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Scripta Materialia 54 (2006) 2137–2141

Grain growth in a bulk nanocrystalline Co alloy duringtensile plastic deformation

G.J. Fan a,*, L.F. Fu b, D.C. Qiao a, H. Choo a, P.K. Liaw a, N.D. Browning c

a Department of Materials Science and Engineering, The University of Tennessee, Knoxville, TN 37996, USAb Department of Chemical Engineering and Materials Science, University of California, Davis, CA 95616, USAc National Center for Electron Microscopy, Lawrence Berkeley National Laboratory, Berkeley, CA 94720, USA

Received 13 December 2005; received in revised form 19 February 2006; accepted 25 February 2006Available online 29 March 2006

Abstract

A bulk nanocrystalline Co–P alloy was subjected to tensile tests. Grain growth from approximately 12 nm in the as-deposited state toabout 25 nm after the tensile test was observed. Grain growth was not observed when the deforming volume was increased, or when thespecimen was annealed at 433 K or above prior to the tensile tests.� 2006 Acta Materialia Inc. Published by Elsevier Ltd. All rights reserved.

Keywords: Nanocrystalline materials; Grain growth; Mechanical properties

1. Introduction

It is well known that grain growth occurs during thethermal annealing of nanocrystalline (nc) metals and alloysat a relatively low annealing temperature, due to the excessfree energy stored in the nc grain boundaries [1–5]. Recentstudies indicate that grain coarsening may occur duringindentation or tensile tests of nc metals and alloys at roomtemperature or even at a cryogenic temperature [6–9]. Thegrain growth caused by the plastic deformation of ncmetals and alloys raises new concerns regarding the micro-structural instability of nc metals and alloys.

In contrast to the thermally-activated grain growth of ncmetals and alloys, where grain coarsening occurs via grainboundary migration driven by the thermal activation toreduce the free energy of the system, the grain growthobserved during plastic deformation may be purely drivenby mechanical stress, since the temperature rise associatedwith the plastic deformation is unlikely to lead to recrystal-lization and eventual grain growth. Molecular dynamicssimulations have suggested that grain coarsening during

1359-6462/$ - see front matter � 2006 Acta Materialia Inc. Published by Else

doi:10.1016/j.scriptamat.2006.02.041

* Corresponding author.E-mail address: gfan@utk.edu (G.J. Fan).

plastic deformation may be due to (a) the grain boundarymigration driven by mechanical energy [10], and (b) thecurvature-driven grain boundary migration and/or grainrotation-induced grain coalescence [11,12].

In this paper, we report that a bulk nc Co–P alloy expe-riences grain growth during the tensile plastic deformation.It was suggested that the grain growth was due to the grainboundary migration and/or grain coalescence involvinggrain boundary activities, which may be responsible forthe observed plasticity during tensile deformation of thebulk nc Co–P alloy.

2. Experimental procedure

The bulk nc Co–1.65 wt.%P alloy sheets (less than0.09 wt.% sulfur and 0.07 wt.% carbon) were producedusing a pulsed-electrodeposition technique by IntegranTechnologies Inc. [13]. The structure of the as-depositedsamples was determined using an X-ray diffractometer(XRD). The as-deposited alloy sheets had dimensionsof 70 mm · 10 mm · 2.5 mm. Dog-bone specimens werecut from the as-deposited sheets with various gaugecross-sections and gauge lengths, and were polished forthe uniaxial tensile tests. An extensometer was used during

vier Ltd. All rights reserved.

30 40 50 60 70 800

1.000

2.000

3.000

4.000

5.000

Inte

nsity

2 Theta Degree

(200)

(220)

(111)

Fig. 2. X-ray diffraction patterns of the as-deposited bulk nc Co–P alloy,showing fcc structure.

2500

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the tests for precise strain measurements. The fracturesurfaces were examined using scanning electron microscopy(SEM). Transmission electron microscopy (TEM) andhigh-resolution TEM (HRTEM) observations were carriedout before and after the tensile plastic deformation, using aSchottky field-emission gun FEI Tecnai F20 UT micro-scope operating at 200 kV with a spatial resolution of0.14 nm. The TEM and HRTEM thin-foil specimens wereprepared by the conventional twin-jet electro-polishingtechnique using a 23% perchloric acid and 77% acetic acidelectrolyte at 18 �C and 12 V. To study the effect of pre-annealing on the grain growth during the tensile deforma-tion, the samples were annealed for 20 min at differenttemperatures ranging from 433 K to 573 K. To examinethe variability of the grain sizes in the as-deposited sheet,three TEM thin-foils cut from different places were pre-pared for the TEM observations; this indicated that the dis-tributions of the grain sizes were fairly uniform across theas-deposited sheet.

3. Results and discussion

Fig. 1 shows the typical dark-field TEM image for theas-deposited nc Co–P alloy. HRTEM observations indi-cated that the as-deposited alloy contains a high densityof nanotwins. The observation of the growth nanotwinsduring deposition has been attributed to the low stack-ing-fault energy of the Co metal [13]. The presence of alarge density of growth nanotwins may influence the defor-mation behavior, as demonstrated by recent reports fornanotwinned metals [14,15]. By counting more than 200grains of the dark-field TEM images, it was found thatthe grain sizes ranged from about 2 to 22 nm for the as-deposited nc Co–P. The average grain size was about12 ± 2 nm, as determined from the grain size distributionstatistics. It is well known that Co exhibits a hexagonal-close-packed (hcp) structure at room temperature and aface-centered-cubic (fcc) structure at high temperatures.In contrary to the hcp structure reported for the as-depos-

Fig. 1. Dark-field TEM image and the corresponding diffraction patternsof the as-deposited bulk nc Co–P alloy with the average grain size of about12 ± 2 nm.

ited nc Co metal by Karimpoor et al. [13], the present bulknc Co–P alloy shows a metastable fcc structure at roomtemperature, as shown in Fig. 2 of the XRD patterns. Byusing the Scherrer formula [16], the grain size is calculatedto be 10 nm from the (20 0) diffraction peak of the XRDpatterns, which is in good agreement with the value calcu-lated from the TEM images.

Fig. 3 shows the engineering stress–strain curves of thenc Co–P alloy with different gauge sections and gaugelengths (Fig. 3a–c). For each gauge section and gaugelength, three tensile specimens were tested (Table 1). Forcomparison, the coarse-grained Co metal (hcp) wasincluded (Fig. 3d). With increasing the gauge dimensionsfrom 15 mm (length) · 1.5 mm (width) · 0.5 mm (thick-ness) (Fig. 3a for Sample I) to 30 mm · 3 mm · 0.8 mm(Fig. 3b for Sample II), and to 30 mm · 3 mm · 2 mm

ε. = 10-3 s-1

0 2 10 12 140

500

1000

1500

2000

d

Eng

inee

ring

Str

ess

(MPa

)

Engineering Strain (%)

c

b

a

4 6 8

Fig. 3. Engineering tensile stress–strain curves for the as-deposited bulknc Co–P alloy with different gauge dimensions (a–c), as well as for thecoarse-grained Co metal (d). Sample I (a) with smallest dimensions showsa highest fracture strength and tensile ductility among the nc samplestested. The fracture strength decreases with increasing the gauge dimen-sions for Sample II (b) and Sample III (c), respectively.

Table 1Mechanical properties (tensile elongation to failure, ef; and the fracturestrength, rf) of the bulk nc Co–P alloy with various gauge dimensions(length · width · thickness) measured from the tensile tests at _e ¼ 10�3 s�1

Gauge dimensions (mm) Test number ef (%) rf (MPa)

15 · 1.5 · 0.5 Sample I 1 2.4 19902 2.4 19403 2.2 1910

30 · 3 · 0.8 Sample II 1 0.9 15502 1.0 16003 1.1 1650

30 · 3 · 2 Sample III 1 0.4 10102 0.4 10403 0.5 1160

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(Fig. 3c for Sample III), the fracture strength decreasedcontinuously from 1,950 ± 80 MPa to 1600 ± 100 MPa,and to 1070 ± 150 MPa, respectively. All of the bulk ncsamples with different dimensions exhibited almost thesame Young’s modulus of 280 GPa. Sample I (with small-est dimensions) showed about 2.3 ± 0.2% tensile elonga-tion to failure, while Sample II and Sample III did notshow any global tensile ductility. Therefore, the mechanicalproperties of the present bulk nc Co–P alloy showed astrong size effect, as has often been observed in other ncmetals and alloys [17–19].

The decrease in the fracture strength and the loss of thetensile ductility with increasing deforming volume wereattributed to the presence of defects (i.e. porosity and/orsecond particle inclusions) in the nc metals and alloys[17–19]. The fracture surfaces after tensile tests were exam-ined using SEM, Fig. 4. A well-developed dimple structurewas observed for Sample I (Fig. 4a). Although Sample IIIwith the large deforming volume did not show any globaltensile ductility, the local plasticity within the fracture areawas still observed, as evidenced by the dimple structure inFig. 4b. It is interesting to note that the dimple sizes areseveral times larger than the original grain sizes, whichhas been explained by the theoretical modeling and com-puter simulations previously [20,21]. For example, Kumaret al. [20] have attributed the observed dimple structure

Fig. 4. The fracture surfaces of Sample I (a) and Sample III (b). The i

to the formation of voids, which act as the nucleation sitesfor the dimples and leads to the fracture of the nc materials.Therefore, the dimple sizes do not depend on the grain sizesince the fracture of nc materials is not along the grainboundaries. However, it is not clear why the dimple sizeswere even larger for Sample III, which did not experienceextensive plastic deformation. It might be due to the lowdensity of voids prior to the fracture resulting from thepremature failure of the tensile specimens.

The gauge sections of the tensile specimens after thefracture were further examined by TEM. Fig. 5 showsthe dark-field TEM images for the nc Co–P alloy afterthe deformation at a strain rate of 10�3 s�1. After the ten-sile plastic deformation, the grain sizes of the Sample I ran-ged from about 2 nm to 40 nm. The average grain size ofSample I increased from about 12 ± 2 nm in the as-depos-ited state to about 25 ± 4 nm after plastic deformation.The microstructure of the grip sections of the Sample I,which did not experience the plastic deformation, wasexamined by TEM. The results indicate that the grain sizesremain unchanged, suggesting that the observed graingrowth in the gauge section of the Sample I is due to theplastic deformation. After the grain growth, the grains stillcontain a high density of nanotwins. However, graingrowth was not observed in Sample II and Sample III(Fig. 5b), which exhibited lower fracture strengths withoutmacroscopic plastic yielding. The results indicate that asufficiently high stress is required to achieve the measurableplasticity perhaps via the grain boundary activities and theassociated grain coarsening, which is supported by thechanges in the fracture strength of the nc Co–P alloyssubjected to annealing at different temperatures prior totensile tests. Fig. 6 shows the fracture strength of SampleI annealed for 20 min at various temperatures prior to ten-sile tests. After annealing at 423 K, the fracture strengthdropped to 1850 ± 80 MPa from 1950 ± 80 MPa for theas-deposited sample. Further increasing the annealing tem-perature caused a rapid drop in the fracture strength. In theliterature, the observed annealing-induced brittleness hasbeen attributed to the impurity (particularly sulfur) segre-gations at the grain boundaries, which may hinder grainboundary mobility [22,23]. Current TEM observations

nsets in (a) and (b) are the fracture surface in low magnifications.

200 400 600 800500

1000

1500

2000

Frac

ture

Str

engt

h (M

Pa)

Annealing Temperature (K)

Fig. 6. The fracture strengths of Sample I annealed for 20 min at varioustemperatures prior to tensile tests. Below 573 K annealing itself did notcause any measurable grain growth.

Fig. 5. Dark-field TEM images and the corresponding diffraction patterns for Samples I (a) and III (b) after the tensile deformation at a strain rate of10�3 s�1. The grain size of Sample I was increased to about 25 ± 4 nm, while that of Sample III remained unchanged compared to the as-deposited sample(grain size of about 12 nm) shown in Fig. 1.

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(not shown here) indicate that the thermal stability of ncCo metal can be significantly enhanced by the alloying ofP element. Annealing the bulk nc Co–P alloy below573 K did not cause any significant grain growth, whichis consistent with recent findings by Choi et al. [24]. Thebrittleness of the annealed nc Co–P specimens may bedue to the reduced grain boundary mobility. In this case,the samples failed prematurely before the stress was highenough to cause the grain growth, and, thereby, the tensileductility.

It is well established that plastic deformation may refinethe grain size of coarse-grained metals and alloys, due tothe introduction of dislocation-cell structures and eventualgrain boundary formation during the plastic deformation.This approach has been widely employed in fabricatingbulk ultrafine-grained (ufg) and nanocrystalline (nc) metalsand alloys, for example, using equal channel angular press-ing [25–28]. The observed grain growth in the present bulknc Co–P alloy, together with those observed in other ncmetals and alloys [6–9], suggests that plastic deformation

of nc metals and alloys cause grain coarsening, instead ofgrain refinement in the coarse-grained metals and alloys.

Furthermore, the simultaneous grain coarsening duringthe tensile test is expected to cause a stress drop in the engi-neering stress–strain curve, which was observed in Fig. 3afor the Sample I immediately after the plastic yielding.However, the necking instability, which is commonlyobserved during the tensile test of ductile samples, couldalso contribute to the observed stress drop. The strengthsoftening due to the simultaneous grain coarsening hastwo competing consequences in terms of the plasticity ofnc materials: (a) strength softening will lead to plastic insta-bility, resulting in a reduced tensile ductility; and (b)strength softening at the same time will cause an increasedtensile ductility due to the trade-off between the strengthand ductility of materials. Compared to the report for anhcp-structured nc Co metal by Karimpoor et al. [13], thelow tensile ductility observed in the present bulk fcc-struc-tured nc Co–P alloy might be due to the early plastic insta-bility caused by the strength softening. The relationshipbetween the grain coarsening and the mechanical proper-ties of nc materials during the tensile plastic deformationhas also been discussed by Gianola et al. [29].

Recent studies have indicated that mechanical twinningis one of the important deformation mechanisms responsi-ble for the plastic deformation of nc metals and alloys [13–15,25,30–33]. However, the presence of a high density ofgrowth nanotwins before and after the plastic deformationmay not be directly related to the observed grain coarsen-ing. Based on the recent experimental observations of thegrain growth in nc Cu, Al, Ni metals [6–9] as well as themolecular dynamics simulations [10–12], it is suggestedthat the grain growth during the plastic deformation mightbe due to either one of, or a combination of, (a) the grainboundary migration and (b) the grain rotation-inducedgrain coalescence. However, the exact mechanisms bywhich grains grow during plastic deformation of nc metalsand alloys and the influence of growth nanotwins requirefurther experimental and theoretical studies.

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4. Conclusions

We reported the observations of the simultaneous graingrowth in a bulk fcc-structured nc Co–P alloy during ten-sile plastic deformation. Grain coarsening was caused whenmechanical stress was sufficiently high to initiate the grainboundary activities (e.g. grain boundary migration and/orgrain coalescence), which are responsible for the plasticdeformation of the nc Co–P alloy. The current results dem-onstrate that deformation-induced grain coarsening is animportant phenomenon associated with plastic deforma-tion of nc metals and alloys.

Acknowledgement

This work was supported by the National Science Foun-dation (NSF) International Materials Institutes (IMI)Program (DMR-0231320). The microscopy was performedin the National Center for Electron Microscopy at Law-rence Berkeley National Laboratory supported by the USDepartment of Energy under Contract No. DE-AC02-05CH11231 and by Grant No. DE-FG02-03ER46057.

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