Effect of retrogression treatments on microstructure, hardness and corrosion behaviors of aluminum alloy 7085

Journal of Alloys and Compounds 814 (2020) 152264

Contents lists available at ScienceDirect

Journal of Alloys and Compounds journal homepage: http://www.elsevier.com/locate/jalcom

Effect of retrogression treatments on microstructure, hardness and corrosion behaviors of aluminum alloy 7085 Yichang Wang a, Lingfei Cao a, b, *, Xiaodong Wu a, **, Xin Tong a, Bin Liao a, Guangjie Huang a, Zhengan Wang c a

International Joint Laboratory for Light Alloys (Ministry of Education), College of Materials Science and Engineering, Chongqing University, Chongqing, 400044, China Electron Microscopy Center of Chongqing University, Chongqing, 400044, China c Southwest Aluminum (Group) Corporation of China Ltd, Chongqing, 401326, China b

a r t i c l e i n f o

a b s t r a c t

Article history: Received 30 April 2019 Received in revised form 24 August 2019 Accepted 12 September 2019 Available online 13 September 2019

The microstructure, hardness and corrosion resistance of aluminum alloy 7085 after different retrogression and re-aging (RRA) treatments were investigated. The results showed that the properties of alloy 7085 are sensitive to the temperature and dwell time of the retrogression treatment. Desirable mechanical property and corrosion resistance can be obtained for the alloy pre-aged at 120
C for 24 h, followed by a retrogression at 160
C for 1.5 h and re-aging at 120
C for 24 h. After the optimized RRA treatment, the hardness of alloy 7085 is improved by 10.2% as compared with the peak-aged one, which is mainly due to the distribution of dispersed rod-like h0 precipitates in the matrix; while the improved corrosion resistance is mainly attributed to the discrete and coarse h phase with higher Cu content, and narrow precipitate-free zone about 45e50 nm in width along the grain boundaries. © 2019 Elsevier B.V. All rights reserved.

Keywords: Aluminum alloy 7085 RRA Hardness Intergranular corrosion Electrochemistry

1. Introduction 7xxx series (AleZneMgeCu) aluminum alloys have been widely used in the aerospace industry due to the high specific strength and toughness [1e3]. In recent decades, these alloys are also utilized for drill pipes to improve drill efficiency with less energy consumption [4e6]. Compared with conventional steel materials, the agehardenable AleZneMgeCu alloy shows remarkable strength to weight ratio, high fatigue resistance and damage tolerance [7e10]. In addition, drill pipes made of aluminum alloys can effectively prevent from stress corrosion crack and electrochemistry corrosion in deep wells containing H2S and CO2 [11]. The alloy 7075 (5.1e6.1 Zn, 2.1e2.9 Mg, 1.2e2.0 Cu, 0.18e0.28 Cr, ~0.3 Mn, ~0.2 Ti, ~0.5 Fe, ~0.4 Si and balance Al, in wt.%) is so far the

* Corresponding author. International Joint Laboratory for Light Alloys (Ministry of Education), College of Materials Science and Engineering, Chongqing University, Chongqing, 400044, China. ** Corresponding author. International Joint Laboratory for Light Alloys (Ministry of Education), College of Materials Science and Engineering, Chongqing University, Chongqing 400044, China E-mail addresses: [email protected] (L. Cao), [email protected] (X. Wu). https://doi.org/10.1016/j.jallcom.2019.152264 0925-8388/© 2019 Elsevier B.V. All rights reserved.

most widely used aluminum alloys for drill pipes. AA7075 was registered as an international designation in 1954, and later in 2002, the Zr-containing aluminum alloy 7085 was developed with a higher Zn/Mg ratio and less impurity elements. The addition of Zr may lead to the formation of metastable and coherent Al3Zr particles, which can inhibit the recrystallization of alloy 7085 and improve its strength, thermostability and corrosion resistance [12e16]. Therefore, the alloy 7085 can be expected to have a good combination of corrosion resistance and strength during heat exposure, which benefits its potential application for drill pipes. In order to further enhance the resistance of AleZneMgeCu alloys to the localized corrosion along grain boundaries (GBs) in a chloride solution, e.g., intergranular corrosion (IGC), one effective method is adjusting their aging conditions. It's reported that the microstructure and chemical component of GBs are the main factors on IGC susceptibility. h (MgZn2) phase is the typical grain boundary precipitates (GBPs) in AleZneMgeCu alloys, which is anodic to the matrix and tends to be attacked preferentially in the corrosive solution [17-21]. However, the addition of Cu shifts the potential of GBPs towards noble direction as Cu shows higher electrochemical potential than Zn and Mg. If the Cu-depleted GBPs are continuous, the anodic dissolution of these active phases will

2

Y. Wang et al. / Journal of Alloys and Compounds 814 (2020) 152264

occur, which can accelerate the crack propagation and thus increase the IGC susceptibility [22-25]. That's the reason why AleZneMgeCu alloys in the peak-aged condition (T6) are easily subjected to IGC [18,26-31]. It is an effective way to improve the mechanical properties and corrosion resistance of aluminium alloys by optimizing heat treatments [23,32e36]. T7x duplex overaging treatments can improve the corrosion resistance however with a partial loss of strength [18,28e32], while retrogression and re-aging (RRA) may result in a good combination of strength and corrosion resistance. The RRA treatment basically contains three steps: (1) pre-aging at an under-aged or T6 temper; (2) retrogression for a short time at higher temperatures to partly dissolve the previous precipitates and coarsen the GBPs; (3) re-aging at a treatment similar to T6 temper, leading to the re-precipitation and growth of strengthening precipitates. The alloy aged by a proper RRA treatment should have near peak strength and good corrosion resistance, due to the formation of fine h0 precipitates in the matrix and coarse h precipitates discretely distributed at GBs [17,32e37]. Table 1 summarizes some findings from published works about corrosion behaviors of AleZneMgeCu alloys after different thermal treatments [36e40]. The RRA procedure is proved to enhance the IGC and exfoliation corrosion (EXCO) resistance of alloy 7085 without reducing the strength [35e38]. It's also documented that Cu content in GBPs has a great influence on the corrosion performance of AleZneMgeCu alloys, and a higher aging temperature leads to higher copper content of GBPs [23,35,36]. However, the retrogression at a high temperature for short time (several minutes at 200e260
C) is not desirable to large-sized products because the material is difficult to be heated evenly under these conditions. Therefore, the retrogression should be designed with a longer holding time to make the process industry applicable. In addition, the effect of such retrogression on the microstructural evolution and corrosion behavior, especially IGC in alloy 7085 is not clear yet. Therefore, in this work, the effects of retrogression treatment on microstructures, IGC behavior, hardness and electrochemical properties of alloy 7085 have been investigated. An optimal RRA treatment was proposed to improve the hardness and IGC resistance of alloy 7085 with fine dispersed precipitates in the matrix and discontinuous precipitates at grain boundaries. The variation in microstructures and properties were studied for samples aged by

conventional T6 or various RRA treatments. 2. Experimental procedures 2.1. Material and heat treatments The material used in this study was a commercial 7085 thick plate (7.41 Zn, 1.50 Mg, 1.79 Cu, 0.09 Zr, 0.04 Ti, 0.04 Fe, 0.02 Si and balance Al, in wt.%). The alloy was hot-rolled and provided by the Southwest Aluminum (Group) Corporation of China Ltd. The specimens with a cross-section of 25  20 mm2 and 4 mm in thickness were cut along the STD (short transverse direction) - RD (rolling direction) from the plate, followed by different heat treatments (Fig. 1). The solution treatment temperature was selected based on the differential scanning calorimetry (DSC) results, where 485
C was the melting temperature of eutectics. Therefore, the solution treatment was designed as a two-stage process, i.e., 430
C for 1 h followed by another 1 h at 480
C in an air furnace and water quenched at room temperature. Then, all samples were exposed to RRA treatments with different retrogression procedures. The RRA treatment parameters were mainly selected according to the empirical values from literatures [27,32,34,36] and our preliminary investigation. The peak-aged condition (sample T) was selected for alloy 7085 to obtain the highest hardness during ageing at 90e200
C. The retrogression treatments were chosen for an application potential for large-sized products. Details of RRA treatments are listed in Table 2, where samples A, B and C were preaged and re-aged at 120
C for 24 h with the retrogression treatment at 160
C, 180
C and 200
C, respectively. All samples were quenched in water at room temperature after heat treatments at different stages, and the quench rates of each steps were measured by a thermocouple and marked in Fig. 1. 2.2. Hardness and electrical conductivity testing Samples for hardness and electrical conductivity testing were flat with surface mechanically polished for optimum accuracy of measurement. The Vickers hardness testing was performed in accordance with the standard test method described in ASTM E9217 [41]. A MH-5L hardness tester was used with a load of 1000 g and

Table 1 Summary of corrosion behavior after different thermal conditions of 7xxx series aluminum alloys as reported in the literature. Alloys (wt.%) Aging Conditions

Corrosion Trends

Corrosion Resistance Mechanisms

7A09

High strength with severe IGC.

Preferential dissolution of the anodic Cu-depleted zone along continuous grain boundaries [23]. Coarse and sparely distribution of GBPs with higher Cu content [35].

AA7085

AA7150

7085

7075

Al-6.02Zn e2.31Mg e2.04Cu 7055

T6: 120
C/24 h (WQ) T74: 110
C/6 h þ 160
C/10 h (WQ after every step) RRA: 60
C/24 h þ 190
C/0.5 h þ 120
C/24 h (WQ after every step) DRRA: 120
C/24 h þ 180
C/0.5 h þ 120
C/ 12 h þ 180
C/0.5 h þ 120
C/24 h (WQ after every step) HTTP: 470
C/1 h (ST)/cooling to 445
C/30 min (WQ) þ 120
C/24 h (WQ) NIA: 40
C/190
C (20
C/h) /100
C (20
C/h) (WQ) ST þ 120
C/24 h (WQ)

Improves EXCO resistance with strength loss. Good combination of both corrosion resistance and strength Improves corrosion resistance without sacrificing the strength.

GBPs are coarser and disconnected, and slow heating rate provides time for Cu diffusion to precipitates which mainly affect the corrosion resistance [32]. Coarse and discrete distribution of GBPs with higher Cu content [36].

Improves the IGC and EXCO resistance with strength loss.

The h0 precipitate density is decreased due to the formation of coarse particles during the pre-precipitation process at high temperature [34].

High mechanical performances and comparable corrosion resistance A lower quench rate leads to lower EXCO resistance

The variation in microchemistry or macrostructure of GBs and GPBs [9].

GBPs in the slowly quenched specimens are more electrochemically active and tend to accelerate corrosion [1].

ST: solution treatment; WQ: water quenching; AQ: air quenching. DRRA: dual-retrogression and re-aging; HTPP: high-temperature pre-precipitation; NIA: non-isothermal aging. The EXCO test solution is 4.0 M NaCl þ 0.5 M KNO3 þ 0.1 M HNO3 at 25
C.

Y. Wang et al. / Journal of Alloys and Compounds 814 (2020) 152264

3

Fig. 1. Schematic of heat treatment routes for alloy 7085 samples.

Table 2 Detailed RRA treatments for alloy 7085. Samples

Pre-aging

Retrogression

Re-aging

A1 A2 B1 B2 C1 C2

120
C, 24 h

160
C, 160
C, 180
C, 180
C, 200
C, 200
C,

120
C, 24 h

0.5 h 1.5 h 0.5 h 1.5 h 0.5 h 1.5 h

a dwell time of 15 s. Ten measurements were carried out across each sample, and an average value of Vickers hardness was calculated for each data point with a standard deviation of about 2 HV. The electrical conductivity (EC) testing was measured by using a D60K digital meter based on the standard test method (ASTM E1004-17 [42]). Ten measurements were made for each sample to calculate the average value, which is expressed as a percentage of conductivity of the International Annealed Copper Standard (% IACS). All tests were carried out at room temperature.

Fig. 2. Schematic diagram and parameters of potentiodynamic polarization curve.

different heat-treated alloys. 2.3. Electrochemical testing Potentiodynamic polarization experiments were carried out using a GAMRY reference 600 þ Electrochemical Workstation in a range from 1.4e0.5 V with a scan rate of 1 mV/s. A three-electrode system comprised the surface of the studied alloy, a Pt counter electrode and a saturated calomel reference electrode (SCE). Samples were ground using SiC paper up to 4000 grit and rinsed by ethanol. The standard exposed surface area was 1.0 cm2 and the measurements were performed at 25
C in a solution mixture of 1.0 M NaCl þ 0.01 M H2O2. Three tests were repeated for each sample to ensure the data reproducibility. The corrosion potential (Ecorr) and corrosion rate (icorr) were estimated using a Tafel-type fit showed in Fig. 2 with the application of Gamry software.

2.5. Microstructural analysis The morphology and chemical composition of precipitates in the interior of the grains and at GBs of samples A2, B2, C2 and T were studied by a Transmission Electron Microscope (TEM, JEOL JEM-2100) equipped with energy disperse spectroscopy (EDS). The foils for TEM were prepared using the twin-jet electro-polishing (solution was 75% CH3OH and 25% HNO3) at 30
C after mechanical polishing to 70 mm. The TEM examination was performed with the specimen oriented along [011] zone axis of Al matrix. In order to ensure the accuracy of EDS results, each test is repeated at least five times on GBPs in different positions. Microcracks on the corrosion surface were observed on a Tescan Vega 2 Scanning Electron Microscope (SEM).

2.4. Intergranular corrosion tests 3. Results The IGC test was undertaken in a solution of 1.0 M NaCl þ 0.01 M H2O2 at 35 ± 3
C for 6 h according to the standard GB/T 7998-2005 [43]. After immersion, all the samples were put in a 10% nitric acid solution for about 10 s, followed by an ethanol rinse to remove the corrosion products on the surface. The corroded surface at a macroscopic scale was captured by a digital camera, and then examined using a Zeiss 40 MAT optical microscope (OM). The corrosion depth was measured at different locations of the sample, and the maximum value was used to represent the IGC resistance of

3.1. Hardness and electrical conductivity The hardness and electrical conductivity of alloy 7085 in different aging conditions are shown in Fig. 3. With the increase of retrogression temperature and dwell time, the electrical conductivity gradually increases, but the hardness varies. For the sample T as peak-aged, the hardness is 182 HV, which becomes higher after retrogression at 160
C (A1 and A2) and 180
C (B1 and B2). The

4

Y. Wang et al. / Journal of Alloys and Compounds 814 (2020) 152264

T. 3.3. Intergranular corrosion

Fig. 3. The hardness and electrical conductivity of alloy 7085 in different ageing conditions.

highest values obtained in samples A2 (160
C for 1.5 h) and B1 (180
C for 0.5 h) are about 200 HV. However, the sample C2 retrogression treated at 200
C for 1.5 h shows a great loss of 17.6% in hardness as compared with sample T, indicating that there is a detrimental effect on the strength of alloy 7085 if a retrogression treatment is designed at 200
C for more than 0.5 h.

3.2. Potentiodynamic polarization The potentiodynamic polarization curves for all samples tested are shown in Fig. 4. All curves present a similar form, in spite of the different aging treatments. Such similarity indicates that the corrosion reactions of all samples are coincident. However, the values of corrosion potential (Ecorr) and corrosion current density (icorr) derived from the polarization curves change significantly for different aging conditions. The corresponding corrosion parameters derived from these curves are present in Table 3 to compare the corrosion tendency of each sample. The larger Ecorr and Rp (polarization resistance) values and the smaller icorr value suggest better corrosion resistance. Compare the values listed in Table 3, it is clear that alloy 7085 after A2 treatment shows the best corrosion resistance, whereas the worst corrosion performance occurs in sample

The optical images of corroded surfaces for all samples after immersion in the IGC solution are shown in Fig. 5. Sample T in the peak-aged condition is severely corroded, in which case the corrosion products and network-like cracks can be observed on the surface (Fig. 5a). However, after different RRA treatments, the corroded surfaces become smoother and cleaner than that of sample T (Fig. 5beg). It is obvious that samples A2 (retrogression at 160
C for 1.5 h) and B1 (retrogression at 180
C for 0.5 h) had smaller pits on the surface, indicating that the corrosion resistance of these two samples are much higher than the others, while the pits on samples C1 and C2 (retrogression at 200
C) become larger and deeper again. Further observation of corroded surfaces for samples T, A2, B2 and C2 are shown in Fig. 6. It is clear that at the macro scale, rough surface and serious corrosion exist in sample T (Fig. 6a). The corroded surfaces become smoother for samples A2 and B2 (Fig. 6c and e), but gets uneven again for sample C2 (retrogression at 200
C for 1.5 h, Fig. 6g). At a higher magnification, as shown by the SEM images of the areas marked by red squares, plenty of microcracks exist on the corroded surface of sample T (Fig. 6b), while few microcracks and shallow corrosion pits can be observed for sample A2 (Fig. 6c). Larger corrosion pits filled with corrosion products (with bright contrast) can be obtained in samples B2 and C2 (Fig. 6f and h). Based on the surface characteristics of samples after immersion tests, it is evident that alloy 7085 can achieve better corrosion resistance after RRA treatments, especially with a retrogression at 160
C for 1.5 h. The corroded depth can also be used to evaluate the corrosion resistance of the alloys after different heat treatments [18,44,45]. Fig. 7 presents the corrosion depth of samples exposed to IGC test solution for 6 h. It can be seen from the cross sections that RRA treated samples show significantly better corrosion resistance than peak-aged samples. As shown in Fig. 7a, the pits extend inside the matrix and the maximum corrosion depth is about 104.1 mm in sample T. All samples after RRA treatments have less corrosion depth, among which samples A2 and B1 present the smaller corrosion depths (less than 30 mm) (Fig. 7c and d). Considering the corrosion morphology and values of corrosion depth (Figs. 5e7), it is reasonable to conclude that samples A2 and B1 have relatively better corrosion resistance. If plotting the values of maximum corrosion depth and Ecorr (from Table 3) together with sample aging conditions (Fig. 8), It can be seen that higher Ecorr corresponds to a lower corrosion depth. Therefore, based on the electrochemical results and IGC measurements, it can be derived that the corrosion resistance of all samples increases in the following order: T < C2 < A1 < C1 < B2 < B1 < A2. 3.4. Transmission electron micrographs

Fig. 4. Potentiodynamic polarization curves of alloy 7085 after different RRA treatments.

The TEM bright field images and corresponding diffraction patterns of four representative samples T, A2, B2 and C2 are shown in Fig. 9. For peak-aged sample T (Fig. 9aec), lots of GP zones and a few h0 phases (<5 nm) exist in the matrix, together with some Al3Zr particles as deduced from the diffraction pattern. For sample A2 (Fig. 9def), dense rod-like precipitates are distributed homogenously in the grains after retrogression at 160
C for 1.5 h. Compared to sample T, the average size of h0 precipitates in sample A2 is increased to 6e10 nm, concomitant with a lower number density. Meanwhile, strong diffraction spots of h0 precipitates at 1/3 and 2/3 of {220}Al positions and spots of Al3Zr phases at {110}Al position can be observed in Fig. 9f [36]. This result indicates that

Y. Wang et al. / Journal of Alloys and Compounds 814 (2020) 152264

5

Table 3 Corrosion parameters of alloy 7085 in different aging conditions. Samples

Ecorr (mVSCE)

icorr (mA/cm2)

ba (mV/dec)

-bc (mV/dec)

Rp (Ucm2)

T A1 A2 B1 B2 C1 C2

757 ± 2 727 ± 2 648 ± 2 664 ± 3 690 ± 2 711 ± 2 739 ± 3

346.7 ± 4.1 167.5 ± 2.0 36.6 ± 2.4 76.8 ± 3.2 104.6 ± 2.7 139.3 ± 2.9 211.5 ± 3.3

42.1 35.4 38.4 44.2 53.9 42.8 30.4

170.4 111.8 88.6 101.5 102.4 98.6 114.0

42.2 69.8 318.2 174.1 146.6 93.4 50.1

Ecorr: self-corrosion potential; icorr: corrosion current density; ba: anodic Tafel slope; -bc: cathodic Tafel slope; Rp: polarization resistance, Rp ¼

ba,bc [24]. 2:303  icorr  ðba þ bcÞ

Fig. 5. Optical images of corroded surfaces of different samples after immersion: (a) T; (b) A1; (c) A2; (d) B1; (e) B2; (f) C1; (g) C2.

both h0 precipitates and Al3Zr particles co-exist in sample A2, which can contribute to the strengthening of alloys [46]. However, the size of precipitates become larger when the retrogression temperature increases to 180
C for 1.5 h (sample B2, Fig. 9g). Stronger spots corresponding to h0 and h phases appear in SAD pattern of sample B2, indicating the coarsening of h0 precipitates and the formation of h phase (Fig. 9i). For sample C2, the spots corresponding to h phases become stronger, which suggests that the number and size of h phase are increased after retrogression at 200
C for 1.5 h (Fig. 9j-l). In addition, the morphology of GBPs also changes obviously after different aging treatments. In the peak-aged sample T (Fig. 9b), continuous GBPs are less than 20 nm in width, and there are no precipitate free zone (PFZ) can be observed clearly. However in sample A2, the h precipitates at GBs become coarser and discontinuous (size of 20e30 nm, interparticle spacing of 30e50 nm), together with a narrow PFZ (45e50 nm in width), which should improve the corrosion resistance (Fig. 9e). As for samples B2 and C2, the PFZs are broadened to 60e70 nm and 80e90 nm in width, respectively. Meanwhile, the GBPs become much coarser in samples B2 and C2 (size of 30e45 nm for B2 and size of 35e65 nm for C2), as shown in Fig. 9h and k. The GBPs tend

to gather gradually and become continuous again in sample C2, because the increase of retrogression temperature can accelerate the phase growth and aggregation. The chemical compositions of GBPs are also varied after different aging treatments. As indicated from EDS analysis, the contents of Zn, Mg and Cu in T sample are 9.85 at.%, 4.74 at.% and 1.05 at.%, respectively. From sample A2 to C2, the content of Zn is increased slightly from 9.85 at.% to 10.59 at.%, while Mg is raised from 4.74 at.% to 4.86 at.%. However, Cu content seems to be enhanced significantly at higher retrogression temperatures. The Cu content of GBPs in samples T, A2, B2 and C2 are 1.02 at.%, 2.08 at.%, 2.45 at.% and 3.04 at.%, respectively. The Cu content of sample C2 is increased by 200% as compared with the peak-aged sample T. With the increase of retrogression temperature, the atoms prefer to diffuse to the high-energy GBs, which results in the increase of element content in GBPs [35,36]. This is also the reason why FPZ becomes wider from sample T to C2. Based on the TEM results mentioned above, it seems that sample A2 (retrogression at 160
C for 1.5 h) shows most desirable microstructure characteristics with uniform rod-like h0 precipitates (6e10 nm) in grains and discrete h phases with narrow PFZ along

6

Y. Wang et al. / Journal of Alloys and Compounds 814 (2020) 152264

Fig. 8. Effect of retrogression treatment on the values of Ecorr and maximum corrosion depth in IGC test.

GBs, which can be beneficial to improve mechanical property and corrosion resistance.

4. Discussion

Fig. 6. Macro digital photos (a, c, e, g) and SEM images (b, d, f, h) of corroded surfaces for sample (a, b) T; (c, d) A2; (e, f) B2 and (g, h) C2.

The results above reveal that retrogression during RRA treatment has significant effects on the microstructure, hardness and corrosion resistance of alloy 7085. Precipitation strengthening is one most effective strengthening method in 7xxx aluminum alloys, and their precipitation sequences are generally listed as follows: supersaturated solid solution (SSS) / coherent GP zones / semicoherent h0 / incoherent stable h [9e12]. The formation and coarsening of precipitates in grain boundaries is prior to that within the grain, due to the strong preferential precipitation at GBs during aging treatment [17]. At the early stage of aging treatment, the

Fig. 7. Micrographs of cross section (perpendicular to rolling direction) for the samples subjected to IGC test: (a) T; (b) A1; (c) A2; (d) B1; (e) B2; (f) C1; (g) C2.

Y. Wang et al. / Journal of Alloys and Compounds 814 (2020) 152264

7

Fig. 9. TEM bright field images and corresponding diffraction patterns of samples: (aec) T; (def) A2; (gei) B2; (jel) C2.

precipitates in grains are small and coherent with the matrix, which can be cut by the dislocation during deformation and improve the strength of the alloy. As the aging process continues, the size of precipitates become larger and the coherent relationship may be destroyed (formation of h phase). During the over-aging step, the strength of the alloy is reduced due to coarsening (Ostwald ripening) of the precipitated phases [8e10]. The schematic of microstructure evolution during RRA treatment is shown in Fig. 10. Generally, the GP zones and some h0 phases appear and grow at the pre-aging stage. During the retrogression treatment, the GP zones re-dissolve into the matrix and some of h0 phases continue growing [34]. Meanwhile, the h phase at GBs become coarser and discontinuous. After the re-aging stage, the h0 phases re-precipitate again and the hardness of alloy gets back to a higher level [45-47]. The temperature and dwell time of retrogression is critical, as higher temperature allows Cu to diffuse

and enter the precipitation process. However, if the retrogression treatment takes an excessive long time or high temperature, the dissolved phases will re-precipitate again and rapidly coarsen, which will contribute to the deterioration of properties. Samples A2 and B1 exhibit higher hardness, which is related to the formation and limited coarsening of rod-like semi-coherent h0 phase in the grain with proper retrogression temperature and dwell time [48,49]. Nevertheless, the coherent metastable GP zones and a few h0 phase in the matrix leads to the lower hardness of sample T. The hardness decreases sharply when the retrogression temperature increases to 200
C for 1.5 h. This phenomenon attributes to the excessive coarsening of h0 phase and its transformation to equilibrium h phase. The electrical conductivity can reflect the precipitation behavior during heat treatment in alloy 7085. The electrical conductivity of the alloy can be calculated by Eq. (1) [50].

8

Y. Wang et al. / Journal of Alloys and Compounds 814 (2020) 152264

Fig. 10. The microstructure evolution during RRA treatment: (a) pre-ageing; (b) retrogression; (c) re-ageing.

s¼

Ne2 1 , 2m r

(1)

Where N is the number of free electrons involved in conducting electricity, e is the electronic charge, m is the mass of electron, and r is the scattering probability of electron waves which is also called the electric resistivity. From Eq. (1), it is indicated that the electrical conductivity of the alloy is mainly affected by the scattering probability of electron waves [51]. The imperfect structures such as atomic vibration, impurities and defects in actual metal crystals can scatter electron waves to affect the conductivity. The electric resistivity r can be figured up with these parameters shown in Eq. (2) [50].

r ¼ r0 þ DrSH þ DrP þ DrV þ DrD þ DrGB

(2)

The subscripts SH, P, V, D, and GB represent the effect on electrical resistance of solution atoms, precipitates, vacancy, dislocation and grain boundaries, respectively. And in alloys with multicomponents, the order of influence on the resistivity is: DrSH ＞ ΔrP＞DrGB DrD ＞ΔrV＞[50]. The precipitation behavior plays an important role in the change of electrical conductivity. As a result, the electrical conductivity of the alloy 7085 increases due to the formation and coarsening of h0 and h precipitates especially during the retrogression and re-aging processes, and it is in accordance with the morphology of precipitates in Fig. 9. Besides, the addition of Zr in alloy 7085 also plays an important role in the mechanical and corrosion properties. AleZr system can form as fine dispersion precipitates of a stable L12 phase, Al3Zr, which can improve the hardness and strength by precipitation strengthening [16,50]. Coherent Al3Zr dispersoid of alloy 7085 can effectively prevent the migration of subgrain boundaries during heat treatments so as to restrain the recrystallization process. However, the solubility of Zr in Al matrix is extremely low at regular aging temperature (<200
C). Al3Zr dispersions formed during homogenization treatment (>400
C) basically remain stable in the subsequent aging treatment [16], so that the change of retrogression treatments seem to have no significant effect on the morphology and distribution of Al3Zr. Therefore, sample A2 was selected as a representative to observe the distribution and size of the Al3Zr particles. The discussion on Al3Zr is introduced here mainly to explain why the hardness and corrosion resistance of alloy 7085 are higher than those of other AleZneMgeCu alloys without Zr mentioned in other literatures [17,18], such as alloy 7075 (traditional aluminum alloy for oil drill pipes). A typical TEM dark field image of Al3Zr precipitates for sample A2 with a L12 crystal structure [47], as marked in the SAED pattern,

Fig. 11. Typical TEM dark field images of Al3Zr in sample A2.

is presented in Fig. 11. To evaluate the influence of Al3Zr precipitates on recrystallization, the Gibbs-Thomson equation is described in Eq. (3) [16,52e54], where RC is the minimum critical radius to determine whether subgrains can be formed as recrystallized grains.

RC ¼

4gGB PD  PZ

(3)

The gGB is the specific grain boundary energy (about 0.3 J/m2 [15]), PD is the stored deformation energy (alloys subjected to the same deformation process will have similar PD values), and PZ is Zener pinning pressure which can represent the recrystallization resistance. PZ can be estimated by Eq. (4).

PZ ¼

3gGB ðf =rÞ 2

(4)

Where f is the volume fraction and r is the average radius of second phase. Eq. (4) makes clear that higher f/r value can achieve higher PZ, which suggests that the critical radius RC is harder to be achieved. In other words, it means that a high volume fraction of small precipitates can retard the growth of recrystallized grains. So after

Y. Wang et al. / Journal of Alloys and Compounds 814 (2020) 152264

adding Zr element into the alloy, there are plenty of Al3Zr dispersoids in the matrix, which can display a markedly high resistance to recrystallization. From Fig. 11, it can be figured out that the average radius of the Al3(Er,Zr) particles is 10.6 nm, the number density (NV) is 770 mm3 and the volume fraction is 0.384%. The value of f/r in the tested alloy 7085 is 0.362 mm1, which is comparable with the results of previous researches and can effectively inhibit the recrystallization process during heat treatments [15,16].As a result, a large number of deformation structures will remain in the alloy during the subsequent heat treatment after deformation, which leads to relative high hardness and strength. Meanwhile, the inhibition of recrystallization by Al3Zr precipitates in alloy 7085 can improve the corrosion resistance because the intergranular corrosion tends to propagate along high-angle boundaries of recrystallized grains. In addition, due to the lower diffusivity of Zr atoms in the Al-based matrix over a wide temperature range, the Al3Zr particles are stable at high temperatures, which can also increase the thermally stable of the Al alloy [55,56]. Until now, IGC of 7xxx aluminum alloys has been extensively investigated, and anodic dissolution is one of the typical theories on the corrosion mechanisms [25]. On the basis of the anodic dissolution theory, there are two methods that can improve the IGC resistance of 7xxx alloys: one is diminishing the potential difference between GBPs and matrix; the other one is reducing corrosion propagation rate along the grain boundary by increasing the size and spacing of GBPs [33]. The corrosion behaviors in 7xxx alloys is strongly affected by the chemical components and morphology of h phases at GBs. Mg and Zn exhibit more negative electrode potential compared to Al matrix, and that can result in anodic dissolution and crack initiation [28,56,57]. However, the potential of GBPs was enhanced due to the higher electrochemical potential of Cu than that of Zn and Mg. Therefore, the severe corrosion resulted from pitting and IGC for sample T in peak-aged condition are related to the existence of continuous anodic h phase with lower Cu content at GBs and a high propagation rate along recrystallization boundaries. The samples A2 and group B have better corrosion resistance due to the coarse and discontinuous h phase with higher Cu content at GBs. However, the GBPs grow up quickly and become continuous again when the regression temperature is increased to 200
C, which results in the reduction of corrosion resistance of samples C (especially for sample C2). Such phenomenon is in agreement with the results of the potentiodynamic polarization experiment. The growth and coarsening of precipitates along the grain boundary result in the formation of PFZs, which can also affect the corrosion resistance of aluminum alloys [28,39,55]. The PFZ is the solute depleted zone adjacent to the grain boundary, which can restrain the anodic dissolution progressing, and the width of PFZ may broaden by increasing aging temperature and time. After RRA treatment, the decrease of corrosion propagation rate is due to inactive discrete GBPs and formation of PFZs with proper width (about 48.7 nm for sample A2). When the PFZ becomes too wide, it will generate localized attack along the grain boundaries and lead to IGC due to the local stress concentration along the interface of PFZ [58]. To sum up, the alloy 7085 can obtain a good combination of strength and corrosion resistance after RRR treatments with the retrogression at 160
C for 1.5 h (A2) or at 180
C for 0.5 h (B1). Such properties improvement is mainly attributed to the dispersed h0 phase in grains and the coarse discontinuous h phase along grain boundaries after proper RRA processes. The RRA treatment with retrogression at 160
C for 1.5 h is more applicable for industrial processing of thick plates because it allows enough time for the material to be heated uniformly with homogeneous microstructure.

9

5. Conclusions (1) The hardness, electrical conductivity and corrosion resistance of alloy 7085 are sensitive to the retrogression processing. The RRA treatment consists of a pre-aging at 120
C for 24 h, retrogression at 160
C for 1.5 h and re-aging at 120
C for 24 h is verified to obtain a good combination of properties for alloy 7085. (2) The hardness of sample A2 (retrogression at 160
C for 1.5 h) is improved by 10.2% as compared with that of peak-aged sample T, which is mainly due to the distribution of dispersed rod-like h0 precipitates in the matrix. When the retrogression temperature is increased higher than 180
C, the hardness decreases due to the excessively coarsening of h0 and h precipitates. (3) After a proper RRA treatment (A2), the corrosion resistance of alloy 7085 in chloride ionic solution can be greatly improved, and the corrosion depth decreases by more than 70% in contrast with that in peak aged condition. Such difference is due to the formation of coarse and discrete h phases with higher Cu content along grain boundaries after retrogression. Acknowledgements This work was supported by the National Key Research and Development Program of China (2016YFB0300901), the National Natural Science Foundation of China (51871033),Chongqing Research Program of Basic Research and Frontier Technology (No. cstc2017jcyjAX0245) and Venture and Innovation Support Program for Chongqing Overseas Returnees (cx2018002). References [1] S.D. Liu, B. Chen, C.B. Li, Y. Dai, Y.L. Deng, X.M. Zhang, Mechanism of low exfoliation corrosion resistance due to slow quenching in high strength aluminium alloy, Corros. Sci. 91 (2015) 203e212. [2] S.P. Knight, M. Salagaras, A.R. Trueman, The study of intergranular corrosion in aircraft aluminium alloys using X-ray tomography, Corros. Sci. 53 (2011) 727e734. [3] M.A. Krishnan, V.S. Raja, Development of high strength AA 7010 aluminum alloy resistant to environmentally assisted cracking, Corros. Sci. 109 (2016) 94e100. [4] J. Liang, J.H. Sun, X.M. Li, Y.Q. Zhang, L. Peng, Development and application of aluminum alloy drill rod in geologic drilling, Procedia Engineering 73 (2014) 84e90. [5] C. Feng, W.B. Shou, H.Q. Liu, D.Q. Yi, Y.R. Feng, Microstructure and mechanical properties of high strength AleZneMgeCu alloys used for oil drill pipes, Trans. Nonferrous Metals Soc. China 25 (2015) 3515e3522. [6] J.C. Liu, M.J. Qin, Q.L. Zhao, L. Chen, P.F. Liu, J.G. Gao, Fatigue performances of the cracked aluminum-alloy pipe repaired with a shaped CFRP patch, ThinWalled Struct. 111 (2017) 155e164. [7] M. Zhou, Y.C. Lin, J. Deng, Y.Q. Jiang, Hot tensile deformation behaviors and constitutive model of an AleZneMgeCu alloy, Mater. Des. 59 (2014) 141e150. [8] D.M. Jiang, Y. Liu, S. Liang, W.L. Xie, The effects of non-isothermal aging on the strength and corrosion behavior of AlZnMgCu alloy, J. Alloy. Comp. 681 (2016) 57e65. [9] J.T. Jiang, W.Q. Xiao, L. Yang, W.Z. Shao, S.J. Yuan, L. Zhen, Ageing behavior and stress corrosion cracking resistance of a non-isothermally aged AleZneMgeCu alloy, Mater. Sci. Eng. A 605 (2014) 167e175. [10] Y. Liu, S. Liang, D.M. Jiang, Influence of repetitious non-isothermal aging on microstructure and strength of Al-Zn-Mg-Cu alloy, J. Alloy. Comp. 689 (2016) 632e640. [11] S.L. Lv, F.Q. Luo, J. Zhou, Prospects of aluminum alloy drill pope in Tarim oilfield, Petroleum Drilling Techniques 37 (2009) 74e77. [12] H.C. Fang, H. Chao, K.H. Chen, Effect of recrystallization on intergranular fracture and corrosion of AleZneMgeCueZr alloy, J. Alloy. Comp. 622 (2015) 166e173. [13] X. Tong, G.Q. You, Y.C. Wang, H. Wu, W.L. Liu, P.Q. Li, W. Guo, Effect of ultrasonic treatment on segregation and mechanical properties of as-cast MgeGd binary alloys, Mater. Sci. Eng. A 731 (2018) 44e53. [14] J. Liu, P. Yao, N.Q. Zhao, C.S. Shi, H.J. Li, X. Li, D.S. Xi, S. Yang, Effect of minor Sc and Zr on recrystallization behavior and mechanical properties of novel

10

Y. Wang et al. / Journal of Alloys and Compounds 814 (2020) 152264

AleZneMgeCu alloys, J. Alloy. Comp. 657 (2016) 717e725. [15] H. Wu, S.P. Wen, H. Huang, B.L. Li, X.L. Wu, K.Y. Gao, W. Wang, Z.R. Nie, Effects of homogenization on precipitation of Al3(Er,Zr) particles and recrystallization behavior in a new type Al-Zn-Mg-Er-Zr alloy, Mater. Sci. Eng. A 689 (2017) 313e322. [16] Z.Y. Guo, G. Zhao, X. Chen, Effects of two-step homogenization on precipitation behavior of Al3Zr dispersoids and recrystallization resistance in 7150 aluminum alloy, Mater. Char. 102 (2015) 122e130. [17] W.T. Huo, J.J. Hu, H.H. Cao, Y. Du, W. Zhang, Y.S. Zhang, Simultaneously enhanced mechanical strength and inter-granular corrosion resistance in high strength 7075 Al alloy, J. Alloy. Comp. 781 (2019) 680e688. [18] F. Andreatta, H. Terryn, J.H.W. de Wit, Corrosion behaviour of different tempers of AA7075 aluminium alloy, Electrochim. Acta 49 (2004) 2851e2862. [19] J.F. Chen, X.F. Zhang, L.C. Zou, Y. Yu, Q. Li, Effect of precipitate state on the stress corrosion behavior of 7050 aluminum alloy, Mater. Char. 114 (2016) 1e8. [20] S.Y. Chen, K.H. Chen, P.X. Dong, S.P. Ye, L.P. Huang, Effect of recrystallization and heat treatment on strength and SCC of an AleZneMgeCu alloy, J. Alloy. Comp. 581 (2013) 705e709. [21] A.C. Umamaheshwer Rao, V. Vasu, M. Govindaraju, K.V. Sai Srinadh, Stress corrosion cracking behaviour of 7xxx aluminum alloys: a literature review, Trans. Nonferrous Metals Soc. China 26 (2016) 1447e1471. [22] F.X. Song, X.M. Zhang, S.D. Liu, Q. Tan, D.F. Li, The effect of quench rate and overageing temper on the corrosion behaviour of AA7050, Corros. Sci. 78 (2014) 276e286. [23] Y.W. Sun, Q.L. Pan, Y.Q. Sun, W.Y. Wang, Z.Q. Huang, X.D. Wang, Q. Hu, Localized corrosion behavior associated with Al7Cu2Fe intermetallic in Al-ZnMg-Cu-Zr alloy, J. Alloy. Comp. 783 (2019) 329e340. [24] S. Li, H.G. Dong, L. Shi, P. Li, F. Ye, Corrosion behavior and mechanical properties of Al-Zn-Mg aluminum alloy weld, Corros. Sci. 123 (2017) 243e255. [25] X.Y. Peng, Q. Guo, X.P. Liang, Y. Deng, Y. Gu, G.F. Xu, Z.M. Yin, Mechanical properties, corrosion behavior and microstructures of a non-isothermal ageing treated Al-Zn-Mg-Cu alloy, Mater. Sci. Eng. A 688 (2017) 146e154. [26] J. Wloka, T. Hack, S. Virtanen, Influence of temper and surface condition on the exfoliation behaviour of high strength AleZneMgeCu alloys, Corros. Sci. 49 (2007) 1437e1449. [27] G.S. Peng, K.H. Chen, S.Y. Chen, H. Fang, Influence of dual retrogression and reaging temper on microstructure, strength and exfoliation corrosion behavior of AleZneMgeCu alloy, Trans. Nonferrous Metals Soc. China 22 (2012) 803e809. [28] S.P. Knight, N. Birbilis, B.C. Muddle, A.R. Trueman, S.P. Lynch, Correlations between intergranular stress corrosion cracking, grain-boundary microchemistry, and grain-boundary electrochemistry for AleZneMgeCu alloys, Corros. Sci. 52 (2010) 4073e4080. [29] S.P. Knight, K. Pohl, N.J.H. Holroyd, N. Birbilis, P.A. Rometsch, B.C. Muddle, R. Goswami, S.P. Lynch, Some effects of alloy composition on stress corrosion cracking in AleZneMgeCu alloys, Corros. Sci. 98 (2015) 50e62. [30] H.Z. Li, S.C. Yao, X.P. Liang, Y.H. Chen, C. Liu, L. Huang, Grain boundary preprecipitation and its contribution to enhancement of corrosion resistance of AleZneMg alloy, Trans. Nonferrous Metals Soc. China 26 (2016) 2523e2531. [31] X.Y. Sun, B. Zhang, H.Q. Lin, Y. Zhou, L. Sun, J.Q. Wang, E.H. Han, W. Ke, Correlations between stress corrosion cracking susceptibility and grain boundary microstructures for an AleZneMg alloy, Corros. Sci. 77 (2013) 103e112. [32] D.K. Xu, N. Birbilis, P.A. Rometsch, The effect of pre-ageing temperature and retrogression heating rate on the strength and corrosion behaviour of AA7150, Corros. Sci. 54 (2012) 17e25. [33] B.M. Cina, Reducing the Susceptibility of Alloys, Particularly Aluminum, to Stress Corrosion Cracking, US, 1974, p. 3856584. [34] J.F. Li, N. Birbilis, C.X. Li, Z.Q. Jia, B. Cai, Z.Q. Zheng, Influence of retrogression temperature and time on the mechanical properties and exfoliation corrosion behavior of aluminium alloy AA7150, Mater. Char. 60 (2009) 1334e1341. [35] S.Y. Chen, K.H. Chen, G.S. Peng, L. Jia, P.X. Dong, Effect of heat treatment on strength, exfoliation corrosion and electrochemical behavior of 7085 aluminum alloy, Mater. Des. 35 (2012) 93e98.

[36] S.Y. Chen, K.H. Chen, P.X. Dong, S.P. Ye, L.P. Huang, Effect of heat treatment on stress corrosion cracking, fracture toughness and strength of 7085 aluminum alloy, Trans. Nonferrous Metals Soc. China 24 (2014) 2320e2325. [37] A.F. Oliveira, M.C. de Barros, K.R. Cardoso, D.N. Travessa, The effect of RRA on the strength and SCC resistance on AA7050 and AA7150 aluminium alloys, Mater. Sci. Eng. A 379 (2004) 321e326. [38] S.Y. Chen, K.H. Chen, G.S. Peng, X. Liang, X.H. Chen, Effect of quenching rate on microstructure and stress corrosion cracking of 7085 aluminum alloy, Trans. Nonferrous Metals Soc. China 22 (2012) 47e52. [39] S.Y. Chen, K.H. Chen, L. Jia, G.S. Peng, Effect of hot deformation conditions on grain structure and properties of 7085 aluminum alloy, Trans. Nonferrous Metals Soc. China 23 (2013) 329e334. [40] B.H. Nie, P.Y. Liu, T.T. Zhou, Effect of compositions on the quenching sensitivity of 7050 and 7085 alloys, Mater. Sci. Eng. A 667 (2016) 106e114. [41] Standard Test Methods for Vickers Hardness and Knoop Hardness of Metallic Materials, 2016. ASTM E92-17. [42] Standard Test Method for Determining Electrical Conductivity Using the Electromagnetic (Eddy Current) Method, 2017. ASTM E1004-17. [43] GB 7998-2005, National Standard of China. [44] U. Trdan, J. Grum, Evaluation of corrosion resistance of AA6082-T651 aluminium alloy after laser shock peening by means of cyclic polarisation and ElS methods, Corros. Sci. 59 (2012) 324e333. [45] Y.C. Lin, J.L. Zhang, G. Liu, Y.J. Liang, Effects of pre-treatments on aging precipitates and corrosion resistance of a creep-aged AleZneMgeCu alloy, Mater. Des. 83 (2015) 866e875. [46] J.H. Li, M. Wiessner, M. Albu, S. Wurster, B. Sartory, F. Hofer, P. Schumacher, Correlative characterization of primary Al3(Sc,Zr) phase in an AleZneMg based alloy, Mater. Char. 102 (2015) 62e70. [47] Z.F. Guo, G. Zhao, X. Chen, Effects of homogenization treatment on recrystallization behavior of 7150 aluminum sheet during post-rolling annealing, Mater. Char. 114 (2016) 79e87. [48] G. Sha, A. Cerezo, Early-stage precipitation in AleZneMgeCu alloy (7050), Acta Mater. 52 (2004) 4503e4516. [49] C. Lei, H. Yang, H. Li, N. Shi, L.H. Zhan, Dependences of microstructures and properties on initial tempers of creep aged 7050 aluminum alloy, J. Mater. Process. Technol. 239 (2017) 125e132. [50] L.F. Mondolfo, Aluminum Alloys : Structure and Properties, Butter Worths, 1976. [51] X. Tong, G.Q. You, Y.H. Ding, H.S. Xue, Y.C. Wang, W. Guo, Effect of grain size on low-temperature electrical resistivity and thermal conductivity of pure magnesium, Mater. Lett. 229 (2018) 261e264. [52] Y. Deng, Z.M. Yin, Q.L. Pan, G.F Xu, Y.L Duan, Y.J Wang, Nano-structure evolution of secondary Al3(Sc1xZrx) particles during superplastic deformation and their effects on deformation mechanism in Al-Zn-Mg alloys, J. Alloy. Comp. 695 (2017) 142e153. [53] H.Y. Li, Z.H. Gao, H. Yin, H.F. Jiang, X.J. Su, J. Bin, Effects of Er and Zr additions on precipitation and recrystallization of pure aluminum, Scr. Mater. 68 (2013) 59e62. [54] X. Tong, G.Q. You, Y. Liu, S.Y. Long, Q. Liu, Effect of C2H2 as a novel gas inoculant on the microstructure and mechanical properties of as-cast AM60B magnesium alloy, J. Mater. Process. Technol. 271 (2019) 271e283. [55] S.P. Wen, K.Y. Gao, Y. Li, H. Huang, Z.R. Nie, Synergetic effect of Er and Zr on the precipitation hardening of AleEreZr alloy, Scr. Mater. 65 (2011) 592e595. [56] T.M.D. Najjar, T.J. Warner, Influence of critical surface defects and localized competition between anodic dissolution and hydrogen effects during stress corrosion cracking of a 7050 aluminium alloy, Mater. Sci. Eng. A 238 (1997) 293e302. [57] W.C. Yang, S.X. Ji, Q. Zhang, M.P. Wang, Investigation of mechanical and corrosion properties of an AleZneMgeCu alloy under various ageing conditions and interface analysis of h0 precipitate, Mater. Des. 85 (2015) 752e761. [58] T. Ogura, S. Hirosawa, T. Sato, Quantitative characterization of precipitate free zones in AleZneMg(eAg) alloys by microchemical analysis and nanoindentation measurement, Sci. Technol. Adv. Mater. 5 (2004) 491e496.