Influence of Cu on microstructure and tensile properties of 7XXX series aluminum alloy

Accepted Manuscript Short Communication Influence of Cu on microstructure and tensile properties of 7XXX series aluminum alloy Yu-guo Liao, Xiao-qi Han, Miao-xia Zeng, Man Jin PII: DOI: Reference:

S0261-3069(14)00365-3 http://dx.doi.org/10.1016/j.matdes.2014.05.003 JMAD 6481

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Materials and Design

Please cite this article as: Liao, Y-g., Han, X-q., Zeng, M-x., Jin, M., Influence of Cu on microstructure and tensile properties of 7XXX series aluminum alloy, Materials and Design (2014), doi: http://dx.doi.org/10.1016/j.matdes. 2014.05.003

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Influence of Cu on microstructure and tensile properties of 7XXX series aluminum alloy Yu-guo Liao, Xiao-qi Han, Miao-xia Zeng, Man Jin1＊ School of Materials Science and Engineering, Shanghai University, Shanghai 200072, China

Abstract: The effect of copper content on tensile properties and microstructure of 7XXX series aluminum alloys at 90% deformation is investigated by tensile test, optical microscopy, scanning electronic microscopy (SEM), X-ray diffraction (XRD) and transmission electron microscopy (TEM). The results show that higher Cu-containing alloy would precipitate more quantity of second-phase particles in rolling process at 420°C, which facilitates the process of recrystallization during solution treatment. With the increase of copper content from 0 to 1.6 wt.%, the density of ç’ phase and the degree of recrystallization increase, meanwhile the strength and plasticity of Al-Zn-Mg-Cu alloy are improved in T6 heat treatment. In addition, the tensile fracture of Cu-free alloy belongs to quasi-cleavage fracture, while the tensile fracture of Cu-containing alloy is dimple fracture and the number of dimples increases with the increasing of copper content. Keywords: Al-Zn-Mg-Cu alloys; Copper; Recrystallization; Tensile properties

1. Introduction Evaluating the composition of microstructure in multi-composition aluminum alloys has been the subject of much research effort in the past decade [1-4]. Copper containing 7000 series Al alloy (based in the system Al-Zn-Mg-Cu) are widely used for aeroplane wing structures. Since Webber developed the first Al-Zn-Mg-Cu series aluminum alloy Al-10% Zn-2% Mg-2% Cu-1% Mn (wt.%) in 1932, the effect of copper on Al-Zn-Mg-Cu aluminum alloy has been extensively studied [5]. It is well documented that addition of copper to high-strength A1-Zn-Mg alloys has been developed to increase the volume fraction of strengthening precipitates, and to improve the yield strength and stress corrosion cracking (SCC) resistance simultaneously [6-8]. In the last decade, Chinh [9] 1＊ Corresponding

author. Tel.:+86 21 56332127; fax: +86 37821398.

E-mail: [email protected]

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found that Cu atom dissolves in GP zones and changes it from spherical to ellipsoidal and also improves the density of the strengthening phases. Hadjadj [10] reported that the addition of Cu can refine the grain and increase the microhardness. Marlaud [4] observed that precipitate is easy coarsening in the lowest Cu containing alloy due to slower diffusivity of Cu in Al compared to Mg and Zn atoms. Even though there is a number of papers about Cu to 7XXX alloy [4, 7, 8, 9, 10, 11], but the influence of copper content on recrystallization and plasticity of 7XXX series aluminum alloys is relatively uncertain. In this paper the effects of Cu addition on the microstructure and tensile properties of 7XXX series aluminum alloy are presented especially about the recrystallization and plasticity.

Microstructures in different heat treatment stages of both ternary and

quaternary alloys were investigated by optical microscopy (OM), scanning electronic microscopy (SEM), X-ray diffraction (XRD) and transmission electron microscopy (TEM), and the tensile properties after T6 heat treatment of alloys were investigated by tensile tests.

2. Materials and methods The investigated Al-7.8% Zn-1.6% Mg-x Cu-0.14% Zr (x=0, 0.8, 1.6 wt.%) alloys were prepared in laboratory by ingot metallurgical process.

The raw materials were

commercially pure Al (99.98%), Zn (99.98%), Mg (99.75%) and Al-4.21% Zr, Al-42.3% Cu (wt.%). The alloys were melted in a graphite crucible heated by an electrical resistance furnace. The liquid metal was then poured into semi-continuous casting mold of round billets with electromagnetic stirring to get a Φ72 mm round ingot. The ingots were homogenized at 450℃ for 24h in industry furnace and cooled in air. The chemical compositions of the alloys were analyzed on PMI-MASTER PRO spectrometer and the results were shown in Table 1. The samples were processed into 20 mm thick plates. All alloys were multi-pass rolled from 20 mm to 2 mm at 420℃ and cooled in air. Five passes rolling were used, and the sheets were placed back to the furnace for 0.5 h between every two passes. After solution treatment at 470℃ for 2h and quenching in water, the samples were aged at 120℃ for 24h (T6 heat treatment). 2

The microstructure was studied using optical microscopy (OM), scanning electronic microscopy (SEM) and transmission electron microscopy (TEM).

An Lv-150

metalloscope was used to obtain the metallographic photos and a HITACHI SU-1500 SEM was used to observe the morphology of precipitated phases and fracture. X-ray diffraction (XRD) analyze was carried out with 18 KW D/MAX2500V+/PC using Cu Kα radiation (λ=0.15406 nm). The samples after rolling at 420 ℃ before solution treatment were etched with Keller’s reagent (1 ml HF, 1.5 ml HCl, 2.5 ml HNO3, and 95ml distilled water), and the samples with T6 treatment were etched with Graff Sargent’s reagent (3 g Cr2O3, 16 ml HNO3, 1 ml HF, 85 ml distilled H2O). Thin foils for TEM observation were punched out directly from samples which were ground down to 0.5 mm thickness after ageing.

There disks were mechanically thinned down to ~ 150 μm then

electropolished using a twinjet Tenupol with a 33% nitric acid solution in methanol maintained at -20 ℃ and 15 V. TEM examinations were performed using a JEM-2010F, operating at 200 kV. Tensile properties testing were carried on a CMT5105 test machine at room temperature with strain rate of 0.001/s. The tensile tests were carried on rolling samples after T6 heat treatment. All tensile samples were 12.5 mm in width and 2 mm in thickness and 30 mm in gauge length. The accepted values for tensile strength and elongation were the average of three measurements.

3. Results and discussion 3.1. Microstructural evolution Fig.1. shows the SEM macrographs of the samples after rolling at 420°C. It can be seen that the number of large second-phase particles increases with the increase of copper content, and the size of second-phase particles is about 1-4 μm. For comparison, the volume fraction of particles remaining after rolling at 420 ℃ of 1#, 2#, 3# alloys is 6.1%, 9.8%, 14.5%, respectively. The second-phase particles of alloys precipitate in the area with high energy along the rolling direction. The large second-phase particles in 1# and 2# alloys distribute along the rolling direction, but disperse in the matrix in 3# alloy. 3

Copper addition improves the supersaturation of alloy and promotes the precipitation of Mg and Zn [12], and some previous researches show that copper addition also increases the quench sensitivity of alloy [11] and alters the kinetics of the precipitation reaction [8, 13, 14], resulting in the improvement of the driving force of nucleation, which increases the amount of the precipitated phases.

Therefore, the volume fraction of the

second-phase particles in 3# alloy is the largest among three samples and the particles disperse in the matrix with no directionality. In order to determine the second-phase particles of Fig.1, the XRD analyses of alloys are tested and showed in Fig.2. It is found that there are three main second-phase particles in Fig.2, namely η, T, Al7Cu2Fe phase. The diffraction peaks of the second-phase particles are stronger and offset to high angle in the alloys with higher copper content. For Cu atom is smaller than Al, Mg and Zn atom, the lattice constant of precipitated phases and aluminum matrix become smaller when Cu atom dissolved in matrix and precipitated phases instead of Al, Mg and Zn atom. Thus the diffraction peaks offset to high angle. The result of X-ray diffraction shows that the spectrum of T phase was Mg32(AlZn)49, and Poganitsch [15] observed Cu could dissolve in Mg32(AlZn)49 in a broad extent, and Li [16] pointed out Al14Mg33Zn37Cu13 had the same crystal structure with Mg32(AlZn)49. Based on the above analysis, the T phase in 1# alloy is Mg32(AlZn)49, and the T phase in 2# and 3# alloys is Al14Mg33Zn37Cu13. The equilibrium phase η has the hexagonal structure MgZn2, however, it has been widely shown that some substitution of Cu and Al occurs at small sizes [2, 17], and the chemical composition of Mg(Zn,Al,Cu)2 has been proposed in Cu containing 7000 series Al alloys [18]. Therefore, Cu atom can also dissolve in η phase and Al7Cu2Fe phase. The diffraction peaks of the second-phase particles are stronger with higher copper content, which agrees with number of second phase particles in Fig.1. Fig.3 shows the metallographic images of 1#, 2#, 3# alloys after T6 heat treatment. Graff Sargent’s reagent preferentially attacks grain and subgrain boundaries, thus unrecrystallized regions appear dark due to the etching of substructure [19]. Fig.3 (a) shows that 1# alloy is characterized by a well-developed subgrain structure with few recrystallization. However, the recrystallized structures of 2# and 3# alloys are elongated along the rolling direction and the degree of recrystallization increases with copper content

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increasing from 0 to 1.6 wt.%.

The statistics of recrystallization volume fraction is

measured by Image-Pro Plus software, and five pictures of each sample were used to get average value. The average recrystallization volume fraction of 1#, 2#, 3# alloys is 3%, 27%, 55% respectively.

Degischer [20] found that the second-phase particles induce

nucleation and facilitate the process of recrystallization during solution treatment when the size of the large second-phase particles is larger than 1 μm. The number of the large second-phase particles in alloys after rolling increases with higher copper addition (Fig.1). The large second-phase particles act as the heterogeneous nucleation site of recrystallization and the recrystallization grains grow in the high energy area along the rolling direction during the solution treatment. And the alloy with higher copper content has a larger supersaturation during solution treatment, which increases the driving force of diffusion that promotes the recrystallization of alloys. Therefore, the recrystallization volume fraction of alloy increases with the increasing of copper content. Fig.4 shows TEM micrographs of the samples after T6 heat treatment. Fig.4 (a), (b), (c) are bright field phases of 1#, 2#, 3# alloy and Fig.4 (d), (f) are dark field of Fig.4 (a), (c), respectively. Fig.4 (e) is the diffraction patterns of Fig.4 (b) and it identifies that the precipitated phases are η’ phase. The major precipitated phase after T6 heat treatment is the spherical η’ phase with a size of 10-30 nm in diameter, which is dispersed in the matrix. The size of η’ phases in three kinds of alloys is substantially the same, but the number of η’ phases increases with copper content increasing.

Fig.5 shows the TEM

microstructures of alloys after T6 heat treatment. The grain boundary precipitates (GBPs) of 1#, 2# are very small and the size is under 30 nm (Fig.5 a, b), and the GBPs are coarser and more sparsely distributed in 3# alloy and the size of GBPs is in range of 15-50 nm (Fig.5 c). As copper content increasing, the size of grain boundary precipitates (GBPs) increases. The generally accepted precipitation sequence for 7XXX series alloys [21, 22] is as follows: supersaturated solid solution (SSS) → coherent GP zones → semi-coherent intermediate η’ (MgZn2) → incoherent stable η (MgZn2). The increase of η’ phase with Cu content increasing (Fig. 4) can be interpreted in two ways. On the one hand, the

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increasing of Cu content improves the supersaturation of alloy, which promotes more vacancy-rich clusters (VRCs) after quenching.

The VRCs act as heterogeneous

nucleation sites of GP zones which improves the number of GP zones and further improves the number of η’ phases. On the other hand, Cu atom can also dissolve in GP zones and η’ phases, and Mg-Cu complexes are more stable than Mg-Zn complexes [11] which can be heterogeneous nucleation sites and promote alloy comments precipitation. Some researchers also point that the addition of Cu increases the overall supersaturation of the alloy and the density of GP zones [9], resulting in a lower nucleation temperature and nucleation radius for η’ phase [11]. Therefore, the volume fraction of η’ phase increases with copper content increasing from 0 to 1.6 wt.%. 3.2. Tensile properties Engineering stress and strain curves of the T6 tempered 1#, 2#, 3# alloys are shown in Fig.6. After T6 heat treatment, the ultimate tensile strength of 1#, 2# and 3# alloys is 515 MPa, 571 MPa and 574 MPa, and the elongation is 6.9 %, 16.7 % and 19.5 %, respectively. At first, as copper content increases, the strength and ductility increases significantly, when copper addition is more than 0.8 wt.%, the performance changed slowly. It is worth noting that the 1# alloy has a much lower elongation and strength than others. As we all know, precipitation hardening plays dominant role in hardening of heat-treatable aluminum alloys, and recrystallization contribute limited effect to the strength. However, recrystallization can play an important role in improving elongation. The increasing sequence of strength corresponds with the number of precipitated phases in Fig.4, and the increasing sequence of elongation is also agreed with the recrystallization volume fraction in Fig.3. The recrystallization volume fraction of 3# alloy is much higher than 2# alloy (Fig.3), but both of them have nearly the same elongation. This is due to the coarse GBPs (Fig.5), which reduce the ductility of the alloy. The precipitated phases improve the strength and recrystallization improves the elongation. That explains why the strength of alloy is improved, while the elongation doesn’t decrease but improves with higher copper content. Fig.7 shows the fracture morphology of three alloys after static tensile. The angle between the macroscopic fracture and tensile axis is 45° for all of the samples. It can be 6

seen that 1# alloy is quasi-cleavage fracture, while 2# and 3# alloys are dimple fracture. As copper content increasing, the amount of the dimples increases. Fig.3 (a) shows that 1# alloy is almost unrecrystallized. It means the level of residential stress in 1# alloy is rather high, thus the stress concentration is prone to micro-cracks during tensile, and thence the crack expands quickly in alloy leading to quasi-cleavage fracture. The 2# and 3# alloys with higher degree of recrystallization and more precipitated phases cause a lower residential stress in matrix, and the stress concentration is prone to holes which slipped together and became dimple during tensile test.

Meanwhile, the fracture

morphology of three alloys is consistent with the elongation result in Fig.6. The addition of Cu to 7XXX series alloys changes the microstructure and tensile properties. All of those changes are interacted to each other.

As Cu was added to

aluminum alloys, it improves the supersaturation and driving force of nucleation of alloy, it thus increases the amount of the precipitated phases during rolling at 420 ℃. And the coarse second-phase particles act as the heterogeneous nucleation site of recrystallization, which increases the degree of recrystallization and further improves the plasticity. As copper addition improves the supersaturation of alloy, which causes more strengthen precipitated phases after T6 heat treatment, the strengthening phases improve the strength of alloy.

Copper addition changes the tensile properties of alloy through the

change of the microstructure.

4. Conclusions The influence of Cu on microstructure and tensile properties of 7XXX series aluminum alloys was studied. The conclusions drawn from this study are as follows: 1) The second-phase particles after rolling are η, T, Al7Cu2Fe phases and the size are about 1-4 μm.

With copper content increasing from 0 to 1.6 wt.%, the volume fraction

of particles after rolling increases from 6.1% to 14.5% and the recrystallization volume fraction increases from 3% to 55%. 2) The main strengthening phase of alloys is η’ phase with a size of 10-30 nm after T6 heat treatment.

The size of η’ phases in three kinds of alloy is substantially the same,

but the number of η’ phases increases with more copper content.

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3) With the increasing of copper content from 0 to 1.6 wt.%, the ultimate tensile strength increases from 515 MPa to 574 MPa, and elongation increases from 6.9 % to 19.5 %, respectively. 4) The tensile fracture of Cu-free alloy belongs to quasi-cleavage fracture, while that of Cu-containing alloys is dimple fracture and the number of dimples increases with the increasing of copper content.

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[11] A. Deschamps, Y. Bréchet, F. Livet. Influence of copper addition on precipitation kinetics and hardening in Al–Zn–Mg alloy. Mater. Sci. Technol 1999; 15: 993-1000. [13] A.J. Bryant. The effect of composition upon the quench-sensitivity of some Al-Zn-Mg alloy [J]. J. Inst. Metals 1966; 94: 94-99. [12] T.H. Sanders, E.A. Starke. The relationship of microstructure to monotonic and cyclic straining of two age hardening aluminum alloys. Metall. Trans. A 1976; 7a: 1407-1418. [14] C.J. Peel, P.Poole. Mechanisms of Environment Sensitive Cracking of Materials, Metals Soc., London: 1977. [15] R. Poganitsch, L. Sigl, F.J. Jeglitsch. Intermetallic Compounds in High Strength Al-Zn-Mg-Cu Alloys. Aluminium(Dusseldorf) 1981; 57: 804-807. [16] Y.X. Li, P. Li, G. Zhao, X.T. Liu, Jianzhong Cui. The constituents in Al-10Zn-2.5 Mg-2.5 Cu aluminum alloy. Mater. Sci. Eng. A 2005; 397: 204–208. [17] A. Deschamps, A. Bigot, F. Livet, P. Auger, Y. Brechet, D. Blavette. A comparative study of precipitate composition and volume fraction in an Al–Zn–Mg alloy using tomographic atom probe and small-angle X-ray scattering. Philos. Mag 2001; 81: 2391-2414. [18] T. Ramgopal, P.I. Gouma, G.S. Frankel. Role of grain-boundary precipitates and solute-depleted zone on the intergranular corrosion of aluminum alloy 7150. Corrosion 2002; 58: 687-697. [19] J.D. Robson, P.B. Prangnell. Predicting recrystallised volume fraction in aluminium alloy 7050 hot rolled plate. Mater. Sci. Technol 2002; 18: 607-614. [20] F.J. Humphreys, M. Hatherly. Recrystallization and Related Annealing Phenomena, second ed., Elsevier; Oxford; 2004. [21] H.P. Degischer, W. Lacom, A, Zahra, C.Y. Zahra. Decomposition processes in an Al-5%Zn-1%Mg alloy.-2. electronmicroscopic investigations. Z. Metallk 1980; 71: 231- 238. [22] Y. Liu, D.M. Jiang, B.Q. Li, T. Ying, J. Hu. Heating aging behavior of Al–8.35 Zn–2.5 Mg–2.25 Cu alloy. Mater. Des 2014; 60; 116-124.

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Table Captions Table 1. Nominal chemical composition of the investigated Al-Zn-Mg-Cu alloy (wt.%).

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Figure Captions Fig.1. SEM images of the specimens after 90% rolling reduction at 420 °C (a) 1#, (b) 2#, (c) 3#. Fig.2. XRD patterns of samples after rolling. Fig.3. Microstructures of the center of cross-sectional part of the three kinds of alloys along the rolling direction at 90% deformation after T6 heat treatment (a) 1#, (b) 2#, (c) 3#. Fig.4. TEM graphs of the specimens after T6 heat treatment (a) 1#, (b) 2#, (c) 3#, (d) dark field of (a), (e) is the diffraction patterns of (b), (f) dark field picture of (c). Fig.5. TEM microstructures of alloys after T6 heat treatment (a) 1#, (b) 2#, (c) 3#. Fig.6. Stress versus strain curves of 1#, 2#, 3# alloys in T6 condition. Fig.7. Fracture morphologies of tensile samples (a) 1#, (b) 2#, (c) 3#.

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Table 1

Alloy 1# 2# 3#

Zn 7.82 7.9 7.8

Mg 1.57 1.58 1.59

Cu 0.003 0.768 1.64

Zr 0.139 0.142 0.14

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Si 0.031 0.035 0.026

Fe 0.068 0.071 0.072

Al Bal. Bal. Bal.

Figure 1 (a)

(b)

(c)

13

Figure 2 Al(110)

Al(200)

10000

Al(311) 

Al(222)

 η T

Relative Itensity

Al 7 Cu 2 Fe

 

 









Al(220) 









500

3#

2#

1#

0

20

30

40

50

60

2 (deg.) /Cuka

14

70

80

90

Figure 3 (a)

(b)

(c)

ND RD

15

Figure 4 (a)

(d)

(b)

(e) η’

(c)

(f)

16

Figure 5 (a)

(b)

(c)

17

Figure 6 600

1# 2# 3#

Strength, MPa

500 400 300 200 100 0 0.00

0.05

0.10

Enlongation

18

0.15

0.20

Figure 7 (a)

(b)

(c)

19