Turkish J. Eng. Env. Sci.31 (2007) , 53 – 60.c© TUBITAK

Influence of Heat Treatment on the Mechanical Properties ofAA6066 Alloy

Evren TAN and Bilgehan OGELMiddle East Technical University, Department of Metallurgical and Materials Engineering,

Ankara-TURKEYe-mail: etan@metu.edu.tr

Received 25.08.2006

Abstract

The microstructural and mechanical characterization of heat treatable 6xxx (Al-Mg-Si-Cu based) wroughtaluminum alloys was studied. The aim of this work was to produce a fine grained, high strength 6xxx seriesaluminum alloy by adjusting the processing conditions, namely deformation, solutionizing and aging. Theireffects were investigated in terms of microstructure using SEM analysis and mechanical properties by tensiletests and hardness measurements. The initial characterizations showed that Mg2Si and (Fe,Mn,Cu)3SiAl12were the primary particles observed in the α-Al matrix. Nearly 140HB hardness was obtained with solution-izing at 530 ◦C and aging at 175 ◦C for 8 h, which was the optimum treatment for obtaining peak hardness.When shaping (deformation) was concerned, 10% swaging before solutionizing yielded lower strength andhardness as compared to the 40% swaging due to lower strains finalizing partial recrystallization in the 10%swaged specimens.

Key words: Al alloy 6xxx series, Deformation, Grain size, Microstructural characterization.

Introduction

Aluminum alloys have been the material of choice foraircraft construction since the 1930s. The aerospaceindustry relies heavily on 2xxx and 7xxx alloys, while6xxx aluminum alloys are of particular interest nowa-days. According to Troeger (2000), 6xxx alloyshave numerous benefits including medium strength,formability, weldability, corrosion resistance, and lowcost. He states that 6xxx can be used in a variety ofapplications including aircraft fuselage skins and au-tomobile body panels and bumpers, instead of moreexpensive 2xxx and 7xxx alloys, after appropriateheat treatments. Hence, microstructural character-ization of the alloy and processing procedure is im-portant for that approach.

Characterization of Al-6xxx alloys has been thesubject of many studies. It is known that the maincomponents of heat treatable 6xxx series Al alloyare Mg and Si, and 6xxx derives its strength from

the precipitation hardening phase, Mg2Si. The vol-ume fraction of Mg2Si is affected primarily throughthe level of Mg within the alloy, but the Si content isalso important. Hirth et al. (2001) concluded thatincreasing Si in 6xxx type alloys increases strengthin the T4 and T6 tempers. Another study on 6xxxcarried out by Matsuda et al. (2002) showed thatthe addition of copper to Al-Mg-Si alloys not onlychanges the precipitation sequence but also enhanceshardness and refines microstructure by segregatingto the Q’ \α – Al interface. In addition to Mg-Si-Cu, manganese and chromium are also used as alloy-ing elements. Dorward and Bouvier (1998) explainedthe beneficial effects of manganese and chromium,namely that these alloying additions inhibit the pre-cipitation of magnesium and silicon on grain bound-aries, thereby reducing intergranular fracture ten-dencies. The beneficial effect of manganese (or otherincoherent dispersoid-forming elements) on tough-ness is homogenization of deformation, leading to a

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reduction in intergranular fracture. In addition, theynoted that a refined grain size, as influenced by man-ganese (a grain structure control agent), was anotherpositive effect.

According to the alloy composition and agingprocedure, different precipitates can be observedin the microstructure, which will affect the finalmechanical properties. Chakrabarti and Laughlin(2004) tried to place some of these precipitates ona phase diagram. By using their line diagram, thestable phases at RT can be observed.

In order to obtain optimum mechanical proper-ties, a suitable production procedure should be se-lected. In addition to aging time and temperature,the presence of deformation, its temperature andplace in the whole production procedure should becarefully checked. Zhen et al. (1997) stress the im-portance of time passed between solution heat treat-ment and artificial aging, which shows that any timespent between quenching and aging leads to naturalaging of the specimen and hence lowers hardness inthe end. In another study, Sun et al. (1999) de-formed samples between quenching and aging andthey concluded that this method has positive effectson hardness and strength. A different method wasused by Cai et al. (2004). They tried a dynamic ag-ing procedure (integrated process combining thermo-mechanical processing and aging). In summary, dy-namic aging was superior to conventional aging interms of both mechanical properties and time topeak strength.

All studies on 6xxx above were based on Al-6061and 6063 type alloys. In our work, the properties ofa 6xxx alloy containing Si, Mg, Cu and Mn in theorder of 1wt% conforming to AA6066 aluminum al-loy were investigated. Al-6063 has balanced Mg-Si toform stoichiometric Mg2Si and Al-6061 has a Cu in-gredient besides balanced Mg-Si; on the other hand,Al-6066 has both Cu and excess Si. As stated pre-viously, different properties were obtained with var-ious amounts of alloying elements. Hence, the aimof the present study was to optimize the heat treat-

ment and to investigate the effect of initial deforma-tion (shaping) process on the mechanical propertiesof Al-6066 type alloy.

Experimental Procedure

This research was conducted using wrought 6xxx se-ries aluminum rods with the dimensions of Ø = 46mm and l = 200 mm. The as-received materials withthe chemical composition stated in Table 1 were inthe annealed temper, which conforms to AA6066 al-loy.

The microstructural investigations and the me-chanical testing were carried out as an initial char-acterization step for the 3 tempers: annealed,solutionized-water quenched (WQ) and aged.

The effect of temperature and time during so-lutionizing and aging were examined as the secondstep. For these experiments, the rods were cutinto 40 equal slices. Throughout the study same-dimensioned test specimens were used. In order toinvestigate the effect of solutionizing temperature,the samples were solutionized for 95 min and for12 h at 4 different temperatures (515, 530, 540 and550 ◦C). A 95 min duration was calculated accord-ing to the dimensions of specimens suggested by theAluminum Association. Twelve hours was selectedarbitrarily to investigate the effect of soaking timeon hardness. All solutionizing treatments were per-formed in a muffle furnace operating with ±1 ◦C.Temperature of the furnace was controlled by an ex-tra secondary thermocouple. Neither air nor gas cir-culation was used in the furnace.

After solutionizing, all samples were quenched inwater at room temperature. Ice was used to keepthe temperature at about 20 ± 2 ◦C. During quench-ing, the medium was stirred to achieve temperaturehomogeneity. It must be noted that, after solution-izing treatment, the samples were always kept in afreezer, which was kept at about –18 ◦C to eliminatethe detrimental effect of natural aging.

Table 1. The nominal chemical composition of the alloy used.

Ni, Cr,Si % Mg % Cu % Mn % Fe % Ti % Zn % Pb, Sn, Al %

Ca, Sr1.21 0.92 0.82 0.65 0.26 0.013 0.010 ≈ 0 balance

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For the ideal solutionizing temperature, 24 differ-ent sets of aging trials were carried out. The selectedaging temperatures were 150, 175 and 200 ◦C. Thespecimens were held at these temperatures for 2 to16 h with a 2-h interval. The aging treatment wascarried in a Julabo heating circulator oil bath, whichoperates with 0.1 ◦C accuracy.

As a final step, the effect of the swaging processprior to heat treatment was investigated. Rods weremachined with 4 different diameters such that ap-proximately 10%, 20%, 30% and 40% deformation(reduction in area) could be established. The de-formation process was done with the use of a Fennrotary die swager at room temperature.

The detailed processing of these specimens is tab-ulated in Table 2.

For microstructural analysis, specimens wereground with emery papers from 600 to 1200, andpolished with 1 µm diamond paste. As an etchant,a mixture of HF, H2SO4 and H2O in the compo-sition 1:2:17 was used. For characterization, opti-cal microscopy and a scanning electron microscope(JEOL JSM 6400r© equipped with Noran r© EDSsystem) were used. Image analysis was carried out

with Clemex Vision Professional software.For hardness measurements, Brinell hardness

numbers were obtained with a 2.5 mm ball indenterunder 613 N load. For tension testing, samples wereprepared according to the aluminum tensile testingstandard of ASTM (B557M-02a).

Results and Discussion

Microstructural Features and MechanicalProperties

As seen in Figure 1, the backscatter SEM imagesand optic micrographs on as-polished surfaces re-vealed 2 types of particles in the α-Al matrix: largeblack particles and tiny (5-10 µm) gray script-likefeatures. With the assistance of EDS analysis (Fig-ure 2), the features were possibly identified as Mg2Siand (Fe,Mn,Cu)3SiAl12.

Highly alloyed 6xxx had complex intermetallicsoriginating from cast ingots. Since iron was the om-nipresent impurity element and had a very low sol-ubility in aluminum, iron-rich phases could be seenin all aluminum alloys. The presence of manganese,chromium or copper leads to the formation of

Table 2. Processing schedule.

SOLN AGING# DEFN Temp Time Temp Time

% (◦C) (min) (◦C) (h)1

-

- - - -2 515

95, 720 175 83 5304 5405 5506 1507 - 530 95 175 2, 4, 6, 8, 10, 128 2009 10

- - - -10 2011 3012 4013 10

530 95 - -14 2015 3016 4017 10

530 95 175 4, 6, 8, 10, 1218 2019 3020 40

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Figure 1. (a) Optical micrograph and (b) backscatter SEM image of as-polished surface.

7000

6000

5000

4000

3000

2000

1000

0

4000

3000

2000

1000

00 1 2 3 4 5 6 7 8 9 10

keV

Full scale count: 6643 Full scale count: 4384

Mn CuFe

Sl

Al

Mn Fe Fe Cu

Si

Al

Mg

0 1 2 3 4 5 6 7 8 9 10keV

Element Weight AtomConc. % Conc. %

Al 86.58 91.91Si 2.52 2.57Mn 4.41 2.30Fe 4.72 2.42Cu 1.78 0.80

Element Weight AtomConc. % Conc. %

Mg 25.92 28.16Al 56.59 55.39Si 17.49 16.45

Figure 2. EDS analysis of microscopic features.

Table 3. Mechanical test results for O, T4 and T6 tempers and their comparison with the standard AA6066.

Experimental Results Literature ASM (1996)Brinell Yield UTS Brinell Yield UTS

Hardness (MPa) (MPa) Hardness (MPa) (MPa)6066 - O 59-60 168-181 220-223 43 82.7 1526066 - T4 100-103 300-335 442-445 90 207 3596066 - T6 128-140 441-461 461-478 120 359 393

(Fe,Mn,Cu)3SiAl12. The other phase, Mg2Si, wasthe main ingredient of 6xxx, which would readily dis-solve during solutionizing and contribute to the pre-cipitation hardening for the period of artificial aging.

The mechanical properties of the alloys at the3 tempers are summarized in Table 3. It is evidentthat the alloy used in the study fulfilled the minimumrequirements given by the Aluminum Association.

Ideal Heat Treatment Cycle

Solutionizing is the main step for precipitation hard-ening, and it must be carried out carefully in orderto avoid grain boundary melting due to overheating.The effect of solutionizing treatment is summarizedvia a column chart in Figure 3. The as-received spec-imen in the annealed temper having a hardness of57HB was treated with different temperatures andtimes before artificial aging at 175 ◦C for 8 h. Peak

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hardness was achieved after solutionizing at 530 ◦Cfor 12 h. The chart demonstrated that, as the tem-perature increased for 95 min solutionizing, the hard-ness increased from 131 to 144. Actually this wasforegone; as temperature increases more solid dis-solves and mixes in the matrix for super saturation,due to increased diffusion kinetics and higher solubil-ity limits at higher temperatures. Another output ofthe chart could be on the soaking times. Below 530◦C, there was a tendency towards increasing hard-ness with increasing soaking time, whereas the trendwas the opposite for temperatures above. Below 530◦C, with the increase in soaking time, a satisfactorydegree of solution of the undissolved or precipitatedsoluble phase constituents, forming a good homo-geneity of solid solution, could be established. After530 ◦C, the loss of hardness for long soaking timesmight be interrelated with the most probable graingrowth tendency. Although grain boundary meltingwas not observed, to be on the safe side 95 min du-ration at 530 ◦C could be chosen as the solutionizingtreatment for the alloy since the hardness values werecloser to each other.

The aging behavior of the alloy is summarized inFigure 4. It shows typical hardness curves of solu-tionized alloys after artificial aging at elevated tem-peratures. A peak hardness of 139 HB could be at-tained after 8 h of aging at 175 ◦C and then the hard-ness decreased slightly. The curve for 200 ◦C did notshow peak hardness but it could be predicted to bein the vicinity of a half an hour or so. The quick dropin the curve was an indication of the fast overaging

due to high diffusion rates at high temperatures. Onthe other hand, the effect of slow kinetics could bethe conclusion of the aging at 150 ◦C. The hardnessvalue reached after 16 h could be obtained by agingat 175 ◦C for 2 h only. Hence it can be concludedfrom Figure 4 that aging at 175 ◦C was the optimumfor the alloy, since it yielded the desired hardness val-ues of 130-140HB in 2-16 h.

Effect of Deformation

The 6xxx type of aluminum alloy is used for gen-eral extrusion purposes. Hence the effect of defor-mation behavior prior to heat treatment must beconsidered. Ultimate tensile strength (UTS) and0.2% offset yield strength of non-deformed and de-formed samples after peak aging were obtained fromthe stress-strain diagram and are summarized in Ta-ble 4. At first sight, deformed specimens had lowerstrength as compared to non-deformed ones. Theonly reason for that would be the larger grain sizesin the deformed samples as compared to the originalnon-deformed ones having a mean grain size of 18 ±4.77 µm. The effect of deformation on final grain sizeis shown in Figure 5. Another output of the mechan-ical test was the increase in strength with an increasein percent deformation. Deformation leads to energystorage for the period of lattice defect creation, i.e.dislocations. During solutionizing at temperatures of530 ◦C, the cold-deformed specimens led to recrys-tallization. It was known that the higher the strains,the lower the recrystallized grain size. Hence, after

62.5

Figure 3. Effect of solutionizing treatment on hardness prior to aging at 175 ◦C for 8 h.

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Figure 4. Comparison of aging behavior of solutionized alloy aged at temperatures of 150, 175 and 200 ◦C.

Figure 5. Optic micrographs after solutionizing and the correspondence grain size distributions; a, c) 10% deformedsamples; b, d) 40% deformed samples.

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solutionizing the 10% deformed sample with meangrain sizes in the order of 26.7 ± 7.02 µm had UTSof 376 MPa, where the 40% deformed sample with20.5 ± 4.72 µm average grain size had UTS of 418MPa.

Table 4. Tensile strength of AA6066 alloy solutionized at530 ◦C and peak-aged at 175 ◦C for 8 h afterdeformation.

Deformation Status UTS Yield StrengthNon-deformed 470 45010% RT 376 29520% RT 380 29830% RT 383 30040% RT 418 319

Figure 6. Effect of deformation on the hardness of alloy.

A similar trend could be obtained from hardnessvalues. The variations in hardness for non-deformed,

10% and 40% deformed specimens are presented inFigure 6. Right after deformation, the 40% deformedspecimen had 79HB due to higher dislocation densitythan the 10% deformed one with 73HB. After recrys-tallization during solutionizing and further artificialaging, a hardness trend similar to that in UTS wasobserved.

Conclusion

In this study, the effect of deformation on me-chanical properties of a 6xxx series aluminum al-loy was investigated. Following the determinationof the ideal conditions for solutionizing and agingprocesses, specimens were mechanically deformed byswaging at 4 different deformations and then heattreated. The primary conclusions obtained from thisstudy can be summarized as follows:

1. Two types of particles were observed inthe alloy: black Mg2Si and gray script-like(Fe,Mn,Cu)3SiAl12, which were equilibriumconstituents coming from the cast ingot.

2. The ideal solutionizing temperature was 530◦C. Below 530 ◦C, there was a tendencytowards increasing hardness with increasingsoaking time, whereas the trend was the op-posite for temperatures above.

3. Aging trials between 150 and 200 ◦C showedthat peak hardness values could be obtainedafter aging at 175 ◦C for 8 h. As comparedto the fast overaging at 200 ◦C and slow hard-ening at 150 ◦C, 175 ◦C was confirmed as theoptimum aging temperature for industrial us-age.

4. As deformation amount increased, the recrys-tallized grain size got smaller, enhancing thestrength and hardness.

Acknowledgment

Special thanks to Mechanical and Chemical Indus-trial Corporation Ammunition Group for their helpwith material provision.

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