Late-stage (LS)

1359-6454/$36.00 � 2013 Acta Materialia Inc. Published by Elsevier Ltd. All

http://dx.doi.org/10.1016/j.actamat.2013.06.010

⇑ Corresponding author. Tel.: +82 42 350 4113; fax: +82 42 869 8920.E-mail address: sjkang@kaist.ac.kr (S.-J.L. Kang).

AGG is generally characterized by an increase in averagegrain size from several hundred nanometers to over severaltens of micrometers [2–5]. The formation of abnormalgrains in ED UFG Ni has been attributed either to the for-mation of a sulfur-based liquid phase at the grain bound-aries [2,6] or to the presence of a texture in the sampleafter the electrodeposition [3,7–11].

The S phase wetting mechanism was first proposed byHibbard et al. [2]. They observed LS-AGG as well asIS-AGG in ED Ni with an initial grain size of several tensof nanometers. IS-AGG readily occurred as the sampleswere heated to 420 �C and the abnormal grains that ini-tially formed several hundred nanometers in size impingedupon each other within 30 s. These impinged abnormalgrains grew very slowly up to 3600 s, showing practicallystagnant grain growth (SGG). At the end of the SGG,some of the impinged abnormal grains grew into micron-sized cube-shaped secondary abnormal grains (LS-AGG).The authors suggested that the LS-AGG was a result of

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Fig. 1. Measured variation of oxygen partial pressure of wet H2 and wet99N2/1H2 with temperature. Variation of the calculated Ni/NiO equilib-rium oxygen partial pressure is also shown.

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boundary wetting of the S phase, which released a pinningforce for boundary movement and increased the boundarymobility. According to our experiments [23,24], however,the boundary velocity of a wet boundary is lower than thatof a dry boundary. This result suggests that the S phasewetting mechanism may not provide an explanation forthe AGG in UFG-Ni. Furthermore, the S-phase wettingmechanism cannot explain why AGG occurs in two stagesand the shape of the abnormal grains after LS-AGG iscubic.

The initial texture mechanism of AGG was first pro-posed by Xiao et al. [3]. These authors suggested that theoccurrence of LS-AGG was related to the presence of aninitial texture. There is agreement among a number ofresearchers [7–11] about the mechanism of initial texture,which was inherently introduced during the preparationof ED UFG Ni. When IS-AGG occurred, a texture witha ND//h111i state developed from a specific initial state(e.g. ND//h100i [7–9], ND//h113i [10] and ND//h114i[11]). Park et al. [7] suggested that IS-AGG was causedby a potential energy difference of 1 J mol�1 betweenND//h100i grains and ND//h11 1i grains. However, asthe total interfacial energy of UFG ED Ni is of the orderof several kJ mol�1 [25,26], the difference of 1 J mol�1 doesnot seem to be sufficient to induce AGG. Klement et al.reported the development of a texture from ND//h113i[10] or ND//h114i [11] to ND//h111i during IS-AGG.They suggested that the growth of abnormal grains withND//h111i was caused by a difference in the surface energyof face-centered cubic (fcc) metals, which was known toincrease in the order of planes {111}, {100}, {11 0} and{31 1} [27]. As will be described in Section 3.1, however,abnormal grains had {100} facets in our samples.

According to recent investigations on AGG in variousmetallic and ceramic systems [28–41], AGG is closelyrelated to the faceting of grain boundaries. Normal graingrowth occurs when grain boundaries are rough (atomi-cally disordered), whereas abnormal (or non-normal) graingrowth occurs when grain boundaries are faceted (atomi-cally ordered), either fully or partially. AGG in faceted sys-tems is attributed to the nonlinear migration of facetedboundaries [42,43].

A relationship between AGG and grain boundary facet-ing was also observed in Ni [29,30]. When the grain bound-ary structure of Ni gradually changed from faceted torough with increasing temperature and decreasing oxygenpartial pressure, the grain growth behavior changed fromabnormal to normal [29]. Although this observation wasmade in microcrystalline Ni, it suggests that AGG in Niof any grain size is closely related to grain boundary facet-ing. It would, therefore, be appropriate to look at the AGGin UFG Ni in terms of grain boundary faceting.

The purpose of this study is to identify the mechanism ofAGG in UFG Ni. To this end, ultrafine Ni powder com-pacts were used, unlike in previous studies with ED Nisheets. The Ni powder used in this experiment had a verylow S concentration of 0.001 wt.% (compared to ED Ni

which has a relatively high S of 0.085 wt.%) [2]. Thus, pos-sible effects of the S phase or initial texture, which havebeen proposed to be the causes of AGG in UFG Ni, couldbe excluded in the present study. Examination of our sam-ples also confirmed that there was no possibility of S meltformation and texture development during sintering. Graingrowth behavior in ultrafine powder compacts wasobserved at different sintering temperatures and oxygenpartial pressures. A strong correlation between grainboundary faceting (i.e. step-free energy variation) andgrowth behavior was found, which supports the previouslysuggested mixed control model of grain growth [42,44,45].

2. Experimental

Samples were prepared from ultrafine Ni powder(99.9 wt.% purity, 180 nm size, JFE Mineral Ltd., Tokyo,Japan). According to the producer’s data, the S contentwas <0.001 wt.%. The powder was granulated by passingit through a 100 lm sieve and lightly pressed in a steeldie into disks 9 mm in diameter and 4 mm thick. The diskswere isostatically compressed at 200 MPa for 10 min usinga cold isostatic press (Autoclave Engineers Inc., Erie, PA,USA). The compacts were placed on Ni spacers in an alu-mina crucible and sintered in a vertical tube furnace in analumina tube capped at one end at various temperatures(500–650 �C) and in two different reducing atmospheres(wet H2 and wet 99N2/1H2). Fig. 1 plots the variation ofthe measured oxygen partial pressure (P O2

) with tempera-ture in wet H2 and wet 99N2/1H2 together with the knownequilibrium P O2

of a Ni/NiO reaction. The oxygen partialpressure in the gases was lower than the equilibrium P O2

of an oxidation reaction. Nickel oxide was not present inthe samples after sintering. The samples were sintered usinga heating rate of about 30 �C min�1. The samples were heldfor various periods up to 6000 min and then pulled out ofthe furnace at the same rate as the heating rate.

The sintered samples were vertically cut and polished toa 1 lm finish, and etched in a solution of 10 ml acetic acid,5 ml nitric acid and 0–3 ml distilled water. The

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microstructures of the central region of the samples wereobserved by scanning electron microscopy (SEM; XL 30SFEG, Philips, Eindhoven, the Netherlands). For transmis-sion electron microscopy (TEM) observation (JEM3010,JEOL, Tokyo, Japan), the samples were ultrasonically cutinto 3 mm disks, mechanically ground to a thickness of100 lm, dimpled to a thickness of less than 5 lm, andfinally ion-milled. The microstructures were also observedby electron backscatter diffraction (EBSD) in a NOVA230scanning electron microscope (FEI Co., Eindhoven, theNetherlands) with an accelerating voltage of 20 kV and astep size of 2 lm. The collected data were analyzed usingTSL orientation imaging microscopy (OIM) Analysis 5software. The specimens for EBSD measurement were pre-pared by a conventional metallographic polishing method.

3. Results and discussion

The results are presented in two parts. The first partdescribes the microstructural evolution and its character-ization of sintered ultrafine Ni at 550 �C in wet H2 up to6000 min. The growth behavior is compared with thatobserved during annealing of UFG ED Ni sheets. Thevalidity of the previous AGG mechanisms (i.e. S phase wet-ting and initial texture mechanism) is discussed to explainthe present results. In the second part, the observedchanges in growth behavior at different sintering tempera-tures and atmospheres are presented and explained interms of the migration behavior of the faceted boundaries.

3.1. Abnormal grain growth in ultrafine Ni powder compacts

Fig. 2 shows SEM micrographs of Ni compacts sinteredat 550 �C in wet H2 for various times from 0 to 6000 min.When the sample was heated up to 550 �C and immediatelycooled down to room temperature (0 min sintering), prac-tically no grain growth occurred, although the sample

Fig. 2. SEM micrographs showing the microstructures of Ni samples sintered aThe inset SEM image in (c) shows an intergranular fracture surface of the 60

was almost fully densified, as shown in Fig. 2a. After20 min sintering, a few large abnormal grains appeared(Fig. 2b). As sintering proceeded, formation and growthof abnormal grains occurred, consuming the fine matrixgrains. Abnormal grains then impinged upon each other(Fig. 2c and d) and covered the sample surface almostentirely (Fig. 2e). After impingement of abnormal grains,hardly any grain growth occurred during extended sinter-ing up to 6000 min. The average grain size was 26.2 lmafter 600 min and 28.0 lm after 6000 min.

Before the impingement of abnormal grains, the bound-aries between the abnormal grains and the matrix grainswere well faceted, as shown in Fig. 2b–d. The regular shapeof an abnormal grain seems to be a cube, as shown in theinset of Fig. 2c.

Fig. 3a is a typical bright-field TEM micrographobserved along the h0 01i zone axis of an abnormal grain,showing the boundary between the abnormal grain and thematrix grains. Macroscopically, the boundary consists of a(100) plane of the abnormal grain. {10 0} planes of abnor-mal grains were consistently observed for more than 10boundaries, indicating that the migrating boundaries ofabnormal grains have {10 0} planes, consistent with theshape of the abnormal grains in the inset of Fig. 2c. Fora fraction of migrating boundaries, an apparently slightdeviation from a (100) plane is also observed, e.g. theregion indicated by a circle in Fig. 3a. However, such aboundary was observed to consist also of {100} facetsand many steps, as shown in the high-resolution (HR)-TEM micrograph in Fig. 3b. This observation implies thatthe growth of abnormal grains is essentially governed bythe migration of {100} boundaries and hence the possibleeffect of the local variation of the apparent boundary planeis insignificant.

The observed AGG behavior in Fig. 2 and the observedshape of abnormal grains are similar to the LS-AGGbehavior and the shape of abnormal grains in ED UFG

t 550 �C in wet H2 for (a) 0, (b) 20, (c) 60, (d) 180, (e) 600 and (f) 6,000 min.min sintered sample.

Fig. 3. (a) Bright-field TEM micrograph showing the boundaries betweenan abnormal grain and the matrix grains in a sample sintered at 550 �C inwet H2 for 60 min. (b) HR-TEM micrograph of the region indicated by acircle in (a). The diffraction pattern in (a) was obtained from the abnormalgrain with a [001] zone axis. Arrows indicate that the migrating boundaryof the abnormal grain has a (100) plane.

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Ni [2–5]. The similarity in AGG phenomena between ourultrafine Ni powder compacts and ED UFG Ni sheets

Fig. 4. SEM micrographs of (a) Ni powder and (c) Ni compact sintered adistribution. The particle size distribution was measured by a particle size analywas measured by an image analysis program (Matrox Inspector 2.1, Matrox E

suggests that the observed AGG is a general phenomenonoccurring in UFG Ni.

In a previous investigation [2], the LS-AGG in ED Niwas attributed to a possible enhancement of the boundarymobility with sulfur accumulation at the boundary and sul-fur phase wetting of the boundary with the growth of grainsfrom 20 nm to a few hundred nanometers in size. Thismechanism, however, raises a question of cause and effect.In order to examine the possibility of sulfur accumulationand wetting, the increase in matrix grain size during sinter-ing was measured and the impinged boundaries betweenabnormal grains were observed by TEM.

Fig. 4a and b show, respectively, an SEM micrograph ofthe Ni powder and its measured size distribution. Fig. 4cand d show, respectively, an SEM micrograph of the Nisample sintered for 60 min and its matrix grain size distribu-tion. The micrographs and the measured size distributionsshow that grain growth in the matrix is inconsiderable. Thissuggests that the appearance of abnormal grains is notrelated to the segregation of impurities and resultant forma-tion of a liquid phase.

The absence of S accumulation and wetting at theboundary was confirmed by energy-dispersive X-ray spec-troscopy (EDS) and HR-TEM observation. Fig. 5a and bare, respectively, a scanning transmission electron micros-copy (STEM) image of an impinged boundary of twoabnormal grains and EDS spectra across the boundary inthe sample sintered at 550 �C for 60 min. No noticeableincrease in S concentration at the boundary is noticedand the average intensity of S Ka peaks around the bound-ary is much lower than that in a previous study (e.g. Fig. 4in Ref. [2]). The HR-TEM micrograph in Fig. 5c also

t 550 �C in wet H2 for 60 min, and their (b) particle and (d) grain sizezer (LS 230, Beckman Coulter Co., USA). The 2-D grain size distributionlectronic Systems Ltd., Dorval, Canada).

Fig. 5. (a) STEM micrograph showing an impinged boundary of twoabnormal grains in a sample sintered at 550 �C in wet H2 for 60 min and(b) EDS spectra at the corresponding positions in (a). (c) HR-TEMmicrograph of the impinged boundary. The two diffraction patterns in (c)show that both of the grains are fcc Ni.

Fig. 6. (a) OIM image showing the microstructure of a Ni sample sinteredat 550 �C for 180 min in wet H2 and (b) its [001] inverse pole figure. Theregions having the same colors in the OIM image are the abnormal grainswith the same crystallographic orientations. The white regions are thematrix with ultrafine grains that are smaller than the beam step size of2 lm. The dashed (black) arrows and the solid (white) arrows on the OIMimage indicate the impinged facets and migrating facets, respectively. (Forinterpretation of the references to color in this figure legend, the reader isreferred to the web version of this article.)

Fig. 7. SEM micrograph showing the shape of an island grain in a samplesintered at 550 �C in wet H2 for 1,800 min. The SEM micrograph in theinset shows the removed shape of an island grain.

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shows the absence of any second phase at the boundary.These results demonstrate that AGG in UFG Niis not caused by the formation of a liquid phase at theboundary.

The initial texture mechanism was also proposed as themechanism of AGG in UFG Ni [3,7–11]. According to thismechanism, AGG in ED Ni is a result of the texture devel-opment from some unstable states to a stable state.Klement et al. [10] suggested that the ND//h111i texturedevelopment during IS-AGG in ED Ni was a result ofthe formation of grains having the lowest surface energy.The surface energy of fcc metals is known to increase inthe order {111}, {100}, {110}, {311} [27]. The grainswith a h111i direction parallel to the normal direction,however, do not necessarily have {111} boundary planes.Indeed, the grain boundary planes of the abnormal grainsin sintered ultrafine Ni are not {111} but {100}, as shownin Fig. 3. Tsurekawa et al. [5] also reported that the abnor-mal grains in ED UFG Ni had {10 0} facets. Therefore, itcan be concluded that ND//h111i texture development isnot related to the formation of planes with the lowest sur-face energy. In addition, our samples do not have anyspecific texture, as shown, for example, in the [001] inversepole figure in Fig. 6b of the OIM image in Fig. 6a.

3.2. Growth behavior with changing step free energy

Fig. 7 shows a typical square-shaped island grainentrapped within a grain in a sample sintered at 550 �Cin wet H2 for 1800 min. When the polished sample wasplaced in a high-energy ultrasonic bath for a few minutes,some of the island grains were removed as shown in theinset of the SEM image in Fig. 7. Theoretically, the shapeof an entrapped island grain represents the minimum grainboundary energy configuration (i.e. the equilibrium crystalshape), which is determined by the variation of the grainboundary energy with its normal, which is usually repre-sented by the polar plot of grain boundary energy (cg) vs.the grain boundary normal or the grain boundary Wulff

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plot cg(n). Note that the shape of island grains is cubic,which is identical to the shape of the abnormal grains. Thisresult suggests that the formation of abnormal grains isrelated to the migration of low-energy boundaries.

The migration of low-energy boundaries during AGGcan also be confirmed from the OIM images. The arrowson the OIM image in Fig. 6a indicate the normal directionsof the impinged facets between abnormal grains (in black)and the normal directions of migrating facets betweenabnormal grains and matrix grains (in white). Faceting isa phenomenon that lowers the energy of a grain boundaryas the boundaries dissociate into facets with different incli-nation angles. Thus, the normals of the impinged facetsrepresent the singular directions on the cg plot that isdrawn for the misorientations between the two adjacentabnormal grains. As shown in Fig. 3a, the normal ofmigrating facets (solid white arrows) are h100i. The nor-mal of impinged facets (dashed black arrows) of abnormalgrains also exhibit h100i, as can be seen in Fig. 6a and alsoin Fig. 5c. These results indicate that the migration of theboundaries of abnormal grains occurs while satisfying theirlow-energy configuration, i.e. conserving the {100} bound-ary facets.

Fig. 8. (a) Changes in grain growth behavior of Ni compacts at different temperin wet H2: (b) at 500 �C for 60 min; (c) at 500 �C for 180 min; (d) at 600 �C fo

The grain boundary energy anisotropy is characterizedby its step free energy which gradually decreases to 0 asthe temperature increases. A faceted grain boundary witha non-zero step free energy becomes rough (i.e. the step freeenergy decreases to 0) at its roughening transition temper-ature [46–48]. The roughening transition can also beinduced by changing P O2

[28,38,41] and incorporating addi-tives [35,37,41], as has been observed in many systems. Forthe Ni system too, faceted grain boundaries graduallybecome rough, i.e. the step free energy gradually decreaseswith increasing temperature [30] and decreasing oxygenpartial pressure [29,49].

To understand the effect of the boundary structure tran-sition on the grain growth in UFG Ni, the grain growthbehavior in UFG Ni was observed under different temper-atures and oxygen partial pressures.

Fig. 8a summarizes the change in grain growth behaviorobserved at different temperatures between 500 and 650 �Cin wet H2. Fig. 8b–e show some typical micrographsobtained after sintering under different experimentalconditions. At low temperatures, abnormal grains formedafter a short incubation period of time, �60 min at500 �C and �20 min at 550 �C, showing incubated AGG.

atures during sintering in wet H2. SEM micrographs of Ni samples sinteredr 0 min; and (e) at 600 �C for 20 min.

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At high temperatures, AGG occurred readily on reachingthe sintering temperature. At any temperature, graingrowth behavior apparently became stagnant afterimpingement of abnormal grains.

The number of cube-shaped abnormal grains in the unitarea and the growth rate of abnormal grains increased withincreasing sintering temperature. At 500 �C, only a fewabnormal grains formed and grew after 60 min with alow growth rate; in contrast, at 600 �C, many abnormalgrains formed and grew from 0 min with a high growth rateas shown in Fig. 8b–e. As a result, the average grain sizeafter impingement of abnormal grains was smaller at ahigher temperature.

Fig. 9a summarizes the grain growth behavior observedat 500–650 �C in wet 99N2/1H2. Two selected micrographsare also shown in Fig. 9b and c. The tendency for AGGwith increasing temperature is similar to that observed inwet H2, but AGG occurs much later than it does in wetH2, indicating a longer incubation time for AGG.

The migration of a faceted boundary, which has non-zero step free energy, is observed and is suggested to occurvia step growth or atom shuffling [50–55], similar to that ofa faceted single crystal in a liquid [45,56,57]. In the case oftwo-phase systems, the growth of a faceted grain is gov-erned by the diffusion of atoms through the matrix andthe reaction at the solid/liquid interface, respectively,above and below a critical driving force, showing mixedcontrol of growth. A recent investigation showed that themigration rate of a faceted boundary is nonlinear withrespect to the driving force, revealing the presence of acritical driving force for appreciable migration [43], as inthe case of two-phase systems. The previously suggested

Fig. 9. (a) Changes in grain growth behavior of Ni compacts at different tempesintered in wet-99N2/1H2: (b) at 550 �C for 600 min; and (c) at 600 �C for 600

mixed-control model of grain growth and the principle ofmicrostructural evolution for two-phase systems [42,44]can therefore be valid also for single-phase systems.According to the model and principle, the grain growthbehavior in a system with a critical driving force for bound-ary migration is a reflection of the coupling effect betweenthe maximum driving force (Dgmax) for growth in the sys-tem, which is for the largest grain, and the critical drivingforce for appreciable migration (Dgc) of faceted boundaries[42,44].

Based on the mixed-control model of grain growth, therelationship between the growth rate of grains and the driv-ing force for growth of grains under various sintering con-ditions can be depicted schematically, as shown in Fig. 10.The initial maximum driving force for grain growth mustbe the same for different samples because the initial particlesize and the distribution are the same. The critical drivingforce (Dgc), however, varies with changing temperatureand atmosphere because Dgc is dependent on the boundarystructure, i.e. the step free energy of the boundary [42,46–48]. For a given P O2

, as the temperature increases, theboundary becomes less faceted and the value of Dgc

decreases, as depicted in Fig. 10. Meanwhile, the absolutemigration rate above Dgc increases with increasing temper-ature because diffusion across the grain boundary increaseswith increasing temperature. At a fixed same temperature,the value of DgT ;N

c (DgTc in wet 99N2/1H2) must be higher

than that in wet H2 (DgT ;Hc ) because the faceting tendency

of the boundary increases as P O2increases [29,49].

For the samples sintered at 500 �C in wet H2, as abnor-mal grains appear after 60 min, Dgmax at the beginning(DgI

max) should be lower than Dg500;Hc , as schematically

ratures during sintering in wet 99N2/1H2. SEM micrographs of Ni samplesmin.

Fig. 10. Schematic representing the relationship between the grain growth rate ( _R) and the driving force for migration (Dg) for various conditions. Therelative values between DgI

max and DgT ;Hc (DgT ;N

c ) are highlighted in a separated box.

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depicted in Fig. 10. This means that all of the grains ini-tially have driving forces lower than Dg500;H

c . As the sinter-ing time increases, however, grain growth occurs, though isnegligible, apparently showing stagnant behavior. Accord-ing to our previous calculation [44], Dgmax can increaseduring extended annealing because the growth rate is non-linear with respect to the driving force [44]. Abnormalgrains can form after a certain period of time, as shownin Fig. 8a and b.

As the sintering temperature is increased to 550 �C inwet H2, Dgc (Dg550;H

c ) decreases towards Dgmax, as schemat-ically depicted in Fig. 10, and thus the incubation time forAGG is reduced to �20 min (Figs. 2b and 8a). Withincreasing temperature, the nucleation rate of abnormalgrains should increase because of the reduction of Dgc

and the enhancement of atomic mobility. It appears thatthe number of abnormal grains is larger in the sample sin-tered at 550 �C (Fig. 2b) than in the sample sintered at500 �C (Fig. 8c). The growth rate of abnormal grains alsoincreases with an increase in temperature from 500 to550 �C due to an increased rate of boundary migration.After impingement of abnormal grains, however, graingrowth may practically stop for a long time, as shown bythe data summarized in Fig. 8a, because of a considerablereduction of Dgmax, which can be much lower than Dgc.The micrographs in Fig. 2 demonstrate that the drivingforce for grain growth is reduced far below Dgc after anincrease in average grain size by �100 times.

A further increases in temperature to 600 and 650 �Cresults in a further reduction of Dgc. The absence of anincubation period for AGG indicates that Dgc is lower than

Dgmax at the sintering temperature. At the same time, thenumber of abnormal grains per unit area increases andthe time necessary for the impingement of abnormal grainsdecreases, as shown in Fig. 8a.

When the sintering atmosphere changes from wet H2 towet 99N2/1H2, the critical driving force increases, asschematically shown in Fig.10, as a result of an increased fac-eting tendency under a higher P O2

. Hence, we can expect toobserve a retardation of the appearance of abnormal grainscompared with that in the sample sintered in wet H2, asobserved in Fig. 9. The time period of the stagnation of graingrowth increases. Only at 650 �C do abnormal grains readilyappear upon reaching the sintering temperature.

The observed change in grain growth behavior withrespect to temperature, time and oxygen partial pressure,from SGG to incubated AGG and finally AGG, is wellexplained in terms of the coupling effect of Dgmax andDgc, which is based on the mixed-control model of graingrowth. The previously suggested principle of microstruc-tural evolution for two-phase systems [42,44] thereforeappears also to be valid for single-phase systems. TheAGG in UFG Ni is related to the intrinsic nonlinear migra-tion behavior of faceted boundaries.

4. Conclusions

When Ni compacts that have negligible S content and arefree of texture were sintered at 550 �C in wet H2, AGGoccurred, as in the case of AGG in ED UFG Ni sheets.The observed AGG, however, cannot be explained by theconventional AGG mechanisms of ED Ni sheets, i.e. the

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S phase wetting mechanism and initial texture mechanism.During AGG, cube-shaped abnormal grains with {100}facets formed, showing that the lowest energy boundariesmigrated with the formation of abnormal grains.

To investigate the relationship between the step freeenergy and grain growth behavior in UFG Ni, powdercompacts were sintered for various times at temperaturesbetween 500 and 650 �C under two different oxygen partialpressures. Depending on the temperature and oxygen par-tial pressure, first-stage SGG, AGG and second-stage SGGoccurred successively with sintering time. The observedgrain growth behavior can be well explained by themixed-control model of grain growth in faceted systemsin terms of the coupling effect of the maximum drivingforce for grain growth and the critical driving force forappreciable migration of the boundary. It is concluded thatAGG in UFG Ni occurs as a result of non-linear migrationof faceted boundaries with respect to the driving force.

Acknowledgements

This work was supported by a National Research Foun-dation of Korea (NRF) grant funded by the Korean gov-ernment (MEST) (No. 2011-0017556) and also bySamsung Electro-Mechanics Co. Ltd. through the Centerfor Advanced MLCC Manufacturing Processes.

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