Microstructure and stability of nanocrystalline aluminum 6061 created ...

Acta Materialia 53 (2005) 4781–4793 www.actamat-journals.com

Microstructure and stability of nanocrystalline aluminum 6061 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 States b 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 2005 Available online 24 August 2005

Abstract We present the properties of aluminum 6061 alloys of various tempers, severely deformed in plane strain at room temperature, by machining. Various values of strain are introduced into the chip, in a single pass, by varying the rake angle of the tool. Chips cut from peak-aged 6061 (T6 condition) are composed of finer microstructures and possess higher hardness than peak-aged 6061 subjected to equal channel angular pressing at elevated temperatures. Thermal stability of the chips with different levels of strain is analyzed by studying evolution of Vickers micro-hardness and microstructure after different heat treatments. Chips produced from the peak-aged temper and over-aged temper soften following heat treatment while those from the solution-treated state first gain strength before softening. The results are rationalized based on prior studies of the characteristics and kinetics of precipitation and coarsening in Al–Mg–Si systems. The observations also suggest processing routes for consolidation of the chips into bulk forms while retaining the UFG microstructure with enhanced mechanical properties. Ó 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 materials are often harder, stronger and more wear-resistant than their coarse grained counterparts [1]. Recent developments in the fabrication of UFG materials have focused on the use of large strain or severe plastic deformation (SPD) as a method for achieving microstructure refinement in metals and alloys. Precipitation-strengthened, age-hardening aluminum alloys are some of the most promising candidates for production of thermally-stable, high-strength UFG materials. The precipitate phases are expected to contribute to the *

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

thermal stability of these materials by pinning the grain-boundaries and dislocation-structures, thus preventing coarsening of the UFG microstructure during subsequent thermo-mechanical processing. The presence of second phase particles also aids the grain refinement process during the SPD of the matrix [2]. Typically, SPD techniques such as equal channel angular pressing/extrusion (ECAP/ECAE) have been used to fabricate thermally stable, lightweight nanostructured aluminum alloys through large strain deformation [3–8]. Recently, the case has been made for machining as a simple process for manufacture of UFG materials [9]. Chip formation, which involves the introduction of uniform, large shear strains (1) in a single pass, also offers a convenient framework for studying the effects of large

1359-6454/$30.00 Ó 2005 Acta Materialia Inc. Published by Elsevier Ltd. All rights reserved. doi:10.1016/j.actamat.2005.07.006

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

strain deformation in a variety of materials, including many that cannot be subjected to traditional SPD processes like ECAP or high pressure torsion. Studies of the evolution of microstructure during chip formation [10] and in situ determination of the strain field in the deformation zone during machining [11] have reaffirmed the case for machining as a convenient method for introducing large plastic strains (1–15) in a single pass. It has also been shown that large-strain deformation of Al 6061-T6 by machining results in a microstructure that is significantly finer and harder than that resulting from ECAP [10]. However, from a practical standpoint, machining suffers from a serious drawback : the dimensions of the chips are usually too small for fabricating any structural or mechanical components directly. Therefore, for producing bulk forms, the chips need to be consolidated using techniques that usually involve some thermal or thermo-mechanical processing. It is necessary to characterize the annealing behavior of the machining chips to provide background information for efforts aimed at consolidation of these chips while retaining the nanocrystalline microstructure. Here, the effect of large strain deformation during chip formation on the grain-refinement characteristics of 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 analyzed to determine the microstructure refinement, thermal-stability/annealing/ageing behavior and resulting mechanical properties. The peak-aged Al 6061 (T6 temper) is characterized by an optimum distribution of fine b00 precipitates and some retention of solutes in the matrix that maximize the yield strength of the material [12,13]. Large strain deformation of such a matrix during chip formation may be expected to lead to a very fine microstructure, reinforced by a fine dispersion of precipitates that could result in chips endowed with significantly improved mechanical properties. For chips produced from Al 6061 in the solutiontreated state, the absence of precipitates in the matrix is expected to lead to a microstructure that is strengthened primarily through grain refinement. The difference in the microstructure between the chips cut from solution-treated 6061 and peak-aged 6061 can, therefore, be used to discern the role played by a fine dispersion of second-phase precipitates in the grain refinement and subsequent strengthening by the deformation. Furthermore, the precipitation kinetics of chips cut from solution-treated 6061 is expected to be significantly accelerated vis-a´-vis the undeformed bulk, as a result of the large defect concentration arising from the deformation. Appropriate thermal processing routes for consolidation can therefore be developed to exploit this kinetics and produce a fine dispersion of precipitates in the deformed matrix while retaining the microstructure resulting from the large strain deformation; this

would result in the final material having considerably enhanced mechanical properties. In the over-aged state, precipitation of solutes from the solid solution is nearly complete and so is the conversion of the fine b00 to coarse B 0 [12]. This coarse dispersion is not likely to be as effective at strengthening the microstructure as the b00 in the peak-aged temper. However, compared to the b00 phase, at the annealing temperatures that will be considered here, B 0 is not likely to coarsen rapidly. Therefore, the microstructure evolution during annealing of chips produced from the over-aged temper should purely be a function of the ability of these coarse precipitates to pin and stabilize the microstructure. 1.1. Machining Large strain deformation results from chip formation in its simplest manifestation, i.e. plane-strain (2-D) machining (Fig. 1). Plane-strain machining is characterized by a sharp, wedge-shaped tool that removes a preset depth of material (ao) by moving in a direction perpendicular to its cutting edge [14]. Chip formation occurs by concentrated shear within a narrow deformation zone, often idealized by a ‘‘shear plane’’ [14–16]. Most of the grain refinement associated with the formation of the UFG chips has been attributed to the large shear strains imposed in this deformation zone [10]. The geometry of the deformation zone and the associated shear strain are determined by the shear plane angle (/) and the rake angle (a). In Fig. 1, the rake angle (a) is taken to be positive when measured clockwise with respect to the normal to the workpiece surface and negative when measured counter-clockwise. Fig. 1 illustrates the case of a positive rake angle tool. The shear strain (c) imposed in the chip during chip formation is given by cos a c¼ ; ð1Þ sin / cosð/ aÞ

Chip Tool

Deformation zone

ac ao

Vo

Bulk Shear plane

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


where / is calculated from a measurement of ao and ac (Fig. 1), as tan / ¼

ao ac

1

cos a ao ac

sin a

ð2Þ

.

2. Experimental procedure The characteristics of chips produced by large strain deformation of Al 6061 alloy in three distinct tempers, viz. peak-aged Al 6061 (T6 temper), over-aged 6061 and solution-treated-6061, were studied through adaptation of 2-D machining as a large strain deformation process. Table 1 gives the experimental conditions. First, Al 6061-T6 plates were machined on a custombuilt plane-strain machining setup with high speed steel tools of different rake angles. A sufficiently low cutting velocity of 10 mm/s was employed to minimize strainrate effects and temperature rise during chip formation. In situ measurements of the deformation zone temperature made using infra-red thermography confirmed that the temperature rise during chip formation was negligible. The excellent machinability of 6061-T6 facilitated the imposition of a wide range of strains in the chip by 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 annealing behavior of the chips was characterized by heating the chips to various temperatures in a furnace (in air), followed by determination of the Vickers micro-hardness (50 g load) and microstructure characteristics. Annealing was done at 175 °C for different lengths of time to study the thermal stability at the characteristic ageing temperature for 6061. Annealing experiments were also done at 210 °C on the chips produced with different values of strain to determine any possible effect of deformation on the over-ageing characteristics. The microstructure of the chips, as-machined and after annealing, was characterized using a JEOL

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2000FX transmission electron microscope (TEM) operating at 200 kV. The chips tended to curl considerably and those produced with the negative rake angle tools tended to fragment. Hence, wedge polishing followed by ion milling was used, rather than electrolytic thinning, to produce most of the TEM chip specimens. In the wedge polishing, the specimens were first mechanically thinned by abrasive polishing to form shallow wedges. Each wedge was then mounted on a copper slot grid and ion milled for a very short duration of time (5 to 10 min) in a Gatan Model 600 dual ion mill to create an electron transparent specimen. Care was taken to avoid any prolonged heating of the specimens during the mechanical thinning and ion milling in order to ensure the integrity of the microstructure of the as-machined chip. Whenever continuous and relatively flat chips were produced, an electrolytic thinning technique was used for the TEM specimen preparation. In this case, the chips were first ground to a thickness of 100 lm using an abrasive grinding wheel. Three millimeter diameter disks were then punched out of the ground chip samples. The disk specimens were made electron transparent by electrolytic jet-thinning (Struers Tenupol-5) using a solution of 25% HNO3–75% CH3OH at 10 °C, 9.5 V. Second, chips were cut from over-aged Al 6061. For this purpose, peak-aged Al 6061-T6 plates were first subjected to a prolonged heat treatment (over-ageing) at 210 °C for 15 h. The choice of this temperature–time condition was guided by a prior TEM study of the precipitation sequence in Al 6061 [12]. Vickers hardness and microstructure of the chips cut from these plates were characterized. The annealing behavior of the chips was analyzed by studying the microstructure and hardness after heat treatment at 175 °C for different lengths of time. The choice of 175 °C was determined by the fact that the precipitates in the chips are likely to be relatively inert at this temperature; this is a consequence of the bulk plates being initially subjected to a prolonged heat treatment at 210 °C. Therefore, the annealing behavior of the chips is likely to be reflective of the

Table 1 Experimental conditions Bulk material characteristics

Machining conditions

Shear strain (c)

Chip annealing conditionsa

Al 6061-T6 (peak-aged) Hardness: 110 kg/mm2 Grain size 75 lm

Vc = 10 mm/s, ao = 150 lm Rake angle (a): +20° to 40°

See Fig. 2

175 °C for 1–15 h 210 °C for 1–10 h

Over-aged 6061 (obtained by annealing 6061-T6 at 210 °C for 15 h) Hardness: 95 kg/mm2 Vc = 10 mm/s, ao = 150 lm Grain size 75 lm Rake angle (a): +20° and 5°

a = +20° ! c = 3.2 a = 5° ! c = 4

175 °C for 1–15 h

Solution-treated 6061 (obtained by annealing 6061-T6 at 550 for 10 h) Vc = 10 mm/s, ao = 150 lm Hardness: 72 kg/mm2 Grain size 30 lm Rake angle (a): +20° and 5°

a = +20° ! c = 2.7 a = 5° ! c = 4

175 °C for 1–21 h 150 °C for 1–10 h

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

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annealing behavior of an UFG matrix that is pinned by coarse precipitates to the exclusion of any other softening phenomena. Lastly, to study the ageing behavior of Al 6061 after large strain deformation, plates of Al 6061 were solution-treated by heating them at 550 °C for 10 h and then quenching them in water at room temperature. Chips were then produced and characterized immediately prior to any natural ageing. The ageing characteristics of these chips were studied by heat treating them at 150 and 175 °C for different lengths of time. The temperature of 175 °C was chosen since it is the typical temperature used to age-harden solution-treated 6061 [17]. The heat treatment at 150 °C (a temperature usually too low to effectively age 6061) was based on prior ECAP work that indicated the sufficiency of lower-than-usual ageing temperatures to effectively harden severely-deformed, solution-treated materials [7]. Table 1 summarizes the experimental conditions including bulk material characteristics, machining parameters and annealing conditions for all of the chips.

3. Experimental results The utility of machining for introducing a range of shear strains in a single pass is evident from Fig. 2, which shows the variation of chip shear strain with tool rake angle for Al 6061-T6. The shear strain varies between 1.7 with a +20° rake tool and 5.2 for a 20° rake tool. When machining was done with tools of rake angle smaller than 20° (i.e. 30° and 40°), the cutting forces were significant (indicative of larger strains) and there was some flow of material in a direction normal

to the plane of machining (side-flow). Under such conditions, the deformation field is no longer one of plane strain and the applicability of Eqs. (1) and (2) is questionable. Hence, the results for the 30° and 40° rake angle chips are not reported here. The shear strains in the chips cut from the over-aged 6061 and solution-treated 6061 were also estimated using Eqs. (1) and (2). These are given in Table 1 for the +20° and 5° rake angle tools. Significant side-flow of material precluded analysis of strains in chips cut with tools of rake angles smaller than 5° for the over-aged and solution-treated 6061. 3.1. Peak-aged Al 6061 (6061-T6) Fig. 3 shows the variation of the measured Vickers hardness (Hv) with shear strain (c) for the as-machined Al 6061-T6 chips. The broken trend-line represents a power-law fit of the chip hardness (strength) data. The Vickers 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 tracking the variation in hardness values following their annealing at 175 and 210 °C for different lengths of time. Chips with c = 3.2 and c = 5.2 were chosen for this analysis. Fig. 4 illustrates the evolution of the Vickers hardness with annealing time at 175 °C. After annealing for 1 h, the hardness value is seen to decrease rapidly from a value of about 150 to about 135 kg/mm2. It is interesting to note that the annealed chip hardness value of 135 kg/mm2 is about the same as that obtained after warm ECAP of 6061-T6. Following this rapid initial decrease, the hardness value decreases at a much slower

10

S hear strain (γ)

8

6

4

2

0 -30

-20

-10

0

10

20

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

Fig. 3. Variation of Vickers hardness (Hv) with shear strain (c) for chips 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 at 175 °C for chips cut from peak-aged 6061.

rate during further annealing. After annealing for 10 h or more at 175 °C, the more highly strained chip is slightly softer than that produced with a smaller strain (Fig. 4). This hardness difference, though very small, is still statistically significant. This is indicative of a possible effect of cold work on the over-ageing characteristics. Fig. 5 shows the Vickers hardness of the chips produced with different levels of strain as a function of annealing time at 210 °C. For reference, an undeformed bulk sample was also annealed and its hardness recorded. The rate of decrease in hardness of the more severely deformed chip (c = 5.2) is considerably more rapid than that of the bulk and somewhat greater than

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that of the chip produced with c = 3.2. It is also interesting to note the one-to-one relationship between deformation strain and the resulting hardness following annealing at 210 °C. Material deformed to the smallest strain, viz. the undeformed bulk Al 6061-T6 (c = 0), softens the slowest and remains the hardest compared to the chips following the annealing. Material deformed to the largest strain value (c = 5.2) anneals the fastest and also reaches the lowest value of hardness after annealing. Fig. 6 shows TEM micrographs illustrating the effect of the strain on the resulting chip microstructure. While the grain morphologies of chips produced with c = 3.2 and c = 5.2 are similarly equi-axed and typically finer than 100 nm, the less strained c = 1.7 chips are composed of elongated and coarser grains of width 150 nm. The diffraction pattern of the c = 1.7 chips is characteristic of a matrix made up of sub-grains of low misorientation angles (Fig. 6(a)), while the nearly uniform rings around the central bright spot in the case of the c = 3.2 and c = 5.2 chips are characteristic of large values of misorientations between the sub-grains (Figs. 6(b) and (c)). The chips annealed at 175 and 210 °C were examined in the TEM to track coarsening effects. After 1 h of annealing at 175 °C, the grain size was found to increase substantially (Figs. 7(a) and (c)) from that of the asmachined chip, although the resulting grains are still sub-micron in size. The same effect was encountered after annealing at 210 °C for 1 h as well (Fig. 8). Fig. 7 also shows that prolonged annealing for more than 1 h at 175 °C does not lead to significant continued coarsening 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 of these precipitates appears to be much larger than those observed in the peak-aged material [12]. These precipitates at the grain boundaries are likely to impede coarsening of the microstructure during extended annealing by pinning the grain boundaries. 3.2. Over-aged Al 6061

Fig. 5. Evolution of Vickers hardness (Hv) with annealing time at 210 °C for chips cut from peak-aged 6061.

Chips were cut from a bulk over-aged 6061 sample with a Vickers hardness of 95 kg/mm2. The hardness of the c = 3.2 chips was 122 kg/mm2 while that of the c = 4 chips was 134 kg/mm2 (Fig. 9). For the same shear strain value of 3.2, chips produced from peak-aged 6061 (6061-T6) were 1.38 times as hard as the bulk while the chips produced from the over-aged 6061 were 1.27 times as hard as the corresponding bulk material. Fig. 9 illustrates the variation of hardness with annealing time at 175 °C. During prolonged annealing, the hardness of the over-aged bulk material is relatively unaffected. The chips however show a steady decrease in hardness that appears to be independent of the extent of

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Fig. 6. Bright field TEM images of chips cut from peak-aged 6061 with different induced strains. Insets are the corresponding selected area diffraction patterns. (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 coarse spherical 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.


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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. Grain size 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 at 175 °C for chips cut from over-aged 6061. For reference, the hardness evolution of the bulk material is also shown.

prior deformation. Unlike the case of the chips cut from the peak-aged 6061, the more highly strained chips from the over-aged 6061 do not anneal at an accelerated rate to a lower value of hardness. Fig. 10 illustrates the microstructure of as-machined chips with different levels of strain. The typical grain sizes are 160 nm, nearly double the grain size of chips cut from peak-aged 6061-T6 for comparable values of strain. The insets of the diffraction patterns in Fig. 10 show relatively uniform intensities of the diffraction rings around the central spot and particularly so with the more highly strained chip (Fig. 10(b)). Fig. 11 shows the microstructures of the chips after 1 h anneal at 175 °C. The resistance to significant grain coarsening is evident from this figure in contrast to the annealing behavior of chips machined from the peakaged 6061 (Fig. 7). The grain-size after the annealing (160 nm) is about the same as in the as-machined state although it is apparent from Fig. 9 that there is some associated decrease in the hardness following the annealing.

Fig. 10. Bright field TEM images of chips cut from over-aged 6061 with different levels of strain. Insets are the corresponding selected area diffraction patterns. Typical grain size 160 nm. (a) c = 3.2. (b) c = 4.0.

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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 = 3.2 chip annealed for 1 h at 175 °C. (b) c = 4 chip annealed for 1 h at 175 °C.

3.3. Solution-treated Al 6061 The hardness of the chips created with 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 solution-treated 6061 had a hardness of 72 kg/mm2 (Table 1). The chips were heat treated at 175 and 150 °C for various lengths of time to study their ageing characteristics. Figs. 12 and 13 show the variation of Vickers hardness with ageing time at 175 and 150 °C, respectively. The ageing characteristics of the chips at 175 °C appear to be independent of the prior strain (Fig. 12); both the highly strained and less highly strained chips are seen to age very rapidly compared to the undeformed bulk. Furthermore, the chips attain a peak-hardness value of 132 kg/mm2 after 1 h of annealing at 175 °C. Prolonged annealing at 175 °C caused a steady decline in the Fig. 13. Evolution of Vickers hardness (Hv) with ageing time at 150 °C for chips cut from solution-treated 6061.

Fig. 12. Evolution of Vickers hardness (Hv) with ageing time at 175 °C for 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 of the chips. The ageing characteristics of the chips were also studied at 150 °C, to explore the possibility of realizing higher values of hardness (strength) by lowering the ageing temperature and inhibiting any coarsening phenomena that may reduce the strength. Fig. 13 shows the variation of hardness with ageing time 150 °C. Not only is a higher hardness value (140 kg/mm2) achieved following the ageing for 1 h, but this value remains relatively unchanged even after 5 h of ageing. However, after 10 h of ageing at 150 °C, the hardness declined from 140 kg/mm2 to 132 kg/mm2, which is the same as that recorded on chip samples after 1 h of ageing at 175 °C. Fig. 14(a) shows a TEM micrograph of a c = 2.7 chip machined from solution-treated 6061. The grains are elongated with a typical width of 150–200 nm; this


microstructure is characteristic of a relatively low level of SPD compared to say the strain (c = 5.2) achieved with a 20° rake tool in peak-aged 6061. Figs. 14(b) and (c) show evolution of the microstructure during the ageing at 150 °C. A side-effect of the ageing process is grain growth that coarsens the UFG microstructure and contributes to a reduction in hardness during the prolonged ageing. At 175 °C, as seen in Fig. 15, there is significant coarsening of the microstructure (grain size 400 nm). But, the hardness value (132 kg/mm2) is still 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 strengthened by a fine dispersion of precipitates in several of the grains in Fig. 14(c), appears to be very similar to that shown in Fig. 15. The diffraction patterns shown as insets in Figs. 14 and 15 indicate qualitatively an evolution in the misorientation between the grains during the ageing process. The diffraction pattern of the as-machined chip is characteristic of a microstructure made up of grains that are considerably misoriented with respect to one another. After the recovery and grain-growth that occur in tandem with the precipitation during ageing, the diffraction pattern is characteristic of a microstructure composed of low-angle boundaries.

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

4. Discussion Strengthening in UFG multi-phase, Al alloys can occur due to a combination of solid-solution strengthening, precipitation strengthening and grain size strengthening. The same phenomena are likely to also affect the formation of UFG microstructures in precipitationstrengthened aluminum alloys during large strain deformation, by modifying the micro-mechanics of plastic

Fig. 14. Bright field TEM images of as-machined chips cut from solution-treated 6061 and after ageing at 150 °C for different lengths of time. (a) c = 2.7 chip cut from solution treated 6061. Grain size 110 nm. (b) c = 2.7 chip annealed at 150 °C for 5 h. Grain size 160 nm. (c) c = 2.7 chip annealed at 150 °C for 10 h. Arrows indicate a fine dispersion of precipitates within a grain. Grain size 270 nm.

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flow and grain refinement. This study has explored the interplay of the role of precipitate size and the singlepass, large-strain deformation as well as the effect of the large strain deformation on the ageing characteristics of the UFG microstructures. 4.1. Large strain deformation of precipitationstrengthened 6061 4.1.1. Grain refinement Large strain deformation by machining is a simple route for studying the effect of precipitates on the scale and stability of UFG Al 6061 microstructures. A range of shear strains and a variety of resulting microstructures can be produced by chip formation through an appropriate choice of the tool rake angle. However, for the same rake angle,two different materials or different tempers of the same alloy can lead to different values of shear strain in the chip. A 20° rake angle tool is seen to produce a shear strain of 1.7 in peak-aged 6061 and a shear strain of 3.2 in over-aged 6061. This is not surprising considering that in machining the geometry of the deformation field is not defined a priori. This does not imply that strains cannot be introduced in a controllable manner. From Fig. 2 it is possible to determine experimentally, the one-to-one correspondence between the tool rake angle and the shear strain in the chip for a specific material state; this relation can then be used to impose a controlled level of strain in the chip. Fig. 3 shows that the chip hardness is much greater than that of the bulk in peak-aged 6061 even at the lower strains, while the incremental increase in the chip hardness at larger strains is diminishing. A switch-over of the grain morphology from elongated structures at smaller chip strains (Fig. 6(a)) to fine, equi-axed grains occurs at larger strains (Figs. 6(b) and (c)). The grain size in the c = 1.7 chips measured along the smallest dimension is 150 nm while the equi-axed grains in the 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 observed in the chips cut from the over-aged 6061 since large values of shear strain (c = 3.2) and an equi-axed grain 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 that of the bulk is found to depend on the characteristics of the bulk material. The chip-to-bulk hardness ratio in the case of the peak-aged 6061 is 1.38 and in the over-aged material is 1.27, even though the initial grain sizes of the bulk material in both cases are very similar(75 lm). This difference can be attributed to more effective trapping of dislocations and associated refinement of microstructure brought about by a fine dispersion of precipitates in the case of the peak-aged material (typically b00 [12,13]) as compared to the coarser precipitates

in the over-aged material (typically B 0 [12]). A finer dispersion provides a larger number of sites at which sub-grain boundaries are pinned thus preventing any significant coarsening during dynamic recovery that can occur in the course of the large deformation. Furthermore, the greater concentration of solutes in the peak-aged state compared to the over-aged state can lead to a greater resistance to dislocation motion and contribute to larger values of hardness (strength). It has been found in rolling of a similar Al–Mg–Si alloy that the precipitates do indeed break up leading to a finer dispersion in the aluminum matrix [18]. Similar instances of precipitates breaking down during deformation are also reported in SPD of Al 5083 [19], Al 7050 [20] and an Al–Cu alloy [21]. We anticipate a similar precipitate breaking phenomenon here that results in fragmented precipitates. These contribute to the fine microstructure and associated strengthening. The hardness values of the chips cut from the peakaged temper are generally greater and the scale of the microstructure observed here is generally finer than that encountered in warm-ECAP of peak-aged 6061 [3]. This could be a result of recovery and loss of dislocations in the ECAP, and thus ineffective grain refinement due to the heating associated with the deformation in the warm ECAP process. The room temperature deformation effected here during machining involves no such softening aspects and thus leads to larger values of chip hardness. 4.1.2. Thermal stability The thermal stability of the UFG 6061 microstructures is of particular interest because of the possible utilization of the machining chips in the fabrication of bulk nanocrystalline materials through powder consolidation techniques. In the case of chips cut from the peak-aged material, the hardness value decreases rapidly in the initial stages of annealing. Prolonged annealing leads to a slow decline in the hardness, as seen in Figs. 4 and 5. The initial decrease in hardness coincides with the rapid increase in the grain size as is evident from Figs. 6–8. Rapid grain growth from what appears to be a relatively unstable configuration of precipitates and grain boundaries to a more stable configuration, rather than a mere ‘‘clean-up’’ of the dislocations from the interior of the grains, appears to be the operative phenomenon here. A study of the annealing behavior beyond 1 h did not indicate any continued increase in grain size. This is consistent with the notion that a relatively stable microstructure forms after about 1 h of annealing at 175 °C, one that is stabilized by the precipitates which pin the grain boundaries preventing grain growth. The large strain deformation is usually associated with fracture of the precipitates and this leads to the formation of considerably finer precipitates with a large surface energy component. Furthermore, it has been suggested in


the case of severely deformed Al–Mg–Si alloys that the sheared, fine precipitates, which have lost their coherency with the matrix as a result of deformation, would rapidly transform to the equilibrium b phase [18]. The resulting stable b phase is characterized by relatively coarse spherical precipitates, similar to those observed here in Fig. 7. The stable b phase that would normally require prolonged heat treatment to form at 175 or 210 °C in the undeformed bulk, has formed quite rapidly in the machining chips due to the instability introduced in the b00 precipitates as a result of large defect concentration and possible breaking up of coherent precipitates. The hardness vs. annealing time data in Figs. 4 and 5 show that the flow stress continues to decrease with annealing time at a rate determined by the extent of cold work during large strain deformation. The c = 5.2 chips are somewhat softer following annealing than the c = 3.2 chips. The difference between the coarsening rates in these two material states might be a result of the differing sizes of fragmented precipitates and differing defect concentrations left behind by the deformation. The greater the defect concentration in terms of vacancies produced by cold work, the greater is the diffusivity of the solute atoms through the matrix and therefore, the greater the rate of diffusion-controlled coarsening [22]. Also, on deforming to larger values of strain, the precipitates are likely to break up further to form an increasingly finer distribution. Such a fine dispersion could be expected to coarsen at a very rapid rate during heat treatment, due to the increased driving force provided by the larger interfacial energy [18]. For these reasons, it is reasonable to expect that precipitates in chips deformed to larger values of strain coarsen at a faster rate, and, therefore, these chips, possess lower strength following annealing. From Figs. 10 and 11 it is apparent that chips cut from the over-aged state are relatively immune to grain growth following annealing at 175 °C. Since the bulk material was in the over-aged state, further coarsening of the precipitates is not likely to be the key aspect influencing the softening of the material. Rather, as seen more clearly in Fig. 11(a), a regular recovery, together with coarsening of some of the broken precipitates, followed by gradual but highly restrained grain growth, appears to be the mechanism that determines the softening of the chips created from the over-aged 6061. 4.2. Ageing characteristics of the chips The properties of the chips machined from the solution-treated state suggest new and exciting possibilities for consolidation of these chips by thermo-mechanical processing. Since some heating is usually involved in any consolidation and if that heating can be associated with a hardness increase rather than softening, then this

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phenomenon can be exploited during the consolidation to produce bulk forms with enhanced strength. That is, chips cut from the solution-treated bulk may be thermo-mechanically processed in such a way that they gain hardness during the consolidation, thus leading to significantly stronger bulk forms. From Figs. 12 and 13, it is apparent that chips machined from the solutiontreated state, unlike those from the peak-aged state or the over-aged state, gain in strength following heat treatment. This clearly implies that the rate of hardening due to precipitation exceeds any softening due to the coarsening of the microstructure from grain growth. Precipitation and grain growth are both diffusion driven, but mutually counter-acting phenomena, as far as the flow stress is concerned; so the relative rate of each process at a given temperature would determine the final strength of the material. By comparing Figs. 14(b) and 15, it is apparent that higher temperatures lead to a greater coarsening of the grains and lower hardness values (see also Figs. 12 and 13). The deformed microstructure significantly enhances the mobility of the solute atoms in the matrix. This when combined with the very large number of nucleation sites made available by the dislocations and the grain boundaries, may be expected to lead to a rapid formation of a fine dispersion of precipitates. The accelerated formation of a fine dispersion of precipitates can occur even at a relatively low temperature. The precipitation hardening of the matrix may then be expected to prevail over the softening effects due to grain growth in a low-temperature annealing process. It is also interesting to note that the hardness and the microstructure resulting from annealing of the solution-treated chips at 150 °C for 10 h are nearly identical to those realized following annealing at 175 °C for 1 h. Both are characterized by a fine dispersion of precipitates in the interior of a grain and a relatively large grain size of about 500 nm. These results are similar to the ageing characteristics of solution-treated 6061 [7] and an Al–Sc alloy [23] processed by ECAP. The same hardness value of 132 kg/mm2 is obtained in the case of the annealed peak-aged or the as-machined over-aged chips, which are both typically composed of 200 nm grains. Hence, significantly improved mechanical properties in the case of precipitation-strengthened materials can either be obtained through simple grain refinement in one set of tempers or by producing an especially fine dispersion of precipitates that accomplishes the same task in a different temper. These results reaffirm the case for machining as a single-pass process for studying large strain deformation of materials. From a practical standpoint, they suggest a parameter window for thermal processing of the chips. In the case of the chips machined from the peak-aged and the over-aged materials, softening during thermal treatment is inevitable. However, by keeping the processing temperature around 175 °C it is possible to

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retain some of the increase in strength resulting from the large strain deformation. In the case of the chips cut from the solution-treated 6061, a significant increase in strength can occur due to the thermal treatment. This aspect combined with the retention of a UFG matrix is likely to be of advantage during thermally-assisted consolidation of these chips into bulk forms.

5. Conclusions We have demonstrated the effectiveness of planestrain machining as a technique for studying large strain deformation of aluminum alloys in various tempers. A variety of strains, microstructures and mechanical properties are shown to be obtained by appropriate choice of the tool rake angle. Furthermore, the large strains are imposed in a single pass of deformation. A very fine microstructure composed of grains smaller than 100 nm is found to result from chip formation in peak-aged 6061. The scale of the microstructure is generally finer, and the associated hardness value for the chips somewhat larger, than those reported for ECAP at elevated temperatures, implying a strong effect of deformation temperature on the resulting microstructure. Chips cut from the peak-aged Al 6061, with larger values of strain, are less stable during annealing than their less-strained counterparts. This difference in stability may be rationalized by considering the effect of the deformation on the precipitates, microstructure and transport properties. Coarsening is found to be a two step process. Initial rapid grain growth is seen to give way to a more stable microstructure that gradually softens during prolonged annealing. The b00 phase that is responsible for the peak-aged strength of Al 6061, following severe deformation is expected to rapidly transform to the relatively coarse equilibrium b phase. This stable b phase, however, is still capable of effectively pinning the grain boundaries thus preventing further coarsening. Deformation of over-aged 6061 by machining results in relatively coarser grains and less hardening compared to similarly deformed peak-aged material. This can be attributed to less-efficient dislocation trapping and stabilizing by the coarser B 0 phase that is dispersed in the over-aged bulk matrix. Annealing of this UFG microstructure results in gradual softening with only minor changes in the microstructure; this is to be contrasted with the annealing behavior of chips cut from the peak-aged material. Chips from the solution-treated material, on the other hand, first harden following heat treatment and then soften after prolonged heat treatment. Accelerated formation of very fine precipitates causes the initial rapid increase in strength in this material, and this effect is then counter-acted by grain growth and coarsening of the fine precipitates.

The observations suggest thermal processing routes for consolidation of the chips into bulk forms while retaining the UFG microstructure with enhanced mechanical properties.

Acknowledgements We thank the Department of Energy (Grant 4000031768 via UT-Batelle), Ford Motor Company, Oak Ridge National Laboratory (ORNL) and the State of IndianaÕs 21st Century Research and Technology Fund for support of this work. Additional thanks are also due to Drs. Ray Johnson (ORNL) and Andrew Sherman (Ford) for their encouragement of the studies.

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