op

.W

g K

foe 2

tion [13,15], which refers to the easy zipping of disloca-tions through the small crystal without accumulation andmultiplication, therefore resulting in the crystal staying in

in the bulk state cannot be met.Ng and Ngan [22] have demonstrated that by coating the

free surface or lling an internal cavity of aluminum micro-pillars with tungsten deposition, dislocations are trappedinside the aluminum and their greatly enhanced multiplica-tion and interactions result in smooth deformation with

Corresponding author.E-mail address: hwngan@hku.hk (A.H.W. Ngan).

Available online at www.sciencedirect.com

Acta Materialia 60 (2012) 6102611. Introduction

The last decade or so has seen tremendous interest in theyielding of micropillar forms of metals. In addition to theirpower-law, size-dependent strength [111], these micropil-lars also deform in a jerky manner with continuous occur-rence of strain bursts in a stochastic manner [1,2,1217]. Toexplain these unusual behaviors, several theories have beenproposed. One important concept is dislocation starva-

a continuously dislocation-starved state. Such a mode ofdeformation has been proven by direct transmission elec-tron microscopy (TEM) observations in nanometric metalvolumes [15,17]. Another group of proposed modelsincludes source truncation [18,19] and exhaustion hard-ening [20,21], which describe the jerky and stochastic nat-ure of deformation as the result of infrequent interactionevents between the small content of dislocations presentin the microcrystal, so that the mean-eld conditions asAbstract

The plastic deformation of micropillars is known to be aected by whether dislocations can escape easily from the material volume,and the extent to which the dislocations mutually interact during the deformation. In this work, pre-straining and coating are used tomodify the initial dislocation content and the constraints on the escape of dislocations. Aluminum micropillars in the size range from 1to 6 lm, with or without thin coating by tungsten deposition and pre-straining by 7%, were compressed using a at-punch nanoind-enter to study their plasticity behavior. The results reveal very dierent behavior between the size regime of a few microns and that of1 lm. For pillars a few microns large, coating leads to signicant strengthening, and pre-straining by 7% also produces a mild strength-ening eect. The proof strength also exhibits good correlation with the square root of the residual dislocation density measured by trans-mission electron microscopy after deformation, indicating that strength in this size regime is controlled by dislocation interactions as intraditional Taylor hardening. Coating evidently helps retain dislocations inside the pillar, and pre-straining increases the initial disloca-tion content; both eects lead to more severe strain hardening during deformation. For smaller pillars 1 lm in size, however, pre-strain-ing results in softening, although coating still leads to strengthening, and the strength exhibits no correlation with the residual dislocationdensity, which remains close to the initial value even with coating. These suggest dislocation starvation in this small size regime and thatstrength is controlled by the availability of mobile dislocations. Coating cannot eectively trap dislocations inside the pillar, but can stillstrengthen the pillar, presumably because dislocation nucleation is more dicult at the coated surface. 2012 Acta Materialia Inc. Published by Elsevier Ltd. All rights reserved.

Keywords: Plastic deformation; Compression tests; Dislocations; NanoindentationEects of pre-straining and coatingmicro

R. Gu, A.H

Department of Mechanical Engineering, The University of Hon

Received 5 June 2012; received in revisedAvailable onlin1359-6454/$36.00 2012 Acta Materialia Inc. Published by Elsevier Ltd. Allhttp://dx.doi.org/10.1016/j.actamat.2012.07.048n plastic deformation of aluminumillars

. Ngan

ong, Pokfulam Road, Hong Kong, Peoples Republic of China

rm 18 July 2012; accepted 19 July 20123 August 2012

www.elsevier.com/locate/actamat

11rights reserved.

atultra-high proof strength and strain-hardening rate. Theseresults are consistent with an earlier short report [23] oncompressive behavior of Al2O3-coated gold micropillars.Recently, Jennings et al. [24] also observed that coppernanopillars with TiO2/Al2O3 coatings not only exhibitedmuch higher strength and the Bauschinger eect as in metal-lic thin lms with a passivation layer [2527], but also signif-icant stochastic nature in the compression behavior. Theseeects of trapping dislocations have also been veried bydislocation dynamics simulations [28,29]. On the otherhand, attempts have also been made to understand theeects of altering the initial dislocation contents in micropil-lars by pre-straining the metal prior to pillar machining,which was found to produce a softening eect in general[3032]. The evidence so far has therefore pointed to veryremarkable eects on the plastic deformation ofmicropillarsif their dislocation contents are conned as in the coatingexperiments, or increased a priori as in the pre-strainingexperiments. In this work, we investigate the coupled eectsof both means for conditioning the dislocation contents inaluminum micropillars over a range of sizes in the micronregime.

2. Experimental details

A piece of ultrapure aluminum (purity >99.99999 wt.%)was cut from a rod of 12 mm diameter. This sample wasannealed at 500 C for 24 h, and then treated by mechani-cal polishing and electropolishing in a 1:9 mixture of per-chloric acid and methanol at 30 C and 16 V for3.5 min. A large grain with orientation [516] and diameter>2 mm was detected by electron back-scattered diraction(EBSD) and was selected for the subsequent experimentsreported below. The bulk yield strength of the grain wasestimated as a third of its Vickers hardness as assessed bymicroindentation just on that grain.

Micropillars were fabricated by focused ion-beam (FIB)milling in a Quanta 200 3D dual beam FIB/SEM systemoperating at 30 kV ion beam voltage, through a series ofconcentric annular pattern millings with the current variedfrom 20 nA for initial coarse milling to 50 pA for nal nemilling. In this study, three micropillar diameters of 5.6, 3.3and 1.2 lm, all with height-to-diameter ratios of around4:1, were fabricated. Their appearances before and aftercompression were imaged by scanning electron microscopy(SEM) in a LEO1530 microscope. A batch of the fabri-cated pillars was then coated by tungsten deposition inthe same FIB system on the side surface. In this procedure,hexacarbonyl tungsten gas W(CO)6 [33] is injected near thespecimen, and this is then decomposed by the gallium ionbeam into a tungsten-rich solid deposit on the specimen.This deposit is not pure metallic tungsten but is alloyedwith a signicant amount of Ga and is in a nanocrystallineclustered state [33]. For the aluminum specimens in thiswork, the ion current used for tungsten coating was

R. Gu, A.H.W. Ngan /Acta M

at6104 R. Gu, A.H.W. Ngan /Acta Machieved with tungsten coating of the pillars, and the eectsare not sensitive to the volume fraction of the tungstencoating. However, for all the three specimen sizes, sudden

Fig. 1. Normal stressnormal strain curves of 7% pre-strained aluminum pillarversion only) are for coated pillars with dierent volume fractions Vw of theRepresentative SEM images of the deformed coated pillars with diameter of (b)references to color in this gure legend, the reader is referred to the web versierialia 60 (2012) 61026111softening manifested as a plateau in the stressstrain curvealways occurs at a later stage of deformation, and similarevents also happen without pre-straining [22]. For the

s of diameter (a) 5.6 lm, (c) 3.3 lm and (e) 1.2 lm. Colored curves (onlinetungsten coating; black curves are for uncoated pillars of similar sizes.5.6 lm, (d) 3.3 lm and (f) 1.2 lm are also shown. (For interpretation of theon of this article.)

montages of the annealed and 7% pre-strained states ofthe aluminum bulk from which the micropillars were fabri-cated. These TEM specimens were prepared by wire-cut-ting thin slices from the bulk piece and electropolishingusing the traditional twin-jet method. It can be seen inFig. 3a that in the annealed bulk dislocation lines distributerather separately without extensive clustering, whereas inthe 7% pre-strained state as in Fig. 3b the dislocation linesare much denser and assemble into subcells. The disloca-tion density in the two states was measured by a line-inter-cept method, i.e. the average distance d betweendislocations was measured from the number of intersectionpoints random lines drawn on the TEM image would makewith the dislocations, and the dislocation density q wasestimated as 1/d2. The dislocation density q for theannealed state is estimated as 9.0 1011 m2, and that

at5.6 and 3.3 lm coated pillars, Fig. 1a,c indicate that thedeformation is smooth with continuous strain-hardeningup to 5% strain, but at larger strains the strain-hardeningceases and a plateau follows. For the smaller 1.2 lm pil-lars in Fig. 1e, the softening at 5% strain is so rapid that itcorresponds to a huge and uncontrollable strain burst of10%. The post-deformation SEM images of the coatedpillars shown in Fig. 1b, d and f reveal the occurrence ofonly one or two intensive slip bands, around which thetungsten coating shows signs of cracking and aking othe pillar surface. As for the 1.2 lm pillars, which deformedwith a large avalanche, the post-deformation SEM mor-phology in Fig. 1f shows an intensive shear step corre-sponding to almost complete shear fracture of the pillarinto two segments. The large strain burst is thought to berelated to the cracking of the coating at a large stress[22,24], and as this happens substantial slip of the pillarcore would proceed at the cracked region.

3.2. Strengthening of tungsten-coated specimens

The stress plotted in Fig. 1a, c and e is the nominal stressr, which is the applied load divided by the gross cross-sec-tional area of the coated pillars including the coating thick-ness. To obtain the stress rAl applied onto the aluminumcore of the coated pillars, a composite model is used asin Ref. [22]:

r rAlV Al rwV w;where VAl and Vw are the aluminum and tungsten deposi-tion volume fractions, respectively, and rw is the stress sus-tained by the tungsten coating. The intrinsic strength of thetungsten coating deposition fabricated by the present FIBsystem was investigated previously [22], and its 2% proofstrength was found to be 100 MPa. Hence the 2% proofstrength rAl of the aluminum core can be calculated fromthe overall nominal stress r using the above rule of mix-tures. Such a rule-of-mixtures calculation assumes simpleload sharing between the constituent phases without mu-tual interactions, and the objective here is indeed to seewhether the pillar cores intrinsic strength can be explainedby such an eect, since if this is the case, the calculated rAlwould agree with the strengths of the uncoated pillars. Theresults for dierent cases, including the uncoated pillarswith no pre-straining (i.e. the pristine group), are shownin Fig. 2. For the coated pillars in both pre-straining con-ditions, their rAl data calculated from the rule of mixturesare always higher than the strengths of the uncoated pillarsof the same size in accordance with previous ndings [22],implying that the strength of the pillar core is aected bysome interaction eects from the coating. However, thecoated and pre-strained pillars (the red dotted line inFig. 2) are substantially weaker than the coated pillars withno pre-straining (the black dotted line) at sizes of 1.2 lm:the 2% proof strength of the coated/pre-strained group at

R. Gu, A.H.W. Ngan /Acta M1.2 lm size is 172.6 24.1 MPa which is only about54% of the coated group without pre-straining. At sizesof 3.3 lm, the strength of the coated/pre-strained groupis 133.2 8.8 MPa which is 78% of that of the coatedgroup without pre-straining, and at 5.6 lm, the strengthof the coated/pre-strained group is 101.4 16.1 MPawhich is 1.1 times that of the coated pillars without pre-straining. Pre-straining and coating evidently result in dif-ferent eects on the strength data in Fig. 2 between large(>3.3 lm) and small (1 lm) pillar sizes. For the uncoatedspecimens (the two solid curves in Fig. 2), their strengthvaries with diameter almost following a power law, andthe pre-strained specimens exhibit a smaller power expo-nent than the pristine case; this phenomenon will be exam-ined in greater detail, separately [36].

3.3. Dislocation distribution in the specimens before

compression

TEM characterization was carried out on representativedeformed specimens of each group in order to understandthe dislocation distributions in the micropillars. First it isnecessary to study the dislocation distribution in the initialstates before pillar fabrication. Fig. 3a and b are TEM

Fig. 2. Average 2% proof strength of aluminum micropillars with andwithout coating and pre-straining.

erialia 60 (2012) 61026111 6105for the 7% pre-strained state is 2.2 times higher at2.0 1012 m2. Vickers hardness (HV) measurement of

inum

atFig. 3. (a and b) TEM montages of annealed and 7% cold reduction alum

6106 R. Gu, A.H.W. Ngan /Acta Mthe annealed and pre-strained states was performed andthis gives yield stress estimates of the two states as 4.5and 5.8 MPa, respectively. From the Taylor hardeningequation r / qp , the dislocation density in the pre-strained state should be about 1.7 times that of theannealed state, and this ratio is close to the 2.2 from directmeasurement of the dislocation densities.

The assessment of the dislocation content in thedeformed pillars, to be reported below, was carried outon FIB-thinned longitudinal sections of the pillars asdescribed in Section 2. Therefore, irrespective of the dis-location structures as seen from twin-jet electropolishedTEM samples, the eects of ion bombardment duringthe FIB process need to be assessed. Fig. 3c shows aTEM image of a FIB-produced section of an undeformed

prepared specimen from the annealed bulk taken near the [101] pole.

Fig. 4. (a and c) Montages of TEM images of the longitudinal sections of 5.6pillar is towards the right side (the same orientation is used for all subsequcondition.samples [36] prepared by the twin-jet method. (c) TEM image of a FIB-

erialia 60 (2012) 61026111pillar without pre-straining, and the dislocation density inthis specimen is estimated to be 1.0 1013 m2, i.e. anorder of magnitude higher than that seen in the twin-jetprepared sample (cf. Fig. 3a). It should be noted thatthe FIB-fabricated micropillars were much thicker thanthe TEM foils, and so the background dislocation densityof 1013 m2 as seen from Fig. 3c may not be representa-tive of the interior of the micropillars. Furthermore, anydislocation density orders of magnitude higher than1013 m2 found in a compressed pillar should be inter-preted as arising from compression process itself, ratherthan due to the FIB damage during the TEM samplepreparation. These two points should be borne in mindwhen interpreting the dislocation density measurementsreported below.

lm uncoated pillars without (a) and with (c) 7% pre-strain. The top of theent TEM images). (b and d) Higher-magnication TEM images of each

atR. Gu, A.H.W. Ngan /Acta M3.4. Dislocation distribution in the deformed specimens

Figs. 47 show TEM microstructures of representativepillars in each condition after compression, and details aregiven Table 1. The dislocation densities are estimated as

Fig. 5. (a and c) Montages of TEM images of the longitudinal sections of Higher-magnication TEM images of each condition.

Fig. 6. (a and c) Montages of TEM images of the longitudinal sections of 5.6Higher-magnication TEM images of each condition.erialia 60 (2012) 61026111 6107the average values from several high-magnication TEMimages taken with diraction vectors g under which visibledislocations were most abundant. Although the deforma-tion strains of the specimens are dierent, experimentalobservations have showed that strain is not an important

1.2 lm uncoated pillars without (a) and with (c) 7% pre-strain. (b and d)

lm tungsten coated pillars without (a) and with (c) 7% pre-strain. (b and d)

at6108 R. Gu, A.H.W. Ngan /Acta Mfactor inuencing the dislocation densities in deformedmicrocrystals [20,37]. Fig. 4 compares the dislocation distri-bution with and without the 7% pre-straining in uncoatedpillars of 5.6 lm diameter. No obvious subcell structurecould be seen in the case without pre-straining (Fig. 4a),and the dislocation density is estimated to be 1.0 1014 m2,i.e. an order of magnitude higher than the 1013 m2

background introduced by the FIB process itself. As forthe pre-strained case, dislocation subcells several micronslarge are discernible in the montage shown in Fig. 4c, andit may be due to the fact that this pillar was fabricated fromthe cold-rolled bulk which already contained dislocationcellular structures as shown in Fig. 3b. From the higher-magnication image shown in Fig. 4d, it is clear that dislo-cations are more abundant than the case with no pre-strain(Fig. 4b), and their density is about 3.0 1014 m2.

Fig. 7. (a and c) Montages of TEM images of the longitudinal sections of 1.2Higher-magnication TEM images of each condition.

Table 1Details of deformed aluminum micropillars for TEM characterization.

Diameter(lm)

Pre-strain Tungsten volumefraction

Max. stress(MPa)

5.6 None None 585.6 7% None 651.2 None None 2401.2 7% None 2095.6 None 12% 1725.7 7% 20% 2271.1 None 10% 3331.2 7% 20% 295erialia 60 (2012) 61026111Fig. 5 compares the dislocation arrangements in thesmaller size of 1.2 lm pillars. For both cases with andwithout pre-straining, the dislocations arrange randomlywith density 2.8 1013 m2 for the case without pre-strain,and 5.4 1013 m2 with pre-strain, i.e. both values are sig-nicantly smaller than those of the larger pillars under sim-ilar conditions, and are close to the 1013 m2 background inthe initial states. Pre-straining evidently results in no obvi-ous eect on the dislocation distribution in the small pillars.

Figs. 6 and 7 show the dislocation structures in thecoated micropillars of large and small sizes, respectively.Fig. 6 shows that after deformation, dislocation cellularstructures were developed in the coated 5.6 lm pillarsboth with or without pre-strain. The higher-magnicationTEM images in Fig. 6b and d show that dislocations inthe subgrains were abundant at density 7.0 1014 m2

lm tungsten coated pillars without (a) and with (c) 7% pre-strain. (b and d)

Final strain g for dislocation densityestimation

Correspondinggures

0.05 111; 020 Fig. 4a and b0.11 200; 111 Fig. 4c and d0.07 111 Fig. 5a and b0.10 200; 111 Fig. 5c and d0.06 111; 111 Fig. 6a and b0.19 111; 111 Fig. 6c and d0.04 111 Fig. 7a and b0.04 111 Fig. 7c and d

result is that, even without coating, the residual dislocation

atfor the case without pre-strain, or 1.0 1015 m2 with pre-strain. On the contrary, for the 1.2 lm coated pillars, nodislocation cell could be observed in the TEM montagesshown in Fig. 7a and c with or without pre-strain. In thesesamples, dislocations are scarce and arrange in a randommanner, and their densities are estimated as 5.0 1013 m2without pre-strain and 7.0 1013 m2 with pre-strain, i.e.still very close to the background value of 1013 m2.Fig. 8 summarizes the measured dislocation densities of dif-ferent types of pillars after compression. It shows that in thecase of the 5.6 lm pillars, coating and pre-straining canresult in a signicant increase in the dislocation density dur-ing deformation. On the contrary, in the small 1.2 lm pil-lars, the dislocation content does not vary much with orwithout coating or pre-straining, and stays at the low mag-nitude of1013 m2, comparable to the initial state prior tocompression.

4. Discussion

The present results in Fig. 2 indicate that coating thealuminum micropillars always leads to strengthening irre-spective of the sizes of the pillars within the range studied,and whether or not pre-straining was applied. However,the eect of the 7% pre-straining is very dierent between

Fig. 8. Estimated dislocation densities in the deformed micropillars.

R. Gu, A.H.W. Ngan /Acta Mlarge and small pillars. For larger 5.6 lm pillars, pre-straining results in a small strengthening eect in boththe coated or uncoated conditions, but for the small1 lm pillars, pre-straining leads to softening in both thecoated and uncoated conditions, and the eect on thecoated condition is larger. The results in Fig. 8 also showthat at the small size of 1.2 lm, the dislocation contentremains close to the initial state in all groups, but this isnot so in the large pillar size of 5.6 lm. These indicatethat the dislocation mechanisms are very dierent betweenthe large and small pillars.

4.1. The larger size (>3.3 lm) regime

Comparison between Figs. 2 and 8 in fact shows that atthe 5.6 lm pillar size, the strength follows the same trendas the residual dislocation density after deformation. At thissize, the strength data in Fig. 2 are in the following descend-ing order: coated with pre-strain > coated without pre-strai-n > uncoated with pre-strain > uncoated without pre-strain, and the dislocation density data in Fig. 8 followexactly the same order. Hence, for the 5.6 lm pillars,strength evidently correlates with the residual dislocationdensity. Fig. 9 plots the strength data vs. the square rootof the residual dislocation density data for the 5.6 lm pil-lars, and a proportionality relation between these two quan-tities is apparent. Hence, the natural conclusion for this sizeregime is that strength is controlled by mutual dislocationinteractions as in traditional Taylor hardening, i.e.r / qp . This conclusion veries the key assumption of Tay-lor hardening in a number of theories on size eect ofstrength [18,19,36], as well as the prediction of dislocationdynamics simulations [21]. In this larger size regime, coatingproduces strengthening because it traps dislocations insidethe pillar and their frequent interactions lead to multiplica-tion, so that the density of mobile dislocations is maintainedat a high level [22]. Without coating, some dislocations canannihilate at the pillars free surface, but since the pillar sizeis not small, not all dislocations can travel freely to the freesurface without interaction with other dislocations. The

Fig. 9. Strength vs. square root of residual dislocation density for the5.6 lm pillars.

erialia 60 (2012) 61026111 6109density is still higher than the initial state with signicantproduction. Pre-straining by 7% in this size regime evidentlyresults in a mild increase in the initial dislocation content(cf. Fig. 3a and b), and indeed in Fig. 8, the residual dislo-cation density is higher with pre-straining in both the coatedand uncoated case. The higher initial dislocation contentfrom pre-straining leads to a small strengthening eect viathe Taylor mechanism as shown in Fig. 2, both for thecoated and uncoated case. In all, in this larger size regime,the eects of coating and pre-straining on strength andresidual dislocation density can be understood as arisingfrom the Taylor hardening mechanism.

4.2. The small size (1 lm) regime

The most remarkable phenomenon to note at the smallpillar size of 1.2 lm is that the residual dislocation

2006;312:1188.

atdensity in Fig. 8 has not grown much from the initial lowvalue, but the strength data in Fig. 2 vary by a large extentdepending on whether the pillar is coated or pre-strained.There is therefore no correlation between strength andthe residual dislocation density, i.e. the results here suggestthat strength in this regime is not controlled by the Taylormechanism. Another important phenomenon from Fig. 2 isthat pre-straining softens the pillars in both the coated anduncoated groups at this small size, and this is in fact oppo-site to the Taylor eect where more existing dislocationsshould produce strengthening. Therefore, the natural con-clusion to draw for this size regime is that strength is con-trolled by the availability of mobile dislocations, in thesense of the dislocation starvation concept as describedin Section 1, or simply that pertaining to the Orowan equa-tion _e qbm if the pillar is not completely depleted of dis-locations. Pre-straining evidently results in more initialdislocations so that strength is decreased at the same strainrate. However, these mobile dislocations zip through thesmall crystal easily, so that the residual dislocation contentstays at the same range. Strength is controlled by the mech-anism(s) by which new dislocations are generated, presum-ably at the pillars surface.

The eects of coating are more dicult to comprehendin this small size regime. Fig. 2 shows that coating stillstrengthens the 1.2 lm pillars as in the larger pillars,but with pre-straining the strengthening eect is less thanthat without pre-straining. The persistently low residualdislocation density even with coating means that the latteris not eective in trapping the generated dislocations in thissmall size regime. Two factors may explain this. First, theapplied stress in this size regime is now very high, and sec-ondly, with the same volume ratio Vw of the coating it isnow very thin, compared to the larger pillars. Therefore,the coating may fail easily when interacted upon by anapproaching dislocation, and may not be able to retainthe latter at the Alcoating interface. However, in Fig. 2,the 1.2 lm coated pillars are always stronger than theuncoated ones, and this is particularly so without pre-straining. This seems to indicate that although the coatingcannot trap the generated dislocations, it can make the dis-location generation process itself more dicult. It is quitepossible that nucleation of dislocations at a coated surfaceof the pillar is more dicult than at a free surface. Thepresent results therefore suggest that the strengtheningmechanism of coating is quite dierent in the small sizeregime, as compared with dislocation trapping in the largersize regime.

5. Conclusions

Eects of coating and pre-straining on deformation ofaluminum micropillars were investigated by compressionexperiments using a nanoindenter with a at-ended dia-mond punch. The results indicate that in the larger size

6110 R. Gu, A.H.W. Ngan /Acta Mregime of a few microns, the strength of the pillars corre-lates well with the square root of the residual density of[15] Greer JR, Nix WD. Phys Rev B 2006;73.[16] Ng KS, Ngan AHW. Acta Mater 2008;56:1712.[17] Shan ZW, Mishra RK, Asif SAS, Warren OL, Minor AM. Nat Mater

2008;7:115.[18] Parthasarathy TA, Rao SI, Dimiduk DM, Uchic MD, Trinkle DR.

Scripta Mater 2007;56:313.[19] Rao SI, Dimiduk DM, Tang M, Parthasarathy TA, Uchic MD,

Woodward C. Philos Mag 2007;87:4777.[20] Noreet DM, Dimiduk DM, Polasik SJ, Uchic MD, Mills MJ. Acta

Mater 2008;56:2988.[21] Rao S, Dimiduk D, Parthasarathy T, Uchic M, Tang M, Woodward

C. Acta Mater 2008;56:3245.[22] Ng KS, Ngan AHW. Acta Mater 2009;57:4902.[23] Greer JR. Mater Res Soc Symp 2007 [0983-LL08-03].[24] Jennings AT, Gross C, Greer F, Aitken ZH, Lee SW, Weinberger CR,dislocations after compression and is therefore controlledby Taylor-type interaction mechanisms. In this regime,coating traps dislocations and pre-straining can furtherraise the initial dislocation contents, and so both are meansof strengthening the material. In the size regime approach-ing 1 lm, however, the residual dislocation density remainsclose to the initial value and pre-straining leads to soften-ing, suggesting that strength is controlled by the availabil-ity of mobile dislocations. Coating cannot trap dislocationseectively but can still produce a signicant strengtheningeect, suggesting that dislocation nucleation at a coatedsurface of the pillar is more dicult.

Acknowledgements

We thank the Electron Microscope Unit of HKU fortheir assistance. The work described in this paper was sup-ported by grants from the Research Grants Council (Pro-ject No. 7159/10E) as well as from the University GrantsCommittee (Project No. SEG-HKU06) of the Hong KongSpecial Administrative Region.

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R. Gu, A.H.W. Ngan /Acta Materialia 60 (2012) 61026111 6111

Effects of pre-straining and coating on plastic deformation of aluminum micropillars1 Introduction2 Experimental details3 Results and analysis3.1 Deformation behavior of tungsten-coated specimens with pre-straining3.2 Strengthening of tungsten-coated specimens3.3 Dislocation distribution in the specimens before compression3.4 Dislocation distribution in the deformed specimens

4 Discussion4.1 The larger size (>3.3m) regime4.2 The small size (?1m) regime

5 ConclusionsAcknowledgementsReferences