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Materials Science and Engineering A 499 (2009) 427–433

Contents lists available at ScienceDirect

Materials Science and Engineering A

journa l homepage: www.e lsev ier .com/ locate /msea

ole of strain reversal in grain refinement by severe plastic deformation

mitry Orlova,∗, Yoshikazu Todakab, Minoru Umemotob, Nobuhiro Tsuji a

Department of Adaptive Machine Systems, Graduate School of Engineering, Osaka University, 2-1 Yamada-oka, Suita, Osaka 565-0871, JapanDepartment of Production Systems Engineering, Toyohashi University of Technology, Tempaku-cho, Toyohashi, Aichi 441-8580, Japan

r t i c l e i n f o

rticle history:eceived 13 June 2008eceived in revised form 26 August 2008ccepted 5 September 2008

a b s t r a c t

Most of severe plastic deformation processes involve strain reversal. Till now quite big number ofresearches has been done on indirect study of its role, which discusses the effect of loading path inultra-fine grained structure formation. This work is aimed to study directly the role of the strain rever-sal in grain refinement by severe plastic deformation. A 99.99% purity aluminum was processed by high

eywords:train reversalluminumPTBSDEM

pressure torsion up to 96◦ rotation (maximal equivalent strain ε ≈ 8) with two deformation modes: mono-tonic and reversal deformations with a step of 12◦ rotation (maximal equivalent strain ε ≈ 1). It was shownthat strain reversal retarded the grain refinement in comparison with the monotonic deformation. Thisexplains slower rate of the nanostructure formation of the SPD processes that involve strain reversal.

© 2008 Elsevier B.V. All rights reserved.

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rain subdivision

. Introduction

At the present time, severe plastic deformation (SPD) is a wellroved tool for fabricating ultrafine-, nano- [1–3] and even amor-hous [4,5] structures in metallic materials. Under SPD processing,desired structure is formed continuously from “top” coarse grains

down” to ultra-fine/nano-grains or amorphous. There are severalPD methods developed during last two decades. Most of themnvolve reversal straining essentially, e.g. repetitive corrugationnd straightening [6], cyclic extrusion and compression (CEC) [7],wist extrusion (TE) [8], and multidirectional forging (MF) [9,10],hile other techniques, e.g. accumulative roll bonding (ARB) [11,12],

pplies (quasi)monotonic straining. In some processes, deforma-ion path might be varied depending on need, e.g. high pressureorsion (HPT) [13] and equal channel angular extrusion (ECAE)14]. However, effects of loading scheme and deformation path ontructure development, especially in case of ECAE [15–21], are stilliscussed vigorously.

Effects of accumulated strain value, deformation temperature,train rate and hydrostatic pressure on microstructure evolution

ave been studied already rather well, while the effect of straineversal under large strain accumulation is studied still moderately.he role of strain reversal in flow stress drop (Bauschinger effect)nd structure evolution has been studied well enough for fatigue

∗ Corresponding author. Tel.: +81 668794173; fax: +81 668794174.E-mail address: orlov@ams.eng.osaka-u.ac.jp (D. Orlov).

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921-5093/$ – see front matter © 2008 Elsevier B.V. All rights reserved.oi:10.1016/j.msea.2008.09.036

est conditions. But in the fatigue tests, plastic strain introduced isp to 1%. For the case of severe plastic straining, the strain rever-al has been studied marginally for the case of ECAE processingy route C that involves rotation of a specimen by 180◦ around itsongitudinal axis before each consequent pass [16,19,20,22]. Directxperimental investigations of the effect of strain reversal at largetrains on structure and mechanical properties evolution were doneor pure Ni and Armco Fe [23], Al–Mg–Sc alloy [24], pure Ti andow carbon steel [25]. This effect was also discussed in theoretical

orks [26,27]. In these reports, it was found that the strain reversalas significant effect on the structure evolution, but it is signif-

cantly different depending on material, and hence other directxperimental investigations of this effect is of great interest.

Therefore, the present research was aimed to clarify the effectf strain reversal on grain refinement in pure aluminum in directxperiments. For this purpose, HPT technique was utilized. Amongther SPD techniques, only HPT allows to introduce precisely con-rolled values of simple shear strain as well as direction of straining,nd therefore is the most appropriate for the present research.

. Experimental

In this study aluminum of 99.99% purity was used. The as-

eceived cold rolled sheet of 1 mm thickness was cut into 10 mmn diameter samples. The samples were mechanically grinded to.6 mm thickness, annealed for 1 h at 500 ◦C and cooled in a fur-ace. As a result, initial material had fully recrystallized structureith an average grain size of 290 �m.

428 D. Orlov et al. / Materials Science and Engineering A 499 (2009) 427–433

FtC

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Table 1Equivalent strain values achieved in this study (estimated by Eq. (1)).

Torsion angle Position in the specimen

M

29

tA

f

ε

wta

sseegopiTFwadiE

Fo

ig. 1. Schemes of HPT processing routes used in this study. The diagram inserted inhe figure shows the principal difference between monotonic and reversal straining.W and CCW mean clockwise and counterclockwise rotations, respectively.

HPT processing was done continuously, without stops betweenorsion direction changes, according to the schemes shown in Fig. 1t room temperature and speed of ∼0.22 rotation per minute. Tor-ions up to 96◦ were used for the case of monotonic loading. Foreversal straining, samples were deformed to 12◦ rotation in clock-ise direction followed by counterclockwise rotation up to the

ame degree. So that one cycle of 12◦ clockwise followed by 12◦

ounterclockwise rotations is mathematically equal to 24◦ rotationn one direction. Four cycles of the reversal deformation is equalo 96◦ of monotonic counterpart. The processing was continuous

or the both deformation modes, and the processing time intervalsere equal. This increment in the step size (12◦ rotation) under thePT processing was chosen since within the specified tool geom-try, it was comparable with the strain for 1 pass of ECAE (in thease if 90◦ of channels intersection die is used) or 1 pass of TE (in

3

i

ig. 2. TEM micrographs of the 99.99%Al after HPT processing. The micrographs a, b and cf reversal straining with amplitude of ±12◦ . The micrographs a and d, b and e, c and f we

onotonic Reversal Axis Middle of radius Edge

4◦ ±12◦ × 1 0 1.01 2.026◦ ±12◦ × 4 0 4.03 8.06

he case if 60◦ of twist line slope angle die is used) or 2 passes ofRB.

Strain values under the HPT processing were estimated by theollowing equation [2]:

= 1√3

r

h� (1)

here ε is equivalent strain, r is a specimen radius, h is a specimenhickness, and � is a torsion angle in radians. Numerical values ofccumulated strains used in this study are shown in Table 1.

The structure characterization was carried out on the mid-ection perpendicular to axial direction of the as-processedpecimens by using transmission electron microscopy (TEM) andlectron backscattered diffraction (EBSD) techniques in a scanninglectron microscope equipped with a field-emission type electronun (FE-SEM). Thin foils for TEM were prepared by cutting samplesf 2 mm × 3 mm using a diamond cutter, grinding them on sandpa-er to thickness of ∼0.1 mm and finally twin-jet electro-polishing

n a 30%HNO3 + 70%CH3OH solution at −20 ◦C to perforation. TheEM observations were done in Hitachi H-800 operated at 200 kV.or EBSD analysis, specimens were grinded on sand paper, polishedith alumina powder suspensions and finally electro-polished in30%HNO3 + 70%CH3OH solution at −20 ◦C. The observations wereone in FEI Sirion FE-SEM facility equipped with TSL orientation

maging microscopy system. For all statistical analysis from theBSD data, at least 1500 (sub)grains were considered.

. Results

Micrographs resulting from the TEM observations are shownn Figs. 2 and 3. Development of dislocation substructures and

correspond to monotonic straining to 24◦ rotation; d, e and f correspond to 1 cyclere obtained at specimens’ axis, middle radius and edge, respectively.

D. Orlov et al. / Materials Science and Engineering A 499 (2009) 427–433 429

F and co d f we

gtttevsm

ts

FcB

ig. 3. TEM micrographs of the 99.99%Al after HPT processing. The micrographs a, bf reversal straining with amplitude of ±12◦ . The micrographs a and d, b and e, c an

rain refinement through grain subdivision are clearly seen fromhese figures. Detailed analysis of mechanisms of the microstruc-ure evolution in pure aluminum under SPD is rather well done in

he previous references [28–32]. Mechanisms of the microstructurevolution observed in this study are quite consistent with the pre-ious knowledge. However, there are some particular differences intructural evolution between monotonic and reversal deformationodes.

ctsct

ig. 4. Boundary maps obtained from EBSD analysis of the 99.99%Al after HPT processingorrespond to 1 cycle of reversal straining with amplitude of ±12◦ . The maps a and d, b anoundaries with angles of misorientation 2◦ ≤ � < 15◦ are shown in grey lines, while boun

correspond to monotonic straining to 96◦ rotation; d, e and f correspond to 4 cyclesre obtained at specimens’ axis, middle radius and edge, respectively.

Comparing the micrographs in Fig. 2, one can see that after theorsion equivalent to 24◦ rotation for both monotonic and rever-al modes, the structures are inhomogeneous in the specimens’

ross-sections. In the areas close to the specimens’ axis, disloca-ion substructures typical for early stages of grain refinement inevere plastic deformation are seen. The microstructures also indi-ate that plastic deformation certainly occurred to some extents,hough the strain value on the specimen axis is ideally zero in HPT,

. The maps a, b and c correspond to monotonic straining to 24◦ rotation; d, e and fd e, c and f were obtained at specimens’ axis, middle radius and edge, respectively.daries with misorientations larger than 15◦ are in black.

430 D. Orlov et al. / Materials Science and Engineering A 499 (2009) 427–433

F essingc d, b anB boun

atosdbrTtam

tr

(tm

Fdsw

ig. 5. Boundary maps obtained from EBSD analysis of the 99.99%Al after HPT procorrespond to 4 cycles of reversal straining with amplitude of ±12◦ . The maps a andoundaries with angles of misorientation 2◦ ≤ � < 15◦ are shown in grey lines, while

s was shown in Table 1. At the specimens’ edge (ε ≈ 2, Table 1) inhe case of monotonic straining, most of subgrains are almost freef dislocations, and the (sub)grain boundaries are quite thin andtraight. In case of the reversal straining, there are still arrays ofislocations within the subgrains and in vicinity of the (sub)grainoundaries even at the specimen edge. In the areas of the middle

adius, transition between the structures on axis and edge is seen.his confirms the gradual evolution of the ultrafine grained struc-ures during severe plastic deformation; i.e., dislocation activation,ccumulation, rearrangement and finally (sub)grain boundary for-ation leading to the ultrafine grains. It was also found that in

csrbt

ig. 6. Microstructural parameters obtained from the EBSD data showing the structure eevelopment at edge part of the specimens (a) and fraction of grains surrounded by HAtrain-induced cells surrounded by boundaries with misorientations ≥2◦ , the minimum anas calculated from at least 1500 (sub)grains.

. The maps a, b and c correspond to monotonic straining to 96◦ rotation; d, e and fd e, c and f were obtained at specimens’ axis, middle radius and edge, respectively.

daries with misorientations larger than 15◦ are in black.

he case of reversal deformation this microstructural evolution isetarded in comparison with monotonic one.

Fig. 3 shows that after the torsion equivalent to 96◦ rotationmaximum strain ε ≈ 8, Table 1), the structure inhomogeneity inhe specimens’ cross-sections is still kept for both modes of defor-

ation. On the specimen axis, almost no change in the structure

ould be seen, compared with Fig. 2(d), in the case of the rever-al deformation; while in the case of the monotonic deformationeduction in dislocation density and formation of limited num-er of subgrain boundaries are apparent. At the specimens’ edgehe (sub)grains are free of dislocations, reasonably equiaxed and

volution with strain accumulation in the 99.99%Al HPT processed. Subgrain sizesBs (the boundaries with misorientations ≥15◦) in total number of subgrains (thegular misorientation resolution detectable by the EBSD technique) (b). The statistics

D. Orlov et al. / Materials Science and Engineering A 499 (2009) 427–433 431

F T proci ic stra

smttstta

nbgdbdsdtridm

scma2agnffirt

sedtdoetast

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ssSfd

ig. 7. Misorientation angle profiles obtained from EBSD analysis of the 99.99%Al HPnserted in the diagrams schematically. The maps a, b and c correspond to monoton

urrounded by sharp boundaries in the both cases of deformationodes. The middle radius areas show significant difference in struc-

ure: well defined (sub)grains but still containing dislocations inhe case of reversal deformation, while clear grains surrounded byharp boundaries and free from dislocations inside in the mono-onic specimen. These observations suggest rather sharp change inhe microstructure depending on the distance from the specimens’xis, i.e., strain.

According to the EBSD observations, initial structure was sig-ificantly refined by the HPT processing, as can be seen from theoundary maps shown in Figs. 4 and 5. The structure inhomo-eneity depending on the distance from the specimens centersiscovered in TEM observations is well confirmed in larger viewsy the EBSD analysis. The inhomogeneity is less in monotoniceformation and decreases with increasing total accumulatedtrain value (rotation degree/number of cycles). Depending on theistance from the specimens’ axis, continuous evolution of struc-ure from subgrains with low angle boundaries (LABs), throughather heterogeneous mixtures of high angle boundaries (HABs)ntersected by LABs, to ultrafine-grained equiaxed structure withominating HABs at the specimens edge could be seen for bothonotonic and reversal specimens.The average (sub)grain sizes obtained from the EBSD data are

ummarized in Fig. 6. For all the data shown there, statistics wasalculated from at least 1500 (sub)grains. For the both deformationodes, very fine subgrain size of about 1.0–1.2 �m was achieved

t the specimens’ edge already after the torsion equivalent to4◦ rotation (equivalent strain ε ≈ 2, Table 1). Further processingffects the average subgrain size very slightly, but number of therains surrounded by HABs increases. This effect is very much pro-

ounced in the case of monotonic deformation, but much slower

or the reversal counterpart. The boundary misorientation pro-les shown in Fig. 7 are in good agreement with the qualitativeesults from the TEM observations and also suggest the retarda-ion in the structure evolution in reversal deformation. On the

esags

essed. Positions where the EBSD maps were obtained are shown in the illustrationsining; d, e and f correspond to reversal straining.

pecimens’ axis (Fig. 7a and d), there is a distribution of misori-ntation angles with a strong peak in LABs region for the botheformation modes. Almost no change in the misorientation dis-ribution is seen with increasing strain in the case of reversaleformation, while monotonic deformation leads to the increasef HAB fraction. In the areas of the specimens’ middle radius anddge for the 96◦ rotated samples, distributions of the misorien-ation angles are bimodal with small peak around 45◦ as wells the peak at very low angle. The reversal deformation to theame equivalent strain does not lead to the second peak forma-ion.

Subgrain boundaries having very low misorientation smallerhan 2◦ and therefore out of the EBSD resolution, do not affect prin-ipally the results shown above. Taking them into account wouldlightly decrease the average subgrain sizes and average angle ofisorientations, but would not affect qualitatively the conclusions

oncerning the effect of strain reversal on structure developmentnder large plastic straining. Slight discrepancy in the EBSD andEM observations for the middle radius area for the deformationsquivalent to 96◦ rotation, might be attributed to local effect owingo the structure inhomogeneity.

. Discussion

The structure inhomogeneity along specimens’ radius at earlytages of the processing reported in the previous section corre-ponds to the strain gradient introduced during the processing.uch a structure inhomogeneity is not surprising but a well knowneature in the HPT processing under monotonic straining con-itions, and usually the structure homogeneity is achieved by

xcessive torsion. At the same time, the absence of ultrafine grainedtructure evolution in vicinity of the specimens’ axis (where strainmplitude is smallest) and the increase in the structure inhomo-eneity with increasing strain are specific features of the reversaltraining by HPT.

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32 D. Orlov et al. / Materials Science a

Such behavior of the structure evolution is thought to be a con-equence of a kind of the Bauschinger effect. It is interesting to findhat the Bauschinger effect works not only in near elastic deforma-ion, but also even in the far plastic region of deformation. And dueo this effect, formation of HABs in the case of reversal straining isignificantly retarded in comparison with monotonic counterpart.

Based on the conclusions from the article by Hasegawa et al.33], in addition to the reversed motion of free isolated disloca-ions within cells or subgrains, cell walls and sub-boundaries arelso unstable against strain reversal. For low strain values, theecrease in dislocation density about 16% at early stages of reversaleformation was reported. These results are very well consistentualitatively with the results of the present work.

When present results are compared with earlier reports, itould be seen that behavior of pure Al is quite different from purei [25]: while strain reversal retards formation of HABs in alu-inum, in titanium it results in faster formation of equiaxed grains

urrounded by HABs. This difference should be attributed to theifference in lattices of these materials (face-centered cubic lattice

n Al, and hexagonal close-packed in Ti). At the same time, materi-ls with similar cubic lattices (Ni and Fe) studied in [23], show veryimilar behavior to our results: fractions of HABs in the deformedtructures were highest in the monotonic deformation mode, andower after the reversal deformation. In [23], minimal grain sizeseached after high strains were also affected by the deformationode: after reversal deformation average grain sizes were larger

han after the monotonic counterpart. In our research, at the spec-mens’ edge, where the strain amplitude was the largest, after theeversal deformation, the average grain size was ∼0.1 �m largerhan that after the monotonic one. This difference is much less thanhe deviations in the grain size distributions (Fig. 6) and could beonsidered as negligibly small. However, the reversal deformationith very low strain amplitude in vicinity of the specimens’ axis didot lead to ultra-fine grains formation at all. So that our results onure Al are very consistent with results of other researches on purei and Armco Fe. Nevertheless, more detailed, study of the effect ofmplitude of strain reversal should be done to further clarify thisssue.

The present results are also very consistent qualitatively withhe theoretical predictions [26,27] suggesting that “effective” strains lesser in the SPD processes which involve strain reversal than inquasi)monotonic one. And therefore, rate of saturation in HABsraction and mechanical properties are retarded in them [27].owever, SPD processes that involve strain reversal are predicted

o produce UFG materials with higher ductility [26], and furtherxperimental work should be done to check this prediction.

The present investigation confirms that the strain reversal,hich takes place in most of the SPD methods, plays a significant

ole in structure evolution during SPD. Specifically, it decreases theate of HABs formation at the stages of large strain accumulation.his explains slower rate of the nanostructure formation of the SPDethods that involve full strain reversal after each processing step,

.g. ECAE by route C, MF, TE, CEC, etc. in comparison with HPT, ECAEy routes A and B, and ARB. The present results suggest also thathe further, more detailed, studies on the effect of strain reversal onhe microstructure evolution in SPD and the mechanical propertiesf the obtained materials should be done for deeper understandingf the nanostructured metals. Numerical analysis of this effect islso of great interest.

. Conclusions

1. In this work, 99.99% purity aluminum was processed by highpressure torsion to large strain values with two deformation

[[

[[

gineering A 499 (2009) 427–433

modes: monotonic and reversal. The effect of strain reversal onintensity of grain refinement under severe plastic deformationwas studied.

. The strain reversal did not affect significantly grain refinementand final grain size, but significantly retarded the formation ofhigh angle boundaries in comparison with monotonic straining.The results explain slower rate of the nanostructure forma-tion of the SPD processes that involve full strain reversalafter each consequent processing step in comparison with the(quasi)monotonic techniques.

cknowledgements

Authors of this work gratefully acknowledge financial supportsy the Grant-in-Aid for scientific research from the ministry ofducation on priority areas “Giant straining process for advancedaterials containing ultra-high density lattice defects” through theinistry of Education, Culture, Sports, Science and Technology of

apan, and by Industrial Technology Research Grant Program’05hrough NEDO of Japan (project ID 05A27502d).

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