Microstructure and stability of nanocrystalline aluminum 6061 … · 2011. 4. 30. · Microstructure and stability of nanocrystalline aluminum 6061 created by large strain machining

Acta Materialia 53 (2005) 4781–4793

www.actamat-journals.com

Microstructure and stability of nanocrystalline aluminum6061 created by large strain machining

M. Ravi Shankar a, Srinivasan Chandrasekar a,*, Alexander H. King b,W. Dale Compton a

a Center for Materials Processing and Tribology, School of Industrial Engineering, Purdue University, 315 North Grant, West Lafayette,

IN 47907-2023, United Statesb School of Materials Engineering, Purdue University, West Lafayette, IN 47907-2023, United States

Received 5 April 2005; received in revised form 30 June 2005; accepted 1 July 2005Available online 24 August 2005

Abstract

We present the properties of aluminum 6061 alloys of various tempers, severely deformed in plane strain at room temperature, bymachining. Various values of strain are introduced into the chip, in a single pass, by varying the rake angle of the tool. Chips cutfrom peak-aged 6061 (T6 condition) are composed of finer microstructures and possess higher hardness than peak-aged 6061 sub-jected to equal channel angular pressing at elevated temperatures. Thermal stability of the chips with different levels of strain is ana-lyzed by studying evolution of Vickers micro-hardness and microstructure after different heat treatments. Chips produced from thepeak-aged temper and over-aged temper soften following heat treatment while those from the solution-treated state first gainstrength before softening. The results are rationalized based on prior studies of the characteristics and kinetics of precipitationand coarsening in Al–Mg–Si systems. The observations also suggest processing routes for consolidation of the chips into bulk formswhile retaining the UFG microstructure with enhanced mechanical properties.� 2005 Acta Materialia Inc. Published by Elsevier Ltd. All rights reserved.

Keywords: Severe plastic deformation; Ultrafine grained microstructure; Aluminium alloys; Thermal stability

1. Introduction

Ultra-fine grained (UFG) and nanocrystalline mate-rials are often harder, stronger and more wear-resistantthan their coarse grained counterparts [1]. Recentdevelopments in the fabrication of UFG materials havefocused on the use of large strain or severe plasticdeformation (SPD) as a method for achieving micro-structure refinement in metals and alloys. Precipita-tion-strengthened, age-hardening aluminum alloys aresome of the most promising candidates for productionof thermally-stable, high-strength UFG materials. Theprecipitate phases are expected to contribute to the

1359-6454/$30.00 � 2005 Acta Materialia Inc. Published by Elsevier Ltd. A

doi:10.1016/j.actamat.2005.07.006

* Corresponding author. Tel.: +1 7654943623; fax: +1 7654945448.E-mail address: [email protected] (S. Chandrasekar).

thermal stability of these materials by pinning thegrain-boundaries and dislocation-structures, thuspreventing coarsening of the UFG microstructureduring subsequent thermo-mechanical processing. Thepresence of second phase particles also aids thegrain refinement process during the SPD of the matrix[2]. Typically, SPD techniques such as equal chan-nel angular pressing/extrusion (ECAP/ECAE) havebeen used to fabricate thermally stable, lightweightnanostructured aluminum alloys through large straindeformation [3–8].

Recently, the case has been made for machining as asimple process for manufacture of UFG materials [9].Chip formation, which involves the introduction of uni-form, large shear strains (�1) in a single pass, also offersa convenient framework for studying the effects of large

ll rights reserved.

Deformation zone

Vo

Tool

Bulk

a o

Chip

a c

Shear plane

Fig. 1. Schematic of plane-strain (2-D) machining.

4782 M.R. Shankar et al. / Acta Materialia 53 (2005) 4781–4793

strain deformation in a variety of materials, includingmany that cannot be subjected to traditional SPD pro-cesses like ECAP or high pressure torsion. Studies ofthe evolution of microstructure during chip formation[10] and in situ determination of the strain field in thedeformation zone during machining [11] have reaffirmedthe case for machining as a convenient method for intro-ducing large plastic strains (1–15) in a single pass. It hasalso been shown that large-strain deformation of Al6061-T6 by machining results in a microstructure thatis significantly finer and harder than that resulting fromECAP [10]. However, from a practical standpoint,machining suffers from a serious drawback : the dimen-sions of the chips are usually too small for fabricatingany structural or mechanical components directly.Therefore, for producing bulk forms, the chips need tobe consolidated using techniques that usually involvesome thermal or thermo-mechanical processing. It isnecessary to characterize the annealing behavior of themachining chips to provide background informationfor efforts aimed at consolidation of these chips whileretaining the nanocrystalline microstructure.

Here, the effect of large strain deformation duringchip formation on the grain-refinement characteristicsof aluminum 6061 in a variety of tempers is examined.Machining chips cut from Al 6061 in the peak-aged(T6), solution-treated and over-aged tempers are ana-lyzed to determine the microstructure refinement, ther-mal-stability/annealing/ageing behavior and resultingmechanical properties. The peak-aged Al 6061 (T6 tem-per) is characterized by an optimum distribution of fineb00 precipitates and some retention of solutes in the ma-trix that maximize the yield strength of the material[12,13]. Large strain deformation of such a matrix dur-ing chip formation may be expected to lead to a very finemicrostructure, reinforced by a fine dispersion of precip-itates that could result in chips endowed with signifi-cantly improved mechanical properties.

For chips produced from Al 6061 in the solution-treated state, the absence of precipitates in the matrixis expected to lead to a microstructure that is strength-ened primarily through grain refinement. The differencein the microstructure between the chips cut from solu-tion-treated 6061 and peak-aged 6061 can, therefore,be used to discern the role played by a fine dispersionof second-phase precipitates in the grain refinementand subsequent strengthening by the deformation. Fur-thermore, the precipitation kinetics of chips cut fromsolution-treated 6061 is expected to be significantlyaccelerated vis-a-vis the undeformed bulk, as a resultof the large defect concentration arising from the defor-mation. Appropriate thermal processing routes for con-solidation can therefore be developed to exploit thiskinetics and produce a fine dispersion of precipitatesin the deformed matrix while retaining the microstruc-ture resulting from the large strain deformation; this

would result in the final material having considerablyenhanced mechanical properties.

In the over-aged state, precipitation of solutes fromthe solid solution is nearly complete and so is the conver-sion of the fine b00 to coarse B 0 [12]. This coarse dispersionis not likely to be as effective at strengthening the micro-structure as the b00 in the peak-aged temper. However,compared to the b00 phase, at the annealing temperaturesthat will be considered here, B 0 is not likely to coarsenrapidly. Therefore, the microstructure evolution duringannealing of chips produced from the over-aged tempershould purely be a function of the ability of these coarseprecipitates to pin and stabilize the microstructure.

1.1. Machining

Large strain deformation results from chip formationin its simplest manifestation, i.e. plane-strain (2-D)machining (Fig. 1). Plane-strain machining is character-ized by a sharp, wedge-shaped tool that removes a presetdepth of material (ao) by moving in a direction perpen-dicular to its cutting edge [14]. Chip formation occurs byconcentrated shear within a narrow deformation zone,often idealized by a ‘‘shear plane’’ [14–16]. Most of thegrain refinement associated with the formation of theUFG chips has been attributed to the large shear strainsimposed in this deformation zone [10]. The geometry ofthe deformation zone and the associated shear strain aredetermined by the shear plane angle (/) and the rake an-gle (a). In Fig. 1, the rake angle (a) is taken to be positivewhen measured clockwise with respect to the normal tothe workpiece surface and negative when measuredcounter-clockwise. Fig. 1 illustrates the case of a positiverake angle tool.

The shear strain (c) imposed in the chip during chipformation is given by

c ¼ cos asin/ cosð/� aÞ ; ð1Þ

M.R. Shankar et al. / Acta Materialia 53 (2005) 4781–4793 4783

where / is calculated from a measurement of ao and ac(Fig. 1), as

tan/ ¼aoaccos a

1� aoacsin a

. ð2Þ

2. Experimental procedure

The characteristics of chips produced by large straindeformation of Al 6061 alloy in three distinct tempers,viz. peak-aged Al 6061 (T6 temper), over-aged 6061and solution-treated-6061, were studied through adapta-tion of 2-D machining as a large strain deformation pro-cess. Table 1 gives the experimental conditions.

First, Al 6061-T6 plates were machined on a custom-built plane-strain machining setup with high speed steeltools of different rake angles. A sufficiently low cuttingvelocity of 10 mm/s was employed to minimize strain-rate effects and temperature rise during chip formation.In situ measurements of the deformation zone tempera-ture made using infra-red thermography confirmed thatthe temperature rise during chip formation was negligi-ble. The excellent machinability of 6061-T6 facilitatedthe imposition of a wide range of strains in the chipby varying the rake angle between +20� and �40�.The shear strain values were estimated using Eqs. (1)and (2) after measuring the deformed chip thickness(ac) and undeformed chip thickness (ao). The annealingbehavior of the chips was characterized by heating thechips to various temperatures in a furnace (in air), fol-lowed by determination of the Vickers micro-hardness(50 g load) and microstructure characteristics. Anneal-ing was done at 175 �C for different lengths of time tostudy the thermal stability at the characteristic ageingtemperature for 6061. Annealing experiments were alsodone at 210 �C on the chips produced with different val-ues of strain to determine any possible effect of deforma-tion on the over-ageing characteristics.

The microstructure of the chips, as-machined andafter annealing, was characterized using a JEOL

Table 1Experimental conditions

Bulk material characteristics Machining conditions

Al 6061-T6 (peak-aged)

Hardness: 110 kg/mm2 Vc = 10 mm/s, ao = 150 lmGrain size �75 lm Rake angle (a): +20� to �40�

Over-aged 6061 (obtained by annealing 6061-T6 at 210 �C for 15 h)

Hardness: 95 kg/mm2 Vc = 10 mm/s, ao = 150 lmGrain size �75 lm Rake angle (a): +20� and �5�

Solution-treated 6061 (obtained by annealing 6061-T6 at 550 for 10 h)

Hardness: 72 kg/mm2 Vc = 10 mm/s, ao = 150 lmGrain size �30 lm Rake angle (a): +20� and �5�

Vickers indentation load = 50 g for all of the hardness measurements.a All annealing was done in air.

2000FX transmission electron microscope (TEM) oper-ating at 200 kV. The chips tended to curl considerablyand those produced with the negative rake angle toolstended to fragment. Hence, wedge polishing followedby ion milling was used, rather than electrolytic thin-ning, to produce most of the TEM chip specimens. Inthe wedge polishing, the specimens were first mechani-cally thinned by abrasive polishing to form shallowwedges. Each wedge was then mounted on a copper slotgrid and ion milled for a very short duration of time (�5to 10 min) in a Gatan Model 600 dual ion mill to createan electron transparent specimen. Care was taken toavoid any prolonged heating of the specimens duringthe mechanical thinning and ion milling in order toensure the integrity of the microstructure of theas-machined chip.

Whenever continuous and relatively flat chips wereproduced, an electrolytic thinning technique was usedfor the TEM specimen preparation. In this case, thechips were first ground to a thickness of �100 lm usingan abrasive grinding wheel. Three millimeter diameterdisks were then punched out of the ground chip samples.The disk specimens were made electron transparent byelectrolytic jet-thinning (Struers Tenupol-5) using asolution of 25% HNO3–75% CH3OH at �10 �C, 9.5 V.

Second, chips were cut from over-aged Al 6061. Forthis purpose, peak-aged Al 6061-T6 plates were first sub-jected to a prolonged heat treatment (over-ageing) at210 �C for 15 h. The choice of this temperature–timecondition was guided by a prior TEM study of the pre-cipitation sequence in Al 6061 [12]. Vickers hardness andmicrostructure of the chips cut from these plates werecharacterized. The annealing behavior of the chips wasanalyzed by studying the microstructure and hardnessafter heat treatment at 175 �C for different lengths oftime. The choice of 175 �C was determined by the factthat the precipitates in the chips are likely to be rela-tively inert at this temperature; this is a consequenceof the bulk plates being initially subjected to a pro-longed heat treatment at 210 �C. Therefore, the anneal-ing behavior of the chips is likely to be reflective of the

Shear strain (c) Chip annealing conditionsa

See Fig. 2 175 �C for 1–15 h210 �C for 1–10 h

a = +20� ! c = 3.2 175 �C for 1–15 ha = �5� ! c = 4

a = +20� ! c = 2.7 175 �C for 1–21 ha = �5� ! c = 4 150 �C for 1–10 h

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annealing behavior of an UFG matrix that is pinned bycoarse precipitates to the exclusion of any other soften-ing phenomena.

Lastly, to study the ageing behavior of Al 6061 afterlarge strain deformation, plates of Al 6061 were solu-tion-treated by heating them at 550 �C for 10 h and thenquenching them in water at room temperature. Chipswere then produced and characterized immediately priorto any natural ageing. The ageing characteristics of thesechips were studied by heat treating them at 150 and175 �C for different lengths of time. The temperatureof 175 �C was chosen since it is the typical temperatureused to age-harden solution-treated 6061 [17]. The heattreatment at 150 �C (a temperature usually too low toeffectively age 6061) was based on prior ECAP workthat indicated the sufficiency of lower-than-usual ageingtemperatures to effectively harden severely-deformed,solution-treated materials [7].

Table 1 summarizes the experimental conditionsincluding bulk material characteristics, machiningparameters and annealing conditions for all of the chips.

3. Experimental results

The utility of machining for introducing a range ofshear strains in a single pass is evident from Fig. 2,which shows the variation of chip shear strain with toolrake angle for Al 6061-T6. The shear strain varies be-tween 1.7 with a +20� rake tool and 5.2 for a �20� raketool. When machining was done with tools of rake anglesmaller than �20� (i.e. �30� and �40�), the cuttingforces were significant (indicative of larger strains) andthere was some flow of material in a direction normal

Rake angle (Degrees)

niartsraeh

S

-300

2

4

6

8

10

)γ(

-20 -10 0 10 20

Fig. 2. Variation of shear strain (c) with rake angle (a) for chips cutfrom peak-aged 6061.

to the plane of machining (side-flow). Under such condi-tions, the deformation field is no longer one of planestrain and the applicability of Eqs. (1) and (2) is ques-tionable. Hence, the results for the �30� and �40� rakeangle chips are not reported here. The shear strains inthe chips cut from the over-aged 6061 and solution-trea-ted 6061 were also estimated using Eqs. (1) and (2).These are given in Table 1 for the +20� and �5� rake an-gle tools. Significant side-flow of material precludedanalysis of strains in chips cut with tools of rake anglessmaller than �5� for the over-aged and solution-treated6061.

3.1. Peak-aged Al 6061 (6061-T6)

Fig. 3 shows the variation of the measured Vickershardness (Hv) with shear strain (c) for the as-machinedAl 6061-T6 chips. The broken trend-line represents apower-law fit of the chip hardness (strength) data. TheVickers hardness of the bulk 6061-T6 is 110 kg/mm2.The chips produced with the highest shear strain (�8)are up to 50% harder than the 6061-T6 bulk.

Thermal stability of the chips was analyzed by track-ing the variation in hardness values following theirannealing at 175 and 210 �C for different lengths of time.Chips with c = 3.2 and c = 5.2 were chosen for this anal-ysis. Fig. 4 illustrates the evolution of the Vickers hard-ness with annealing time at 175 �C. After annealing for1 h, the hardness value is seen to decrease rapidly froma value of about 150 to about 135 kg/mm2. It is interest-ing to note that the annealed chip hardness value of135 kg/mm2 is about the same as that obtained afterwarm ECAP of 6061-T6. Following this rapid initial de-crease, the hardness value decreases at a much slower

Fig. 3. Variation of Vickers hardness (Hv) with shear strain (c) forchips cut from peak-aged 6061. Hardness of peak-aged bulk is 110 kg/mm2. Vickers indentation load: 50 g.

Fig. 4. Evolution of Vickers hardness (Hv) with annealing time at175 �C for chips cut from peak-aged 6061.

M.R. Shankar et al. / Acta Materialia 53 (2005) 4781–4793 4785

rate during further annealing. After annealing for 10 hor more at 175 �C, the more highly strained chip isslightly softer than that produced with a smaller strain(Fig. 4). This hardness difference, though very small, isstill statistically significant. This is indicative of a possi-ble effect of cold work on the over-ageing characteristics.

Fig. 5 shows the Vickers hardness of the chips pro-duced with different levels of strain as a function ofannealing time at 210 �C. For reference, an undeformedbulk sample was also annealed and its hardness re-corded. The rate of decrease in hardness of the more se-verely deformed chip (c = 5.2) is considerably morerapid than that of the bulk and somewhat greater than

Fig. 5. Evolution of Vickers hardness (Hv) with annealing time at210 �C for chips cut from peak-aged 6061.

that of the chip produced with c = 3.2. It is also interest-ing to note the one-to-one relationship between defor-mation strain and the resulting hardness followingannealing at 210 �C. Material deformed to the smalleststrain, viz. the undeformed bulk Al 6061-T6 (c = 0),softens the slowest and remains the hardest comparedto the chips following the annealing. Material deformedto the largest strain value (c = 5.2) anneals the fastestand also reaches the lowest value of hardness afterannealing.

Fig. 6 shows TEM micrographs illustrating the effectof the strain on the resulting chip microstructure. Whilethe grain morphologies of chips produced with c = 3.2and c = 5.2 are similarly equi-axed and typically finerthan 100 nm, the less strained c = 1.7 chips are com-posed of elongated and coarser grains of width�150 nm. The diffraction pattern of the c = 1.7 chipsis characteristic of a matrix made up of sub-grains oflow misorientation angles (Fig. 6(a)), while the nearlyuniform rings around the central bright spot in the caseof the c = 3.2 and c = 5.2 chips are characteristic oflarge values of misorientations between the sub-grains(Figs. 6(b) and (c)).

The chips annealed at 175 and 210 �C were examinedin the TEM to track coarsening effects. After 1 h ofannealing at 175 �C, the grain size was found to increasesubstantially (Figs. 7(a) and (c)) from that of the as-machined chip, although the resulting grains are stillsub-micron in size. The same effect was encounteredafter annealing at 210 �C for 1 h as well (Fig. 8). Fig. 7also shows that prolonged annealing for more than 1 hat 175 �C does not lead to significant continued coarsen-ing compared to that encountered after 1 h of annealing.The occurrence of coarse, spherical precipitate particles,typically at grain boundaries, may be noted. The size ofthese precipitates appears to be much larger than thoseobserved in the peak-aged material [12]. These precipi-tates at the grain boundaries are likely to impede coars-ening of the microstructure during extended annealingby pinning the grain boundaries.

3.2. Over-aged Al 6061

Chips were cut from a bulk over-aged 6061 samplewith a Vickers hardness of 95 kg/mm2. The hardnessof the c = 3.2 chips was 122 kg/mm2 while that of thec = 4 chips was 134 kg/mm2 (Fig. 9). For the same shearstrain value of 3.2, chips produced from peak-aged 6061(6061-T6) were 1.38 times as hard as the bulk while thechips produced from the over-aged 6061 were 1.27 timesas hard as the corresponding bulk material.

Fig. 9 illustrates the variation of hardness withannealing time at 175 �C. During prolonged annealing,the hardness of the over-aged bulk material is relativelyunaffected. The chips however show a steady decrease inhardness that appears to be independent of the extent of

Fig. 6. Bright field TEM images of chips cut from peak-aged 6061 with different induced strains. Insets are the corresponding selected area diffractionpatterns. (a) c = 1.7. Grain size (along smaller dimension) �150 nm. (b) c = 3.2. Grain size �80 nm. (c) c = 5.2. Grain size �80 nm.

Fig. 7. Bright field TEM images of chips cut from peak-aged 6061 and annealed at 175 �C for different lengths of time. Arrows indicate the coarsespherical b phase. Typical grain sizes are �250 to 300 nm. (a) c = 3.2 chip annealed at 175 �C for 1 h. (b) c = 3.2 chip annealed at 175 �C for 10 h.(c) c = 5.2 chip annealed at 175 �C for 1 h. (d) c = 5.2 chip annealed at 175 �C for 10 h.

4786 M.R. Shankar et al. / Acta Materialia 53 (2005) 4781–4793

Fig. 8. Bright field TEM images of chips cut from peak-aged 6061 and annealed at 210 �C for 1 h. (a) c = 3.2 chip annealed at 210 �C for 1 h. Grainsize �250 nm (b) c = 5.2 chip annealed at 210 �C for 1 h. Grain size �300 nm.

Fig. 9. Evolution of Vickers hardness (Hv) with annealing time at175 �C for chips cut from over-aged 6061. For reference, the hardnessevolution of the bulk material is also shown.

Fig. 10. Bright field TEM images of chips cut from over-aged 6061 with differpatterns. Typical grain size �160 nm. (a) c = 3.2. (b) c = 4.0.

M.R. Shankar et al. / Acta Materialia 53 (2005) 4781–4793 4787

prior deformation. Unlike the case of the chips cut fromthe peak-aged 6061, the more highly strained chips fromthe over-aged 6061 do not anneal at an accelerated rateto a lower value of hardness.

Fig. 10 illustrates the microstructure of as-machinedchips with different levels of strain. The typical grainsizes are �160 nm, nearly double the grain size of chipscut from peak-aged 6061-T6 for comparable values ofstrain. The insets of the diffraction patterns in Fig. 10show relatively uniform intensities of the diffractionrings around the central spot and particularly so withthe more highly strained chip (Fig. 10(b)).

Fig. 11 shows the microstructures of the chips after1 h anneal at 175 �C. The resistance to significant graincoarsening is evident from this figure in contrast to theannealing behavior of chips machined from the peak-aged 6061 (Fig. 7). The grain-size after the annealing(�160 nm) is about the same as in the as-machined statealthough it is apparent from Fig. 9 that there is someassociated decrease in the hardness following theannealing.

ent levels of strain. Insets are the corresponding selected area diffraction

Fig. 11. Bright field TEM images of chips cut from over-aged 6061 and annealed at 175 �C for 1 h. Typical grain size �160 to 200 nm. (a) c = 3.2chip annealed for 1 h at 175 �C. (b) c = 4 chip annealed for 1 h at 175 �C.

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3.3. Solution-treated Al 6061

The hardness of the chips createdwith the 20� rake tool(c = 2.7) was 120 kg/mm2 while that generated with the�5� rake tool (c = 4) was 127 kg/mm2; the bulk solu-tion-treated 6061 had a hardness of 72 kg/mm2 (Table 1).

The chips were heat treated at 175 and 150 �C forvarious lengths of time to study their ageing characteris-tics. Figs. 12 and 13 show the variation of Vickers hard-ness with ageing time at 175 and 150 �C, respectively.The ageing characteristics of the chips at 175 �C appearto be independent of the prior strain (Fig. 12); both thehighly strained and less highly strained chips are seen toage very rapidly compared to the undeformed bulk. Fur-thermore, the chips attain a peak-hardness value of132 kg/mm2 after 1 h of annealing at 175 �C. Prolongedannealing at 175 �C caused a steady decline in the

Fig. 12. Evolution of Vickers hardness (Hv) with ageing time at 175 �Cfor chips cut from solution-treated 6061.

Fig. 13. Evolution of Vickers hardness (Hv) with ageing time at 150 �Cfor chips cut from solution-treated 6061.

hardness value of the chips and after 21 h of annealing,the hardness of the bulk is indistinguishable from that ofthe chips. The ageing characteristics of the chips werealso studied at 150 �C, to explore the possibility of real-izing higher values of hardness (strength) by loweringthe ageing temperature and inhibiting any coarseningphenomena that may reduce the strength. Fig. 13 showsthe variation of hardness with ageing time 150 �C. Notonly is a higher hardness value (�140 kg/mm2) achievedfollowing the ageing for 1 h, but this value remains rel-atively unchanged even after 5 h of ageing. However,after 10 h of ageing at 150 �C, the hardness declinedfrom 140 kg/mm2 to 132 kg/mm2, which is the same asthat recorded on chip samples after 1 h of ageing at175 �C.

Fig. 14(a) shows a TEM micrograph of a c = 2.7 chipmachined from solution-treated 6061. The grains areelongated with a typical width of 150–200 nm; this

Fig. 15. Bright field TEM image of c = 2.7 chips cut from solution-treated 6061 and annealed at 175 �C for 1 h. Grain size �400 nm.

M.R. Shankar et al. / Acta Materialia 53 (2005) 4781–4793 4789

microstructure is characteristic of a relatively low levelof SPD compared to say the strain (c = 5.2) achievedwith a �20� rake tool in peak-aged 6061. Figs. 14(b)and (c) show evolution of the microstructure duringthe ageing at 150 �C. A side-effect of the ageing processis grain growth that coarsens the UFG microstructureand contributes to a reduction in hardness during theprolonged ageing. At 175 �C, as seen in Fig. 15, thereis significant coarsening of the microstructure (grain size�400 nm). But, the hardness value (�132 kg/mm2) isstill greater than that of the solution-treated bulk(�72 kg/mm2), the bulk peak-aged 6061 (�110 kg/mm2) and the chips cut from the solution-treated bulk(�120 kg/mm2). The microstructure that is strengthenedby a fine dispersion of precipitates in several of thegrains in Fig. 14(c), appears to be very similar to thatshown in Fig. 15.

The diffraction patterns shown as insets in Figs. 14and 15 indicate qualitatively an evolution in the misori-entation between the grains during the ageing process.The diffraction pattern of the as-machined chip is char-acteristic of a microstructure made up of grains that areconsiderably misoriented with respect to one another.After the recovery and grain-growth that occur in tan-dem with the precipitation during ageing, the diffractionpattern is characteristic of a microstructure composed oflow-angle boundaries.

Fig. 14. Bright field TEM images of as-machined chips cut from solution-(a) c = 2.7 chip cut from solution treated 6061. Grain size �110 nm. (b) c =chip annealed at 150 �C for 10 h. Arrows indicate a fine dispersion of pr

4. Discussion

Strengthening in UFG multi-phase, Al alloys canoccur due to a combination of solid-solution strengthen-ing, precipitation strengthening and grain size strength-ening. The same phenomena are likely to also affect theformation of UFG microstructures in precipitation-strengthened aluminum alloys during large strain defor-mation, by modifying the micro-mechanics of plastic

treated 6061 and after ageing at 150 �C for different lengths of time.2.7 chip annealed at 150 �C for 5 h. Grain size �160 nm. (c) c = 2.7

ecipitates within a grain. Grain size �270 nm.

4790 M.R. Shankar et al. / Acta Materialia 53 (2005) 4781–4793

flow and grain refinement. This study has explored theinterplay of the role of precipitate size and the single-pass, large-strain deformation as well as the effect ofthe large strain deformation on the ageing characteristicsof the UFG microstructures.

4.1. Large strain deformation of precipitation-

strengthened 6061

4.1.1. Grain refinement

Large strain deformation by machining is a simpleroute for studying the effect of precipitates on the scaleand stability of UFG Al 6061 microstructures. A rangeof shear strains and a variety of resulting microstruc-tures can be produced by chip formation through anappropriate choice of the tool rake angle. However,for the same rake angle,two different materials or differ-ent tempers of the same alloy can lead to different valuesof shear strain in the chip. A 20� rake angle tool is seento produce a shear strain of 1.7 in peak-aged 6061 and ashear strain of 3.2 in over-aged 6061. This is not surpris-ing considering that in machining the geometry of thedeformation field is not defined a priori. This does notimply that strains cannot be introduced in a controllablemanner. From Fig. 2 it is possible to determine experi-mentally, the one-to-one correspondence between thetool rake angle and the shear strain in the chip for a spe-cific material state; this relation can then be used to im-pose a controlled level of strain in the chip.

Fig. 3 shows that the chip hardness is much greaterthan that of the bulk in peak-aged 6061 even at the low-er strains, while the incremental increase in the chiphardness at larger strains is diminishing. A switch-overof the grain morphology from elongated structures atsmaller chip strains (Fig. 6(a)) to fine, equi-axed grainsoccurs at larger strains (Figs. 6(b) and (c)). The grainsize in the c = 1.7 chips measured along the smallestdimension is �150 nm while the equi-axed grains inthe c = 3.2 and c = 5.2 chips are typically 80 nm in size.Such a switch-over has also been observed in ECAP [3].A similar switch-over in grain morphology was not ob-served in the chips cut from the over-aged 6061 sincelarge values of shear strain (c = 3.2) and an equi-axedgrain microstructure were obtained even with the 20�rake angle tool, the most positive rake angle studied.

The relative increase in hardness of the chip over thatof the bulk is found to depend on the characteristics ofthe bulk material. The chip-to-bulk hardness ratio in thecase of the peak-aged 6061 is 1.38 and in the over-agedmaterial is 1.27, even though the initial grain sizes of thebulk material in both cases are very similar(�75 lm).This difference can be attributed to more effective trap-ping of dislocations and associated refinement of micro-structure brought about by a fine dispersion ofprecipitates in the case of the peak-aged material (typi-cally b00 [12,13]) as compared to the coarser precipitates

in the over-aged material (typically B 0 [12]). A finerdispersion provides a larger number of sites at whichsub-grain boundaries are pinned thus preventing anysignificant coarsening during dynamic recovery thatcan occur in the course of the large deformation. Fur-thermore, the greater concentration of solutes in thepeak-aged state compared to the over-aged state canlead to a greater resistance to dislocation motion andcontribute to larger values of hardness (strength). Ithas been found in rolling of a similar Al–Mg–Si alloythat the precipitates do indeed break up leading to afiner dispersion in the aluminum matrix [18]. Similar in-stances of precipitates breaking down during deforma-tion are also reported in SPD of Al 5083 [19], Al 7050[20] and an Al–Cu alloy [21]. We anticipate a similarprecipitate breaking phenomenon here that results infragmented precipitates. These contribute to the finemicrostructure and associated strengthening.

The hardness values of the chips cut from the peak-aged temper are generally greater and the scale of themicrostructure observed here is generally finer than thatencountered in warm-ECAP of peak-aged 6061 [3]. Thiscould be a result of recovery and loss of dislocations inthe ECAP, and thus ineffective grain refinement due tothe heating associated with the deformation in the warmECAP process. The room temperature deformationeffected here during machining involves no such soften-ing aspects and thus leads to larger values of chiphardness.

4.1.2. Thermal stability

The thermal stability of the UFG 6061 microstruc-tures is of particular interest because of the possible uti-lization of the machining chips in the fabrication of bulknanocrystalline materials through powder consolidationtechniques. In the case of chips cut from the peak-agedmaterial, the hardness value decreases rapidly in the ini-tial stages of annealing. Prolonged annealing leads to aslow decline in the hardness, as seen in Figs. 4 and 5.The initial decrease in hardness coincides with the rapidincrease in the grain size as is evident from Figs. 6–8.Rapid grain growth from what appears to be a relativelyunstable configuration of precipitates and grain bound-aries to a more stable configuration, rather than a mere‘‘clean-up’’ of the dislocations from the interior of thegrains, appears to be the operative phenomenon here.A study of the annealing behavior beyond 1 h did notindicate any continued increase in grain size. This is con-sistent with the notion that a relatively stable micro-structure forms after about 1 h of annealing at 175 �C,one that is stabilized by the precipitates which pin thegrain boundaries preventing grain growth. The largestrain deformation is usually associated with fractureof the precipitates and this leads to the formation ofconsiderably finer precipitates with a large surface en-ergy component. Furthermore, it has been suggested in

M.R. Shankar et al. / Acta Materialia 53 (2005) 4781–4793 4791

the case of severely deformed Al–Mg–Si alloys that thesheared, fine precipitates, which have lost their coher-ency with the matrix as a result of deformation, wouldrapidly transform to the equilibrium b phase [18]. Theresulting stable b phase is characterized by relativelycoarse spherical precipitates, similar to those observedhere in Fig. 7. The stable b phase that would normallyrequire prolonged heat treatment to form at 175 or210 �C in the undeformed bulk, has formed quite rapidlyin the machining chips due to the instability introducedin the b00 precipitates as a result of large defect concen-tration and possible breaking up of coherent precipi-tates.

The hardness vs. annealing time data in Figs. 4 and 5show that the flow stress continues to decrease withannealing time at a rate determined by the extent of coldwork during large strain deformation. The c = 5.2 chipsare somewhat softer following annealing than thec = 3.2 chips. The difference between the coarseningrates in these two material states might be a result ofthe differing sizes of fragmented precipitates and differ-ing defect concentrations left behind by the deforma-tion. The greater the defect concentration in terms ofvacancies produced by cold work, the greater is the dif-fusivity of the solute atoms through the matrix andtherefore, the greater the rate of diffusion-controlledcoarsening [22]. Also, on deforming to larger values ofstrain, the precipitates are likely to break up further toform an increasingly finer distribution. Such a fine dis-persion could be expected to coarsen at a very rapid rateduring heat treatment, due to the increased driving forceprovided by the larger interfacial energy [18]. For thesereasons, it is reasonable to expect that precipitates inchips deformed to larger values of strain coarsen at afaster rate, and, therefore, these chips, possess lowerstrength following annealing.

From Figs. 10 and 11 it is apparent that chips cutfrom the over-aged state are relatively immune to graingrowth following annealing at 175 �C. Since the bulkmaterial was in the over-aged state, further coarseningof the precipitates is not likely to be the key aspect influ-encing the softening of the material. Rather, as seenmore clearly in Fig. 11(a), a regular recovery, togetherwith coarsening of some of the broken precipitates, fol-lowed by gradual but highly restrained grain growth,appears to be the mechanism that determines the soften-ing of the chips created from the over-aged 6061.

4.2. Ageing characteristics of the chips

The properties of the chips machined from the solu-tion-treated state suggest new and exciting possibilitiesfor consolidation of these chips by thermo-mechanicalprocessing. Since some heating is usually involved inany consolidation and if that heating can be associatedwith a hardness increase rather than softening, then this

phenomenon can be exploited during the consolidationto produce bulk forms with enhanced strength. Thatis, chips cut from the solution-treated bulk may be ther-mo-mechanically processed in such a way that they gainhardness during the consolidation, thus leading to sig-nificantly stronger bulk forms. From Figs. 12 and 13,it is apparent that chips machined from the solution-treated state, unlike those from the peak-aged state orthe over-aged state, gain in strength following heat treat-ment. This clearly implies that the rate of hardening dueto precipitation exceeds any softening due to the coars-ening of the microstructure from grain growth. Precipi-tation and grain growth are both diffusion driven, butmutually counter-acting phenomena, as far as the flowstress is concerned; so the relative rate of each processat a given temperature would determine the finalstrength of the material. By comparing Figs. 14(b) and15, it is apparent that higher temperatures lead to agreater coarsening of the grains and lower hardness val-ues (see also Figs. 12 and 13). The deformed microstruc-ture significantly enhances the mobility of the soluteatoms in the matrix. This when combined with the verylarge number of nucleation sites made available by thedislocations and the grain boundaries, may be expectedto lead to a rapid formation of a fine dispersion of pre-cipitates. The accelerated formation of a fine dispersionof precipitates can occur even at a relatively low temper-ature. The precipitation hardening of the matrix maythen be expected to prevail over the softening effectsdue to grain growth in a low-temperature annealing pro-cess. It is also interesting to note that the hardness andthe microstructure resulting from annealing of the solu-tion-treated chips at 150 �C for 10 h are nearly identicalto those realized following annealing at 175 �C for 1 h.Both are characterized by a fine dispersion of precipi-tates in the interior of a grain and a relatively large grainsize of about 500 nm. These results are similar to theageing characteristics of solution-treated 6061 [7] andan Al–Sc alloy [23] processed by ECAP. The same hard-ness value of �132 kg/mm2 is obtained in the case of theannealed peak-aged or the as-machined over-aged chips,which are both typically composed of 200 nm grains.Hence, significantly improved mechanical properties inthe case of precipitation-strengthened materials caneither be obtained through simple grain refinement inone set of tempers or by producing an especially fine dis-persion of precipitates that accomplishes the same taskin a different temper.

These results reaffirm the case for machining as a sin-gle-pass process for studying large strain deformation ofmaterials. From a practical standpoint, they suggest aparameter window for thermal processing of the chips.In the case of the chips machined from the peak-agedand the over-aged materials, softening during thermaltreatment is inevitable. However, by keeping the pro-cessing temperature around 175 �C it is possible to

4792 M.R. Shankar et al. / Acta Materialia 53 (2005) 4781–4793

retain some of the increase in strength resulting from thelarge strain deformation. In the case of the chips cutfrom the solution-treated 6061, a significant increase instrength can occur due to the thermal treatment. Thisaspect combined with the retention of a UFG matrixis likely to be of advantage during thermally-assistedconsolidation of these chips into bulk forms.

5. Conclusions

We have demonstrated the effectiveness of plane-strain machining as a technique for studying large straindeformation of aluminum alloys in various tempers. Avariety of strains, microstructures and mechanical prop-erties are shown to be obtained by appropriate choice ofthe tool rake angle. Furthermore, the large strains areimposed in a single pass of deformation.

A very fine microstructure composed of grains smal-ler than 100 nm is found to result from chip formationin peak-aged 6061. The scale of the microstructure isgenerally finer, and the associated hardness value forthe chips somewhat larger, than those reported forECAP at elevated temperatures, implying a strong effectof deformation temperature on the resulting microstruc-ture. Chips cut from the peak-aged Al 6061, with largervalues of strain, are less stable during annealing thantheir less-strained counterparts. This difference in stabil-ity may be rationalized by considering the effect of thedeformation on the precipitates, microstructure andtransport properties. Coarsening is found to be a twostep process. Initial rapid grain growth is seen to giveway to a more stable microstructure that gradually soft-ens during prolonged annealing. The b00 phase that isresponsible for the peak-aged strength of Al 6061, fol-lowing severe deformation is expected to rapidly trans-form to the relatively coarse equilibrium b phase. Thisstable b phase, however, is still capable of effectively pin-ning the grain boundaries thus preventing furthercoarsening.

Deformation of over-aged 6061 by machining resultsin relatively coarser grains and less hardening comparedto similarly deformed peak-aged material. This can beattributed to less-efficient dislocation trapping and stabi-lizing by the coarser B 0 phase that is dispersed in theover-aged bulk matrix. Annealing of this UFG micro-structure results in gradual softening with only minorchanges in the microstructure; this is to be contrastedwith the annealing behavior of chips cut from thepeak-aged material. Chips from the solution-treatedmaterial, on the other hand, first harden following heattreatment and then soften after prolonged heat treat-ment. Accelerated formation of very fine precipitatescauses the initial rapid increase in strength in this mate-rial, and this effect is then counter-acted by grain growthand coarsening of the fine precipitates.

The observations suggest thermal processing routesfor consolidation of the chips into bulk forms whileretaining the UFG microstructure with enhancedmechanical properties.

Acknowledgements

We thank the Department of Energy (Grant4000031768 via UT-Batelle), Ford Motor Company,Oak Ridge National Laboratory (ORNL) and the Stateof Indiana�s 21st Century Research and TechnologyFund for support of this work. Additional thanks arealso due to Drs. Ray Johnson (ORNL) and AndrewSherman (Ford) for their encouragement of the studies.

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