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Materials Science and Engineering A 528 (2011) 8060– 8064

Contents lists available at ScienceDirect

Materials Science and Engineering A

journa l h o me pa ge: www.elsev ier .com/ locate /msea

nhanced mechanical properties in an Al–Cu–Mg–Ag alloy by duplex aging

ao Lia,b, Zhiyi Liua,b,∗, Song Baia,b, Xuanwei Zhoua,b, Heng Wanga,b, Sumin Zenga,b

Key Laboratory of Nonferrous Metal Materials Science and Engineering, Ministry of Education, Central South University, Changsha 410083, ChinaSchool of Materials Science and Engineering, Central South University, Changsha 410083, China

r t i c l e i n f o

rticle history:eceived 1 April 2011eceived in revised form 30 May 2011ccepted 25 July 2011

a b s t r a c t

A type of duplex aging heat treatment was developed to improve the mechanical properties at room tem-perature and elevated temperatures in a pre-strained Al–Cu–Mg–Ag alloy. In contrast to the conventionalT8 temper at 165 ◦C and 200 ◦C, the hardening response of the alloy to aging was increased by duplex agingtreatment, the ultimate tensile strength and yield strength of duplex aging temper were improved by

vailable online 30 July 2011

eywords:luminum alloysechanical characterization

approximately 3–7%, which was attributed to the fact that the recovery of dislocations occurred and theprecipitation of �′ phase was restrained effectively at high aging temperature, and more � precipitateswere formed during secondary aging.

© 2011 Elsevier B.V. All rights reserved.

gingrecipitation

. Introduction

Al–Cu–Mg–Ag alloys are promising materials for aerospacepplications due to their high strength and excellent thermal sta-ility, as well as creep resistance [1–3]. These superior propertiesre attributed to the formation of a fine and uniform dispersion ofexagonal-shaped plate-like � precipitates on the {1 1 1}� matrixlanes, which is promoted by Ag addition to the alloy with highu/Mg ratios. There is competitive precipitation between �′ and

precipitates in Al–Cu–Mg–Ag alloys, although the dominanthase is � phase. The precipitation sequences of the alloy cane represented as: SSS → GP zones → �′ ′ → �′ → � and SSS → Mgluster/Mg–Ag co-cluster → � → � [4–6]. The thickening rate of �hase, however, was confirmed to be much smaller than that of �′

hase even when exposed at elevated temperature up to 300 ◦C,howing a great thermal stability [7].

With the development of the aerospace industries, mechanicalroperties of aluminum alloys are required to be further improvedo satisfy the applications [8]. New heat treatments can be devel-ped to obtain the prescribed microstructures and properties in aide range of age-hardenable aluminum alloys. The Al–Cu–Mg–Ag

lloys show superior creep resistance in the underaged condition

ather than in fully hardened T6 temper [9]. Recently, multiple-tage aging heat treatments (interrupted aging treatment or T6I6)ave been developed to increase the strength and fracture tough-

∗ Corresponding author at: School of Materials Science and Engineering, Centralouth University, Lu Mountain South Road, Changsha 410083, China.el.: +86 731 88836011; fax: +86 731 88876692.

E-mail address: liuzhiyi@mail.csu.edu.cn (Z. Liu).

921-5093/$ – see front matter © 2011 Elsevier B.V. All rights reserved.oi:10.1016/j.msea.2011.07.055

ness, which involves that the T6 treatment is interrupted by agingat a lower temperature (25–65 ◦C) before resuming the final agingat the temperature for the initial T6 treatment, or at anotherdifferent elevated temperature [10–15]. These heat treatmentswere successfully applied to Al–Cu, Al–Cu–Mg–(Ag), Al–Mg–Si andAl–Zn–Mg–Cu alloys [10–15]. The increase of strength and fracturetoughness is attributed to a secondary precipitation that occurs at alower temperature, a finer and denser dispersion of strengtheningphase is formed when T6 aging is resumed.

In commercial practice, a pre-straining processing is usuallyapplied to straightening of products of wrought aluminum alloys inthe as-quenched state. The presence of dislocations that generatedin this process significantly influences the subsequent precipitationprocess and final mechanical response of the alloy. Based on a studyby Ünlü et al. [16], the pre-straining process increases the mechan-ical properties in Al–5.0Cu–0.5Mg (wt%) alloy due to an increasingnumber density and a refinement of precipitates. However, Ringeret al. [17] revealed that dislocations introduced by cold-workingprior to aging interfered with nucleation of the � phase and pro-vided sites to facilitate heterogeneous nucleation of �′ precipitates,the peak hardness values of Al–4Cu–0.3Mg–0.4Ag (wt%) alloy agedat 165 ◦C and 200 ◦C are reduced by 4.5% and 7.5%, respectively.

Although the T6I6 heat treatment in an Al–Cu–Mg–Ag alloy isreported [11], it cost too much time during the lower-temperatureaging and there is no report so far about improvement of themechanical properties of the alloy subjected to the pre-strainingprocessing. In order to reduce the unfavorable effect of pre-

straining on the mechanical properties and assist the precipitationof � phase in Al–Cu–Mg–Ag alloys, the present work aims todevelop a type of duplex aging treatment, indicating that the firstaging at 200 ◦C for 20 min was interrupted by aging at 165 ◦C.

Y. Li et al. / Materials Science and Engineering A 528 (2011) 8060– 8064 8061

Table 1Description of heat treatments for Al–Cu–Mg–Ag alloy.

Number Heat treatment Description and aging times

◦ Aging at 200 ◦C for up to 24 hAging at 165 ◦C for up to 100 hUnder aging at 200 ◦C for 20 min Secondary aging at 165 ◦C for up to 100 h

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Table 2Tensile properties of Al–Cu–Mg–Ag alloy in T8 and T8L6 peak-aged conditions testedat RT and elevated temperatures.

Temperature (◦C) Temper UTS (MPa) YS (MPa) Elongation (%)

20 T8/200 492 461 9.3T8/165 495 460 10.7T8L6/165 508 476 10.7

250 T8/200 309 300 12.4T8/165 313 303 12.2T8L6/165 323 315 11.4

300 T8/200 212 209 13.2

1 T8/200 ST at 515 C for 6 h + WQ + 2% pre-straining2 T8/1653 T8L6/165

. Materials and methods

The chemical composition of the experimental material used inhe present work was Al–4.94Cu–0.43Mg–1.04Ag–0.3Mn–0.15Zrwt%). The ingot was homogenized after casting, and hot rolledo 2 mm thick strip at about 460 ◦C. All the samples wereolution-treated (ST) at 515 ◦C for 6 h, water-quenched (WQ) andmmediately stretched with 2%. Then the pre-strained specimens

ere aged at 165 ◦C and 200 ◦C (hereafter termed T8/165 and8/200, respectively), or aged at 200 ◦C for 20 min followed by65 ◦C for times up to 100 h. This new type of duplex aging heatreatment was termed T8L6/165. Details of these heat treatmentsre shown in Table 1.

Hardness measurement was performed on a HV-10B Vickersardness tester with a load of 3 kg. The hardness values reportedere represent the average of at least ten measurements. Tensileesting was conducted on a SANS-CMJ5105 testing machine at roomemperature (RT) and elevated temperatures (250 ◦C and 300 ◦C)ith 2 mm/min loading speed. The values of strength and ductil-

ty were the mean values of three specimens. Differential scanningalorimeter (DSC) analysis was carried out on a NETZSCH SAT 449Calorimeter using high purity aluminum as a reference. Disc-likeamples with a thickness of about 0.5 mm and diameter 5 mmere scanned at a heating rate of 10 ◦C/min in the temperature

ange from 20 ◦C to 400 ◦C. Specimens for transmission electronicroscopy (TEM) were prepared by using twin-jet electrolytically

olishing with a voltage of 10–15 V in a solution of 70% methanolnd 30% nitric acid at −20 ◦C. The TEM observations were carriedut on a TECNAI G220 transmission electron microscopy operatedt 200 kV.

. Results and discussion

Fig. 1 illustrates the hardness curves for Al–Cu–Mg–Ag alloy

n the T8 and T8L6 conditions. The peak hardness of the T8 alloyecreased as the aging temperature was elevated from 165 to00 ◦C, meanwhile the peak-aged time decreased from 20 h to 2 h.

Fig. 1. Comparison of T8/165, T8/200 and T8L6/165 hardness curves.

T8/165 203 199 12.8T8L6/165 217 213 12.4

The hardness of the T8L6/165 temper significantly increased incompared with those of T8/165 and T8/200 tempers, whereas thepeak-aged time for T8L6/165 and T8/165 tempers is almost thesame. The peak hardness value of T8L6/165 temper is 162 HV, whichexceeded the T8/165 and T8/200 peak-aged condition by 3% and 5%,respectively.

The samples were all peak-aged treated for T8 and T8L6/165tempers, and then tested at temperatures up to 300 ◦C. Table 2 pro-vides a comparison of the tensile properties for the three tempersexamined. It can be seen that T8L6/165 heat treatment producedimprovements in the strength at RT and elevated temperatures,while the elongation was still at a high level. Comparing with theultimate tensile strength (UTS) and yield strength (YS) of T8/165temper, those of T8L6/165 temper were increased by 13 and 16 MPaat RT, respectively. Furthermore, an improvement of approximately7% in the UTS and YS at 300 ◦C was achieved. It should be noted thatthe strength of T8/165 temper at RT and 250 ◦C was superior to thatof T8/200 temper, whereas it was less at 300 ◦C. This was related tothe type and number density of strengthening particles.

To demonstrate the effect of aging temperature on precipitationof � and �′ precipitates in the Al–Cu–Mg–Ag alloy, DSC samples

were prepared after hardness and tensile tests. Fig. 2 shows a typi-cal DSC thermogram for the specimens of the Al–Cu–Mg–Ag alloy.According to literature [17–19], three peaks can be identified in

Fig. 2. DSC thermograms for Al–Cu–Mg–Ag alloy under different conditions.

8062 Y. Li et al. / Materials Science and Engineering A 528 (2011) 8060– 8064

Fig. 3. TEM images of the Al–Cu–Mg–Ag alloy under different conditions: (a) underaged at 200 ◦C for 20 min (UA), dominantly �; (b) peak-aged to T8/200, � + a minor �′

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hase; (c) peak-aged to T8/165, � + large quantities of �′ precipitates; (d) peak-ag0 1 1〉� .

urve 1 (as-quenched stage), three peaks can be identified, whichre associated with the GP zone dissolution (peak A), the � pre-ipitation (peak B), and the �′ phase precipitation (peak C). Afterging at 200 ◦C for 20 min (UA), the area of peak B reduces and theariation of the peak C is not obvious, suggesting that � precipi-ation predominated the aging process at 200 ◦C. Comparing withhe curves 3–5, the precipitation behavior of � phase was differentven though all three samples were peak-aged temper. The peak

of both the T8/200 and T8L6/165 tempers disappeared (curves 4nd 5), indicating the complete precipitation of the � phase. Theccurrence of great reduction of the peak B in curve 3 indicatedn incomplete precipitation of � phase in T8/165 temper condi-ion. Subsequent precipitation should continue when a high agingemperature was employed. It is thus confirmed that high agingemperature facilitates the precipitation of the � phase.

Fig. 3(a) is a TEM image from the Al–Cu–Mg–Ag alloy agedt 200 ◦C for 20 min together with the corresponding SAED pat-ern, a uniform dispersion of fine and nano-scale precipitates wasetected, reflections at the 1/3 and 2/3 g {0 2 2}˛ positions in0 1 1〉˛ SAED patterns suggested the presence of � phase although

he intensity was extremely weak. Fig. 3(b–d) shows the peak-ged T8/200, T8/165 and T8L6/165 microstructures, respectively.

dense and uniform distribution of � phase was predominant and′ phase was only present as a minor phase in the T8/200 temper.

T8L6/165, � + small quantities of �′ precipitates. The electron beam is parallel to

The corresponding SAED pattern in Fig. 3(b) did not show appar-ent diffraction spots of �′ phase, only reflections of � phase werefound. Fig. 3(c) shows the peak-aged T8/200 microstructure, wherea large amount of � and �′ precipitates was found. The diffractionspots of �′ phase at 1/2 g {0 2 2}˛ were also observed in SAED pat-tern except for the reflections of the � phase. Fig. 3(d) is a TEM andcorresponding SAED pattern of the T8L6/165 temper. The dominantphase was � phase and small quantities of �′ precipitates were alsodetected, in accordance with the SAED pattern.

It is well known that cold deformation could obviously increasethe density of dislocations in a metallic material. Fig. 4 is the TEMimage for the Al–Cu–Mg–Ag alloy in the as-quenched and pre-strained condition; high density of dislocations was clearly visiblewithin the grains. Ringer et al. [20] investigated the microstructureevolution of Al–Cu–Mg alloy during the aging process, and foundthat a very high density of dislocation loops was detected in theas-quenched stage and the dislocation loops still existed until theaging time reached 150 min at 180 ◦C. Nevertheless it was hard toobserve the dislocations in the UA condition (Fig. 3(a)), which indi-cated the density of dislocations would be significantly reduced and

it is easier for the recovery of dislocations at high temperature.

The strengthening of age-hardenable aluminum alloys dependsgreatly on the type, size, orientation, distribution, morphology, andnumber density of the strengthening phases [13]. Based on the

Y. Li et al. / Materials Science and Engin

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ig. 4. TEM micrograph showing dislocations of pre-strained Al–Cu–Mg–Ag alloyrior to aging.

esearches by Zhu and Starke [21] and Liu et al. [22], providingnly one kind of strengthening phase exits in the alloy and it isonsidered as unshearable, moving dislocations have to bow outnd bypass between adjacent precipitates in order to slip further,he increment in the yield strength (��) due to the strengtheninghase can be expressed as:

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here M denotes the Taylor factor, G the shear modulus, b theagnitude of Burgers vector, r0 the inner cut-off radius for the cal-

ulation of dislocation line tension, and r, h and fv are the radius ofabit plane, the half-thickness of peripheral plane and volume frac-ion of plate-like precipitates, respectively. It can be seen from Eqs.1) and (2) that the strengthening effect of precipitates on {1 1 1} isarger than precipitates on {1 0 0}. Moreover, the thickening rate of

phase is smaller than that of �′ phase, especially at high tem-erature. As a result, although both �′ and � contribute to ageardening, the strengthening effect of �′ phase on {1 0 0} is weakerhan � phase on {1 1 1}. Aging temperature and dislocations inhis study have significant effects on the competitive precipitationetween � and �′ phases, thereby influencing the mechanical prop-rties of the alloy. The driving force for particle nucleation and theucleation rate for precipitates are much greater at 165 ◦C thant 200 ◦C [23]. More strengthening particles precipitated at 165 ◦C,he tensile properties of the T8/165 temper at RT and 250 ◦C wereigher than those of the T8/200 temper.

On the other hand, the dislocations introduced by pre-strainingrior to aging served as sites for the heterogeneous nucleation of′ phase. The � phase evolves from Ag and Mg co-clusters con-inuously, rather than discretely by heterogeneous nucleation toome precursor phase [24]. Both � and �′ phases contained Cu, the

recipitation of a large quantity of the �′ phase consumed the Cutoms in the matrix, suppressing the � precipitation [25]. How-ver, the recovery of dislocations began to occur at approximately82 ◦C [16] and a greater extent of recovery took place at 200 ◦C.

eering A 528 (2011) 8060– 8064 8063

(for {1 1 1} plate-like precipitates) (1)

(for {1 1 1} plate-like precipitates) (2)

Therefore, DSC and TEM indicated that the precipitation of �′ phasewas restrained and the formation of � phase was promoted at hightemperature, consistent with the results of Refs. [4,26]. More �phase and less �′ plates were observed at 200 ◦C (Fig. 3(b)); both thethickening rate and dissolution rate of the �′ phase is in excess thanthat of the � phase at high temperatures, thereby the strength ofT8/200 temper at 300 ◦C was improved in contrast to that of T8/165temper.

For the T8L6 temper, the recovery of dislocations generated bypre-straining prior to aging occurred when the alloy pre-aged at200 ◦C for 20 min. The aim of first aging at higher temperature wasto obtain fine � precipitates and restrain the formation of �′ phaseduring the early stage of aging (Fig. 3(a)), reducing the unfavor-able influence of dislocations on � precipitation. If following theUA treatment, further aging is continued at 200 ◦C (T8/200), smallqualities of strengthening particles were formed even though the� phase was predominant (Fig. 3(b)). Similarly, less � phase pre-cipitated although large qualities of strengthening particles werepresent in the T8/165 sample (Fig. 3(c)). However, if 200 ◦C/20 minaging process was interrupted, the next step of aging was contin-ued at a lower temperature of 165 ◦C, the � phase and co-clustersformed during the UA treatment would be respectively expected touniformly grow and subsequently transform to precipitates duringthe lower temperature aging. Moreover the dislocation density wasdramatically degraded during first stage of aging, greatly reducingthe nucleation site of �′ particles for the subsequent second stage ofaging processing at 165 ◦C. In addition, the critical size nucleus for �and �′ formation is smaller at 165 ◦C than at 200 ◦C, the probabilityof the co-clusters evolving into precipitates would be increased atthe lower temperature, leading to the formation of more � precip-itates and improving mechanical properties of Al–Cu–Mg–Ag alloyat RT and elevated temperatures.

4. Conclusions

In summary, a type of duplex aging was employed to obtainmore � precipitates by controlling the competitive precipitationbetween � and �′ phases, increasing the aging hardness of pre-strained Al–Cu–Mg–Ag alloys. Comparing with the conventionalT8 peak-aged temper, the UTS and YS of the T8L6/165 peak-agedtemper at RT and elevated temperatures were enhanced by approx-imately 3–7%, whereas the elongation kept at a high level.

Acknowledgement

The authors are grateful for the financial support of National KeyFundamental Research Project of China.

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