Enhanced grain refinement in AA7050 Al alloy by deformation-induced precipitation

Materials Science and Engineering A 549 (2012) 100–104

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Enhanced grain refinement in AA7050 Al alloy by deformation-induced precipitation Yuanhua Cai ∗ , Yujing Lang, Lingyong Cao, Jishan Zhang State Key Laboratory for Advanced Metals and Materials, University of Science and Technology Beijing, Beijing 100083, China

a r t i c l e

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Article history: Received 9 December 2011 Received in revised form 22 March 2012 Accepted 5 April 2012 Available online 13 April 2012 Keywords: Al Alloys Grain refinement Double-step thermomechanical treatment Mechanical properties

a b s t r a c t In this study, an improved double-step thermomechanical treatment was proposed for producing fine grain structure of high-strength Al–Zn–Mg–Cu series alloy based on deformation-induced precipitation. The new method composed of a low temperature pre-deformation and a traditional hot rolling. Effects of this proposed treatment on microstructure and properties of the alloy were studied using optical microscopy, transmission electron microscopy, X-ray diffraction, and tensile tests. The microstructural analysis results show that the pre-deformation can accelerate both formation and spheroidization of the fine precipitates. Those deformation-induced precipitates could exert more drag force on migrations of the grain boundaries. So, recovery and coarseness of those deformed grains can be suppressed during subsequent heating and hot rolling, contributing to the final fine-grain structure of the AA7050 alloy. The properties test results reveal that the alloy processed by the double-step thermomechanical treatment displays superior strength and elongation than the conventional hot-rolled alloy because of the fine grain. Crown Copyright © 2012 Published by Elsevier B.V. All rights reserved.

1. Introduction As one of the key materials, the Al–Zn–Mg–Cu alloys are widely used in the aircraft and automotive industries for their excellent combination of high strength to weight ratio, toughness, corrosion resistance and so on [1]. And further improvement of properties and performance are greatly desired to exert the materials potential for safety, light-weighting and fuel efficiency. Grain refinement often results in favorable changes of mechanical properties in Al alloys, which can meet the increasing performance demand of these alloys in aerospace and automobile industries [2]. So, many efforts have been conducted to achieve fine grain for the high-strength aluminum alloy by improving thermomechanical treatment (TMT) [3–6] and developing severe plastic deformation (SPD) techniques [7–18] as well composition adjustment [19]. For structural applications, Al–Zn–Mg–Cu alloys are generally required of reasonable dimensions in substantial quantities and low fabrication cost. Thermomechanical processing which involves the deformation and annealing of bulk alloys is usually considered to be the optimum method for producing such fine-grained alloys. Traditional thermomechanical treatment of alloys often results

∗ Corresponding author at: No. 30, Xueyuan Road, Haidian District, Beijing 100083, China. Tel.: +86 10 62334862; fax: +86 10 62333447. E-mail address: [email protected] (Y. Cai).

in grain sizes about 30–250 ␮m [20]. In order to achieve further fine grains with size less than 30 ␮m, Wert et al. developed a new or improved TMT scheme [4]. In their TMT, great number of precipitates larger than ∼0.75 ␮m should be formed first. Those large particles stimulate recrystallization of deformed grains during subsequent annealing (which was known as particle-stimulated nucleation of recrystallization, PSN). Thus, the alloy grains could be refined to the scale of about 10 ␮m suiting for superplastic forming. However, this improved TMT caused an increase of cost, and producing finer grain structures needed careful controls of the material and processing parameters. Furthermore, those large precipitates could not redissolve into Al alloy completely during post heat treatment, leading to the loss of tensile strength. SPD has its own advantages in producing stable fine or ultrafine grain for the simultaneous increase of strength and ductility. However, those SPD methods could hardly be used directly to produce large-size Al alloy parts, especially those ultra-thick plate parts used in aerospace industry [21]. Up to now, innovative thermomechanical treatment is still preferred method for grain refinement of aluminum alloys during industrial production and is also of great practical importance. Here, we propose a new double-step thermomechanical method to produce fine-grained Al–Zn–Mg–Cu alloys according to deformation-induced precipitation (DIP) widely used in advanced steel production [22]. And the effect of such method on microstructural evolution and mechanical properties was studied.

0921-5093/$ – see front matter. Crown Copyright © 2012 Published by Elsevier B.V. All rights reserved. http://dx.doi.org/10.1016/j.msea.2012.04.011

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Table 1 Chemical composition of AA7050 alloy (in wt%). Elements

Zn

Mg

Cu

Zr

Fe

Si

Mn

Cr

Ti

Al

wt%

5.7–6.7

1.9–2.6

2.0–2.6

0.08–0.15

<0.15

<0.12

<0.10

<0.04

<0.06

Bal.

Table 2 Description of prescribed procedures. Designation

Processing regimes

HQ

The SQ-treated sample was heated to 573 K with a holding of 120 s The SQ-treated sample was heated to 573 K, held for 120 s, then deformed with a reduction of 60% Conventional hot rolling. The SQ-treated sample was heated to 673 K, held for 120 s, deformed with a reduction of 80%, one-pass deformation Double-step hot rolling. The SQ-treated sample was heated to 573 K, held for 120 s, then pre-deformed with a reduction of 60% (first-step, DIP treatment), and sequentially heated to 673 K with a holding of 10 s at 673 K followed by deforming with a 50% reduction of DIP sample (second-step), the total reduction of double-step hot rolling is also 80%

DIP CHR

DHR

2. Experimental The investigations were carried out on a commercial AA7050 aluminum alloy, and the nominal composition of the alloy was shown in Table 1. The alloy was solutionized at 753 K for 16 h and 756 K for 8 h, and then rapidly quenched into room-temperature water (SQ). Cylindrical samples with 10 mm in diameter and 15 mm in height were machined from the SQ-treated alloy. All compressing deformations were carried out under the procedures as listed in Table 2 on the Gleeble-1500 thermomechanical simulator (GTMS) with a heating rate of 2 K/s and a strain rate of 10 s−1 . The deformed specimen was immediately water quenched to room temperature for preserving the hot-deformed microstructures. The microstructure was studied using optical microscopy (OM), transition electron microscopy (TEM) and high resolution electron microscopy (HREM). Specimens for OM observations were prepared by the standard metallographic procedures using SiC abrasive paper, 2.5 ␮m diamond polishing paste and Keller’s reagent. The thin foils for TEM and HREM studies were prepared by twinjet electropolishing in 33% nitric acid solution in methanol. TEM and HREM studies were performed on Hitachi H-800 TEM and FEI-TF20 HREM. X-ray diffraction (XRD) analyses were carried out

with Rigaku X-ray diffractometer (D/MAX-RB) using Cu K␣ radiation ( = 0.15406 nm). Mechanical property tests were carried out on a MTS-810 test machine at room temperature with a nominal strain rate of 10−3 s−1 . For tensile tests, both CHR-treated and DHR-treated samples were solution treated at 748 K for 30 min and aged at 393 K for 24 h (T6 treatment). All tensile test samples had 3 mm width and 1.5 mm thickness and a 15 mm gauge length. The accepted value for tensile strength or elongation was the average of three measurements. 3. Results and discussion 3.1. Microstructural evolution Fig. 1(a) shows the optical graph of the SQ-treated materials. Obviously, after a long time holding at the high temperature, the recrystallized grains are dominant in the SQ-treated alloy. And the recrystallized grains were grown to a large size of 600 ␮m in length and 50–100 ␮m in transverse. The TEM graph of such SQ-treated alloy shows the complete dissolution of the soluble second particles of (MgZn2 ), T(Al2 Mg3 Zn3 ) or S(Al2 CuMg) phases usually present in Al–Zn–Mg–Cu alloys, and the large size subgrain of about 8 ␮m (Fig. 1(b)). Fig. 2(a) revealed grain microstructures of the DIP-treated alloy. As one could see that plenty of dislocations were formed, piled up and tangled together when the SQ-treated alloy was processed during DIP stage. Those dislocations formed small elongated cells with different sizes rather than polygonizing to well-ordered sub-grains because of low deformation temperature and high deformation rate [20]. The formation of dislocation cells indicated that the hot deformation was predominant during DIP stage. Simultaneously, large number of fine precipitates were formed and uniformly distributed throughout the alloy. Most precipitates were globular and had a size of 20 nm in diameter, only a few of spherical precipitates were about 40 nm in diameter (Fig. 2b). XRD analysis, shown in Fig. 2(c), indicated that fine precipitates were mainly MgZn2 . For discovering the effect of low temperature deformation on precipitates, the SQ-treated alloy was also HQ-treated. TEM graphs

Fig. 1. Microstructures of the SQ-treated alloy (a) OM image and (b) TEM image.

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Y. Cai et al. / Materials Science and Engineering A 549 (2012) 100–104

Fig. 2. TEM graphs and XRD pattern of the DIP-treated alloy (a) grain structure, (b) precipitates and (c) XRD pattern.

Fig. 3. TEM images of the HQ-treated alloy (a) grain structure and (b) precipitates (the insert showed the whole morphology of a precipitate).

about grain structure and precipitate of the HQ-treated alloy were shown in Fig. 3(a) and (b) respectively. Based on the graphs, one could see precipitates were formed both along grain boundaries and in grain interiors. The grain boundary precipitates were larger than those in grain interior. For the HQ-treated alloy, all the precipitates were rod-like or elliptical with a size rang of 20–100 nm in length and 10–20 nm in transverse. The number density of those precipitates in HQ-treated alloys was smaller than that of DIPtreated alloys, and contrarily, the precipitates size in HQ-treated alloys was larger than that of DIP-treated alloy.Obviously, the low temperature pre-deformation enhanced the formation, the spheroidization and the refinement of those precipitates, which was consistent with prior research results [23,24]. During DIPtreatment, the deformation could introduce many defects, such as vacancies and dislocations, into the alloy. Those defects could afford adequate energy fluctuation, structure fluctuation and rapid diffusion paths of solutes for easy formation of precipitates. The internal deformation strain may restrict the growth of those precipitates toward preferred direction leading to spheroidization and refinement of such precipitates. On the other hand, the deformation could also result in fragment of those rod-like precipitates. So the precipitate in DIP-treated alloy was finer than HQ-treated alloy with a larger number density. Microstructures of the DHR-treated and CHR-treated alloys were shown in Fig. 4. After DHR treatment, the grains of SQ-treated alloy were markedly refined to the size less than 2 ␮m in transverse,

and most of subgrains were refined to about 2 ␮m or less in transverse. The grains of CHR-treated alloy were comparatively large and composed of well-developed subgrains. The grains of CHR-treated alloy were about 10 ␮m in transverse, and the subgrains were about 5 ␮m in length and 2 ␮m in transverse. Obviously, the DHR-treated alloy had fine and uniformly distributed precipitates, and the round precipitates were domain with the present of a few rod-like particles in matrix. For the CHR-treated alloy, both diameter and length of the rod-like particles were larger than those in the DHR-treated alloy. The CHR could also be seen as a two-step thermomechanical treatment consisting of HQ-treatment and post high temperature hot rolling. The only difference between the CHR and the DHR was that DHR scheme had a low temperature deformation. As mentioned above, the DIP-treatment could produce more and finer precipitates than that HQ-treatment did. Commonly, the more and the smaller are the precipitates, the larger the pinning force is. So, migrations of grain and subgrain boundaries or dislocations were prevented or even inhibited. Thus, fine grains or subgrains could be formed during DHR processing. The reasons for more fine size of the precipitate in DHR alloy than its CHR counterpart can be described as follow. Firstly, more particles were there in Al alloys at low temperature, more of them could survive after a similar post heating and hot rolling. Secondly, the 50% deformation reduction of DIP-treated alloy needed less time than 80% reduction if the strain rate was the same. Thus, the

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Fig. 4. TEM micrographs of the DHR-treated (a, b) and CHR-treated alloys (c, d).

3.2. Tensile properties After solid solutionized at 748 K for 30 min and artificial aged at 393 K for 24 h (T6 condition), the DHR-processed AA7050 alloy was subjected to tensile testing at room temperature. For easy comparison, tensile testing was also performed on the CHR-treated 7050 alloy heat treated as that did for DHR-treated alloy.

650

DHR CHR

600 550 500

500

Strength, MPa

Sterngth, MPa

total time needed in the post heating and hot rolling for DHR was less than that of CHR. This was also favorable to keep a small size and a high number density of precipitates. After the high temperature conventional hot rolling, the fine globular precipitates formed during DIP stage grew to a size of 30–60 nm in diameter. At the same time, some rod-like particles of nearly 30 nm in transverse and about 100 nm in length were formed. The number density of precipitates in DHR-treated alloy decreased markedly comparing to the DIP-treated alloy. This decrease in number density of precipitates could be explained based on the solid solubility of solutes and precipitate growth. The solubility of alloying element would increase when the alloy was heated to a higher temperature. The content of alloying solutes which cannot dissolved into Al lattice became lower, leading to low equilibrium fraction and small number of precipitates remained in Al alloy. The growth of precipitate followed Ostwald ripening principle when the temperature was increased [25]. The shrinking and re-dissolving of smaller precipitates with further growing of larger one would lead to the decrease of the number density of precipitates.

450 400

400 300 200 100 0 0.0

350

0.2

0.4

0.6

0.8

1.0

Enlongation, %

300

0

2

4

6

8

10

12

14

Enlongation, % Fig. 5. Stress versus strain curves of DHR-treated and CHR-treated alloys in T6 condition. The insert is the curves in initial deformation stage.

Engineering stress and strain curves of both the T6 tempered DHR- and CHR-treated alloys were shown in Fig. 5. After T6 aging, the ultimate tensile strength and elongation of DHR-treated alloy are 600 ± 1.7 MPa and 12.1 ± 2.7%, and 592 ± 2.45 MPa and 9.1 ± 0.75% for CHR-treated alloy. It was worth noting that the DHRtreated alloy had a much higher elongation than CHR-treated alloy in T6 condition with a slight higher strength. As we know, precipitation hardening play domain role in hardening response of

Fig. 6. Microstructure of aged DHR alloy (a) grain structure and (b) precipitates.

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heat-treatable Al alloy, and grain refinement contribute limited effect to the strength. But grain refinement can play important role in improving elongation. Obviously, grain refinement played a key role for the increases of elongation because of the similar solution and aging treatments conducted on both CHR-treated and DHR-treated alloys [26,27]. The microstructure of DHR-treated alloy in T6 condition could be seen in Fig. 6. After solid solutionized and T6 tempered, the DHR-treated alloy had very fine grain structure, and large number of nano-scaled particles precipitated throughout the matrix, which mainly contributed to the aging hardening response during artificial aging for AA7050 alloy. It also can be seen that the precipitation free zone (PFZ) is very narrow DHR-treated alloy, which is also beneficial to the mechanical properties. Apart from the fine nano-scaled particles, a lot of discrete large precipitates are also formed along the grain boundary. According to some literatures [28–30], the distribution of those isolated large grain boundary precipitates is favorable to the increase of stress corrosion cracking (SCC) resistance. 4. Conclusions (1) DIP-treatment enhances both the formation and the spheroidization of a large quantity of fine precipitates through the alloy. This great number of fine particles can result in very strong pinning force on the migration of the grain boundary or dislocation. Subsequently, the coarsening of dislocation cells or grains is suppressed. (2) The double-step thermomechanical rolling (DHR-treatment) can lead to not only more fine grain structure but also more fine global and elliptical precipitates than the CHR-treatment. After T6 temper, the DHR-treated AA7050 alloy displays superior combination of strength and elongation than the CHR-treated alloy. The grain refinement plays key effect on the increase of elongation of DHR-treated alloy. Acknowledgment The authors are grateful for the support provided by State Key Laboratory for Advanced Metals and Materials of China.

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