The mechanism of comprehensive properties enhancement in Al–Zn–Mg–Cu alloy via novel thermomechanical treatment

Journal Pre-proof The mechanism of comprehensive properties enhancement in Al–Zn–Mg–Cu alloy via novel thermomechanical treatment Zhiguo Chen, Zhengui Yuan, Jieke Ren PII:

S0925-8388(20)30809-4

DOI:

https://doi.org/10.1016/j.jallcom.2020.154446

Reference:

JALCOM 154446

To appear in:

Journal of Alloys and Compounds

Received Date: 5 December 2019 Revised Date:

18 February 2020

Accepted Date: 20 February 2020

Please cite this article as: Z. Chen, Z. Yuan, J. Ren, The mechanism of comprehensive properties enhancement in Al–Zn–Mg–Cu alloy via novel thermomechanical treatment, Journal of Alloys and Compounds (2020), doi: https://doi.org/10.1016/j.jallcom.2020.154446. This is a PDF file of an article that has undergone enhancements after acceptance, such as the addition of a cover page and metadata, and formatting for readability, but it is not yet the definitive version of record. This version will undergo additional copyediting, typesetting and review before it is published in its final form, but we are providing this version to give early visibility of the article. Please note that, during the production process, errors may be discovered which could affect the content, and all legal disclaimers that apply to the journal pertain. © 2020 Published by Elsevier B.V.

Author contributions: Zhiguo Chen contributed to conceptualization, research design, data analysis, writing and supervision. Zhengui Yuan performed research, analyzed data and wrote the draft. Jieke Ren analyzed data. All authors contributed to refining the ideas, carrying out additional analyses and finalizing this paper.

The mechanism of comprehensive properties enhancement in Al-Zn-Mg-Cu alloy via novel thermomechanical treatment Zhiguo Chena,b,*, Zhengui Yuana , Jieke Renc a

School of Materials Science and Engineering, Central South University, Changsha 410083, P.R.China

b

Department of Materials Engineering, Hunan University of Humanities, Science and Technology, Loudi 417000, P.R.

China c

School of Materials Science and Engineering, Zhejiang University, Hangzhou 310027, P.R.China

Abstract The influence of novel thermomechanical treatment (NTMT) on the microstructure and mechanical behavior was studied by methods such as tensile test, slow strain rate test (SSRT), X-ray diffraction (XRD), transmission electron microscope (TEM), electron back-scattered diffraction (EBSD) and atom probe tomography (APT). Results show that the comprehensive properties can be significantly improved by NTMT in Al-Zn-Mg-Cu alloy; the alloy treated by NTMT not only obtains excellent combination of strength and ductility, but also has good stress corrosion resistance. Analysis shows that the synergy effects of co-clusters, distribution of dislocations, texture configuration and nanoprecipitates contribute to the improvement of the strength. Ductility is maintained mainly due to precipitation characteristics and slight recovery of dislocations. Meanwhile the formation of discontinuously distributed grain boundary precipitates, dislocations distribution and a high proportion of low-angle grain boundaries work together to improve the corrosion resistance of the alloy. Keyword: Al-Zn-Mg-Cu alloy; thermomechanical treatment; stress corrosion resistance; co-cluster 1. Introduction High-strength 7000 series aluminum alloys have been widely used as critical structure materials in aircraft, marine and nuclear industries due to their excellent comprehensive properties, such as high strength to weight ratio, high fatigue resistance and ease of fabrication [1-3]. However, these alloys are prone to localized corrosion, like pitting, inter-granular and stress corrosion cracking (SCC) [4-5], which limits its practical applications. Since the strength and the corrosion resistance of the alloy are both important in the application, it is essential to achieve a good combination of strength and corrosion resistance. The optimization of alloy composition via microalloying is an effective means by which the alloy properties may be enhanced [6-8]. Typical alloy design approaches select and applies alloying elements at levels near their solid-solubility limit. However, the supersaturated solid solution is prone to decomposition and excessive constituents result in poor performance. Reasonable choice of heat treatment can also improve the alloy properties. It is well-documented that the Al-Zn-Mg-Cu alloy in the T6 state has high strength, but has high stress corrosion sensitivity [9]. The SCC resistance can be greatly improved by over-aging, namely the T7x temper [10-11]. Unfortunately, the strength is 1

inevitably reduced by 10 to 15% compared with the T6 state. Cina et al. [12] developed a multi-step aging process, known as retrogression and re-aging (RRA), which increased the corrosion resistance while maintaining a strength level similar to that of T6 temper. But it may be unsuitable for large-profiled materials due to the fact that a short retrogression time is necessary during Retrogression. The comprehensive properties of 7xxx alloys could also be modified via severe plastic deformation (SPD) like equal channel angular pressing (ECAP), accumulative roll bonding (ARB) and high-pressure torsion (HPT) [13-16]. However, the difficulties in producing large-scale parts limit their practical application. Numerous studies have shown that thermomechanical treatment consisting of “solution treating-cold-rolling-ageing” can effectively increase the dislocation density and refine grain size, thereby increasing strength and providing good ductility for aluminum alloy [17-20]. The more prominent feature of TMT is its great feasibility in industrial applications. Lin et al. [21] achieved a yield strength of 301 MPa and a conductivity of 58.9% IACS for Al-Mg-Si alloy by the process of “pre-aging-cold deformation-re-aging”. Huo et al. [22] combined the warm rolling and continuous rolling treatment to successfully refine the grain of 7075 aluminum alloy to less than 10µm, improving mechanical performance of the aluminum alloy. Zuo et al. [23] propose a modified TMP which includes pre-deformation, short time intermediate annealing and final hot rolling to produce ultra fine grained structure, exhibiting good mechanical properties efficiently. It is noteworthy that ultrafine-grained alloys tend to have a high proportion of high-angle grain boundaries, which leads to increased intergranular corrosion and stress corrosion sensitivity [24-25]. One of our recent works [26] has designed a new thermomechanical treatment based on cluster strengthening, which effectively enhanced the tensile mechanical strength of Al-Mg-Si alloy while maintaining good tensile ductility. Given the results above, we propose a novel thermomechanical treatment to achieve better mechanical properties as well as stress corrosion resistance in Al-Zn-Mg-Cu alloy. We discuss the involved mechanisms in detail, especially the strengthening mechanism of NTMT is studied down to the atomic scale. 2. Material and methods The chemical compositions of Al-Zn-Mg-Cu alloy used in this research are shown in Table 1. In the NTMT schedule: the Al-Zn-Mg-Cu alloy sheets with a thickness of 4 mm was solution treated at 475°C for 1h and directly followed by hot rolling. During hot rolling process, the best thickness reduction value is selected to be 30%, which is obtained by one-pass, the temperature is maintained above 450°C, and then rapidly quenched into water at room temperature (We name this process as SSHR). The main function of this process is to obtain a hot-rolled structure in a supersaturated state which will affect the subsequent pre-aging. Later, the alloy was immediately subjected to pre-aging. Next, to create different deformation microstructure, the samples were separately processed at room temperature by procedures of cold rolling (CR), and asymmetrical rolling (ASR) which can feasibly be adapted at an industrial scale. In ASR procedure, the speed ratio between the rolls was 1.3 and 1.6, which is achieved by rolls on one rolling stand with the same radius but different linear speeds. Each pass thickness reduction was maintained to be about 2.5% to avoid temperature ramp. Finally, the samples are treated by the final aging (natural aging or low-temperature artificial aging). 2

Different experimental parameters are listed in Table 2. TABLE 1 Chemical composition of Al-Zn-Mg-Cu alloy used in the experiment (wt.%). Zn

Mg

Cu

Fe

Si

Zr

Mn

Al

6.20

2.35

2.23

0.15

0.12

0.10

0.10

Bal.

TABLE 2 Heat treatment procedures used for Al-Zn-Mg-Cu alloy

Sample

ASR-TMT

Pre-aging

A1A

ASR(10%, r=1.3)

A2A

ASR(10%, r=1.6)

A1N

SSHR

80°C /6h

A2N CR-TMT

Deformation

C1A C1N

80°C /6h

T6

475°C /1h+W.Q.+120°C /24h

RRA

T6+200°C /40min+120°C /24h

100°C /6h

ASR(10%, r=1.3) ASR(10%, r=1.6)

SSHR

Final aging

CR10%

NA 15 d 100°C/6h NA 15 d

* r is the speed ratio between the rolls during ASR. NA is natural aging. W.Q. is water quenching

The tensile properties of samples were assessed at room temperature (25±2°C) and at an initial strain rate of 10-3s-1 on a MTS-858 electronic tensile machine along rolling direction with a constant cross head velocity of 2 mm/min. These uniaxial tensile samples possessed dimensions consistent with ASTM E8/E8M-16a. Three parallel samples were tested for each data point with the magnitudes of the error bars being the standard deviations. Slow strain rate test (SSRT) was carried out to study the SCC resistance in air and in 3.5% NaCl solution (corrosive environment) at room temperature (25±2°C). The long axes of the specimens were to the rolling direction. Dog-bone shaped specimens were tested in a conventional tensile machine with a strain rate of 10−6 s−1 with the gauge sections being completely immersed. The XRD measurements were conducted by Bruker D8 DISCOVER automatic X-ray diffractometer to measure the dislocation density of samples. The micro-crystallite size and micro-strain are obtained according to the Williamson-Hall method [27], which can be expressed by the equation of (B·cosθ)/λ = 1/D + (4ε·sinθ)/λ, where B is the broadening at FWHM, λ=0.154 nm, is the X-ray wavelength, and θ is the Bragg angle (half of the diffraction angle). In order to characterize the texture evolution of the alloy during the thermomechanical treatment, three incomplete pole figures of {111}, {200} and {220} were measured by a Bruker D8 Discover X-ray diffractometer. Based on these incomplete pole figures, the orientation distribution functions (ODFs) were calculated by the series expansion method, at the same time the ghost correction was 3

made by the positivity method. And the volume fractions of different texture components were calculated by the decomposing method according to particle swarm optimization algorithm [28]. The TEM observations were performed on a TECNAI G220 transmission electron microscopy. The samples for TEM observation were prepared by the standard twin-jet electropolishing method performed below -25°C using a solution of methanol and nitric acid (3:1 in volume). EBSD experiments were carried out on Sirion 200 field emission scanning electron microscope controlled by TSLTM software. The samples are mechanically polished and then electrolytically polished. Experimental results were analyzed using OIM software. Atom probe tomography (APT) experiments were carried out on a LEAP 3000 HR instrument at sample temperature of about 20 K under an ultrahigh vacuum condition of ~1.2×10-10 Torr with a pulse fraction of 15% and a pulse repetition rate of 200 kHz.

3. Results 3.1. Tensile properties To obtain an insight into the effect of heat treatment on the mechanical properties of Al-Zn-Mg-Cu alloy, the tensile tests were performed. The results are reported in Table 3. It can be seen that the mechanical properties of T6 and RRA state are close, but the elongation of the RRA sample is slightly lower than that of T6 state. Normally, aluminum alloys tend to suffer a significant decrease in elongation while increasing their strength. However, the samples in TMT states possess a more excellent combination of strength and ductility compared to the conventional heat treatment states. Especially, A1A samples with a deformation rate of 10% and an asymmetric ratio of 1.3 exhibited higher YS and UTS, the UTS was increased by 13% compared to T6, whist the elongation was 13.4%, which is close to the RRA state. When the asymmetric ratio is 1.6, the strength is increased, but the elongation is drastically reduced (A2A and A2N). The comparison between the tensile mechanical properties of the different conditions is intuitively displayed in Fig. 1. It can be concluded that the NTMT here involving asymmetric rolling (r=1.3) and artificial aging has delivered better properties. What’s more, the strength of the samples using the cold rolling deformation is generally lower than that of the asymmetric rolling method. Furthermore, it is worth noting that the strength of NTMT state alloy has slightly decreased when final aging is lowtemperature artificial aging, but the elongation has improved compared with the corresponding natural aging process.

4

TABLE 3 Results of tensile testing performed on samples in all conditions

ASR-TMT

CR-TMT

Process

UTS/MPa

YS/MPa

Elongation/%

A1A

639±0.4

605±1.6

13.4±0.2

A2A

661±2.9

626±3.2

7.1±0.5

A1N

651±0.8

613±0.6

10.2±0.3

A2N

679±0.9

637±5.3

5.6±0.4

C1A

582±1.5

548±1.1

12.6±0.5

C1N

596±3.4

573±2.7

11.3±0.1

T6 RRA

563±2.5 567±3.9

525±3.1 537±3.2

14.3±0.3 13.9±0.3

Fig. 1 Tensile properties of Al-Zn-Mg-Cu alloy with various heat treatments

3.2 SCC performance Slow strain rate tests (SSRT) have been carried out to assess the SCC susceptibility of Al-Zn-Mg-Cu alloy under different states. The sensitivity of stress corrosion cracking can be characterized by ISSRT, which is evaluated by [29]: ISSRT=1−[Rm(3.5%NaCl)×(1+A(3.5%NaCl))]/[Rm(Air)×(1+A(Air))]

(1)

Where Rm (3.5% NaCl) is the tensile strength of samples in 3.5% NaCl solution; Rm (Air) is the correspondingly value determined in dry air. A (3.5% NaCl) is the elongation at fracture of samples in 3.5% NaCl solution; A (Air) is the correspondingly value determined in dry air. We use this index to 5

evaluate the SCC resistance: ISSRT from 0 to 1 means increased SCC sensitivity. The SSRT experimental results are listed in Table 4. As the Table 4 shows that the strength loss of the T6 alloy is the most severe, reaching 11.23%, and T6 sample has the lowest elongation in the 3.5% NaCl solution, the ISSRT value of T6 sample reaches 0.58, indicating that the T6 sample has the worst SCC resistance among these samples. While the ISSRT value of the RRA alloy is only 0.23, implying that RRA alloy has excellent SCC resistance. It is noted that compared with T6, the ISSRT value of NTMT state samples shows a downward trend, and the time-to-fracture in the 3.5%NaCl solution increases. For example, the ISSRT value of A1A state alloy is only 0.20, which maintaining a level similar to that of RRA temper, signifying that A1A has superior SCC resistance. Fig. 2 depicts the slow strain rate tensile curves of various state samples. In general, the larger elongation at failure in the SSRT indicates that the SCC resistance of the alloy is better. The RRA sample reveals the largest elongation, followed by the A1A alloy. Similar to the results in Table 4, it can be seen intuitively that the T6 state alloy has the worst SCC resistance, since its strength and the displacement in the corrosive environment are drastically reduced, the plastic plateau of the stressdisplacement curve of CR-TMT, ASR-TMT and RRA state alloys in corrosive environment is obviously longer than that of T6 alloy, indicating better stress corrosion resistance.

TABLE 4 Slow strain tensile test results of Al-Zn-Mg-Cu alloy

Alloy state

Medium

UTS/MPa

Fracture time/h

Elongation/ %

ISSRT

T6

Air Solution Air

528.4 469.1 533.2

26.3 11.1 25.2

9.5 4.0 9.1

0.58

Solution

518.8

19.5

7.0

A1A

Air

613.7

23.8

8.6

0.20

A1N

Solution Air

595.0 636.1

19.1 22.7

6.9 8.0

0.31

Solution

598.9

15.4

5.5

Air Solution Air

549.7 530.5 563.6

25.9 16.4 24.5

9.3 5.9 8.8

Solution

549.9

12.9

4.7

RRA

C1A C1N

6

0.23

0.35 0.43

Fig. 2 Tensile curves of Al-Zn-Mg-Cu alloy at slow strain rate with different processes: In 3.5% NaCl solution.

3.3. TEM microstructures The relationship between the microstructure and properties of materials is a key concern in materials science and engineering. The different microstructures revealed by TEM along the <001>α directions and select area electron diffraction (SAED) patterns are shown in Fig. 3 (circles are the selected areas). It can be observed from the Fig. 3 (a) and (b) that the size and distribution of the matrix precipitates (MPts) in the T6 and RRA states is similar, fine MPts are uniformly distributed. Combined with the corresponding diffraction spots, it can be judged that these precipitates are mainly η' phase appeared at 1/3 and 2/3{220} Al sites, weak diffraction spots of GP zone can also be found near 2/3{220} Al sites. Therefore, the strengthening phase of T6 and RRA are mainly η' phase and minor GP zone. The average diameter of η' phase is about 8nm. Furthermore, the T6 state sample shows the continuous presence of precipitates along the grain boundary while the RRA state sample shows some discontinuity in the arrangement of precipitates along the grain boundary. It is well known that continuous GBPs (grain boundary precipitate) tend to become an anodic corrosion channel in the corrosive medium, accelerating stress corrosion cracking [30]. That’s why the T6 state alloy has high-stress corrosion sensitivity. Fig.3 (c)-(h) represents the TEM images of CR-TMT state and ASR-TMT state samples. Observing Fig.3 (c) and (g), it can be found that the dislocation-free regions of C1A and A1A are uniformly distributed with a large number of precipitated phases which are finer than the T6 and RRA sample. And the average diameter of precipitated phases in A1A sample is only about 4nm. According to the diffraction pattern, the diffraction spots of η' phase and GP zone can be seen. The GBPs in C1A sample are discontinuously distributed (Fig.3 (d)). It is obvious the GBPs in A1A sample are larger and more separated from each other shown in Fig.3 (g); thereby this further proves that A1A has excellent stress corrosion resistance. However, in C1N and A1N sample, a large 7

number of dislocations can be observed. And dislocations are severely deposited at the grain boundaries (Fig.3 (f) and (i)). According to the diffraction pattern, only weak spots of GP zone and Al3Zr exist. Interestingly, the existence of some sporadic GBPs can be seen, as marked in the Fig.3 (f) and (i). Deformation introduces strain hardening, so the dislocations present at the grain boundaries make it possible to increase the strength while significantly reducing the ductility. This agrees with our previous tensile mechanical studies. For A1A and C1A, after low temperature artificial aging (100 °C/6h), the dislocation density is reduced, and the concentration of dislocations at the grain boundary is alleviated. It can be concluded that the improvement of elongation is related to the high homogeneity of the distribution of nano-scale precipitates and the slight recovery of dislocations during the low temperature aging process.

Fig. 3 TEM maps and corresponding diffraction pattern of Al-Zn-Mg-Cu alloy after different process: (a) T6 temper; (b) RRA; (c), (d) C1A; (e), (f) C1N; (g), (h)A1A; (i) A1N; (a), (b), (d), (f), (h), (i): TEM micrographs of grain boundary; The inserts are the corresponding diffraction patterns.

8

3.4 Dislocation densities The dislocation density is calculated using the Equation (2): ρ=2√3ε/Db

(2)

Where ρ represents the dislocation density, ε is the micro-strain, D is the crystallite size, and b=0.286 nm is the magnitude of the Burgers vector for α-Al. The micro-strains and micro-crystallite sizes and the corresponding dislocation densities are summarized in Table 5 for different processes. Due to the sufficient recovery, the dislocation density of T6 peak aging and RRA state is approximately negligible. And comparing TMT state samples, the dislocation density of the cold rolling method is lower than the dislocation density introduced by the asymmetric rolling, which is consistent with their tensile strength. But the dislocation density in the natural aging state is significantly higher than that in the artificial aging, which indicates that during low-temperature artificial aging, the dislocations underwent a certain degree of recovery, causing some entangled dislocations to disappear and the dislocation density to decrease. The results agree well with the TEM observations. The A1N sample has the highest dislocation density, reaching 3.8x1014/m2, followed by the A1A sample. TABLE 5 Dislocation density of the Al-Zn-Mg-Cu alloy under different TMT states Treatments

Crystallite size (nm)

Strain (%)

Dislocation density(/m2)

A1A A1N

792 640 820 751

0.171 0.203 0.163 0.184

2.6x1014 3.8x1014

C1A C1N

2.4x1014 2.9x1014

3.5 Texture analysis Fig. 4 shows the orientation distribution functions (ODF) of the RRA, CR-TMT and ASR-TMT samples, the measured texture components and volume fractions are shown in Table 6. As shown in Fig. 4a and Table 6, the texture components in the T6 sample consist mainly of the recrystallization textures of Cube {001}<100>, R-Cube {001}<110> and P {011}<122>, and small amount of Goss {011}<100> and S {123}<624> are also observed. However, in the CR-TMT sample the volume fractions of Cube {001}<100> and R-Cube {001}<110> decrease whereas the volume fractions of P {011}<122> increase, and Brass {011}<211> appeared. As compared with the T6 sample and C1A sample, the rolling textures of Brass {011}<211> increase remarkably while the recrystallization textures of Cube {001}<100> and R-Cube {001}<110> decrease in the ASR-TMT samples (Fig. 4b and Table 2). This is attributed to the strong rearrangement of grain orientations occurred during asymmetric rolling. Furthermore, the low-temperature artificial aging at 100 °C /6h doesn’t obviously change the texture characteristics in the A1A samples, which can be confirmed by the fact that the A1A sample has almost the same texture components and volume fractions as the A1N sample. Based on the measured volume fraction of each texture component, the average Taylor factor 9

(MA) values of the samples can be estimated by the equation of MA =ΣfiMi [31], where fi and Mi are the volume fraction and Taylor factor value of a given texture, respectively. According to the measured texture components (Mi) and corresponding volume fractions (fi), the MA values of the RRA, CR-TMT and ASR-TMT samples are calculated and shown in Table 6. Obviously, the A1A sample has almost the same MA value with the A1N sample, which is 3.16 and 3.18, respectively, and higher than that of the T6 sample and C1A sample.

Fig. 4 Orientation distribution function of the Al-Zn-Mg-Cu alloy in various TMT states: (a) T6; (b) C1A (CR10%, AA); (c) A1A (ASR10%, AA); (d) A1N (ASR10%, NA)

TABLE 6 The ODF decomposition results and Taylor factors of the Al-Zn-Mg-Cu alloy S(3.33) Cube(2.45) R-Cube P(3.97) Goss(2.45) B(3.17) Treatments (2.45) T6 18.5 9.6 5.6 2.1 0 6.6 C1A 6.9 5.2 11.9 5.6 7.6 3.2 A1A 9.3 0 12.1 0 41.5 7.7 A1N 9.7 0 13.8 2.3 49.2 4.8

10

MA 2.95 3.05 3.16 3.18

3.6 EBSD analysis High-angle grain boundaries are more susceptible to corrosion than the low angle grain boundaries, so it is important to research the characteristics distribution of grain boundaries to predict the corrosion resistance in a number of materials [32-34]. In this paper, a grain boundary angle between 2.0° and 15.0° is defined as low angle grain boundary (LAGB), and a grain boundary angle higher than 15.0° is defined as high angle grain boundary (HAGB). Fig. 5 shows the EBSD reconstruction map of the Al-Zn-Mg-Cu alloy in various states. The inserts are the distribution of the grain boundary angles of Al-Zn-Mg-Cu alloy in different states. The characteristics distribution of grains and grain boundaries are summarized in Table 7. Observing the EBSD reconstruction map provided in Fig. 5, it can be seen that the grains in RRA state are equiaxed grains; because of the static recrystallization during the solution treatment can promote the formation of the recrystallized grains without substructures. In sharp contrast, the grains of C3A, A1A and A1N state are elongated into a fibrous shape due to deformation, wherein many shear zones of substructure can be observed. As the results summarized in Table 7, we can conclude that compared with the RRA, the proportion of LAGBs in C1A state increased from 51.42% to 85.70%, where the grain boundary angle of 2~5° accounted for a larger proportion as shown accordingly in Fig.5b. Concomitantly, the average grain size has slightly decreased. UTS and YS should increase according to the Hall-Petch effect. When the ASR-TMT was employed, the proportion of LAGBs was further increased. It can be seen that the proportion of LAGBs exceeds 90%. Especially, the proportion in sample A1N even reaches 95.48%, which is the highest of all selected states; whilst the ratio of misorientation angles ranging from between 2~5° is even above 80% as shown in Fig.6c.

TABLE 7 Characteristic distribution of grains and grain boundaries in different state samples Sample

Low angle boundary fraction/%

Average grain boundary angle

Average grain diameter (µm)

RRA

51.4

21.9

11.9

C1A

85.7

6.9°

10.5

A1A A1N

91.5 95.5

9.1° 5.6°

9.3 7.4

11

Fig. 5 Full Euler angle reconstruction of Al-Zn-Mg-Cu alloy in various states: (a) RRA; (b) C1A; (c) A1A; (d) A1N. The inserts are the distribution of grain boundary misorientation angles

3.7 APT analysis Fig. 6 illustrated the APT elemental maps of Al-Zn-Mg-Cu alloy in different states. It can be intuitively seen that the position of the Al, Zn and Mg atoms in the A1A sample are highly coincident, a small amount of Cu atoms can also be detected (Fig. 6(a)), indicating a strong tendency to form Al-Mg-Zn co-clusters. Interestingly, the size of the atom segregation in the RRA state is larger and the shape is more stable (Fig. 6(b)); the segregation of Cu atoms increases. Combining with the quantitative APT results of atomic segregations provided in Table 8. It is noteworthy that the average atomic number of atom segregations in the RRA state even reaches 2034, with an average equivalent radius of ~2.82 nm. According to previous researches [35-36], we tentatively determined that those large atom segregations in RRA are probably GP zones. Meanwhile, the concentration profiles of Al, Zn, Mg and Cu atoms across the atom segregation in RRA state are analyzed within a selected cylinder (Fig. 6c), where the atom segregation is defined by the iso-concentration surface of 15 at.% (Zn+Mg). It can be clearly seen that the Zn concentration is in the range of 20-25 at.%, the Mg concentration is in the range of 15-20 at.%, and the Cu concentration is in the range of 2-4 at.%. The ratio of Zn to Mg is approximately 1.3. This result is roughly consistent with the composition of the GP zones in previous work by Hono et al [37]. So we conclude that the atom segregations in the RRA have evolved into GP zones. And from the results 12

shown in the Table 8, the number density of atom segregations in the RRA state are considerably lower than those of A1A, whereas the average equivalent radius significantly increases, suggesting that during re-aging at 120 °C/24h, large atom segregations in RRA state absorb the surrounding small atom segregations, resulting in a decrease in the cluster number and an increase in the cluster radius, eventually evolved into GP zones. This explanation is consistent with the TEM observation.

Fig. 6 Three-dimensional solute atom maps in the Al-Zn-Mg-Cu alloy with matrix atoms and non-segregated solute atoms removed: (a) A1A (ASR10%, AA); (b) RRA. (c)The composition profile of Al, Zn, Mg and Cu measured using a selected cuboid of RRA sample.

13

TABLE 8 Solute atom clustering statistics: dispersive properties

Average equivalent radius (nm)

Average number of solute atoms

Number density (/µm3)

Volume fraction (%)

clusters number

A1A

0.94±0.001

75.5

(7.6±0.5)×106

0.49±0.04

872

RRA

2.82±0.003

2034.2

(9.5±0.4)×105

0.89±0.05

454

4. Discussion 4.1 Effect of NTMT on SCC of Al-Zn-Mg-Cu alloy Taking the experimental observations above into consideration, the influencing mechanism of SCC was interpreted mainly from three aspects, i.e. GBPs, dislocation distribution and grain boundary type. It is well known that grain boundaries are rich in solute atoms and vacancies, which are beneficial to the nucleation and growth of aging precipitates. Meanwhile, the NTMT treatment introduces dislocations, which also provide an effective site for heterogeneous nucleation of phases and serve as a fast diffusion channel during the aging process, so the phases can be precipitated at a lower temperature. What’s more, the dislocations usually accumulate at grain boundaries, which serve as diffusion channels for grain boundary precipitation [38-40], Moreover, dislocations create links between precipitates, increasing the efficiency of the transformation of GBPs from small precipitates to large ones in low-temperature artificial aging (pre-aging and final aging). Consequently, when the final aging is natural aging, sporadic coarse precipitates can be seen at the grain boundary (Fig.3 (f) and (i)), and in the low temperature artificial aging state sample, the GBPs are discontinuously arranged and coarser (Fig.3 (d) and (g)). The larger size and spacing of the GBP are helpful in resisting the SCC. It is explained by the Hydrogen embrittlement theory [41], which is thought to be the predominant mechanism of SCC in these alloys: the coarse GBPs can act as the trapping sites for free hydrogen atoms and form bubbles to reduce the concentration of hydrogen atoms at the grain boundary and the rate of crack growth. Therefore, the coarser the grain boundary precipitates and the larger the spacing between them, the better stress corrosion resistance. Secondly, the diffusion coefficient of hydrogen atom in aluminum alloy is usually very low, and its diffusion mainly depends on the movement of dislocations [42]. In the natural aging (A1N), there are only coherent GP zones and a small amount of Al3Zr particles, the dislocations can easily cut through the GP zones and continue to move forward. Hence the dislocation accumulation at the grain boundary is very severe, which causes stress concentration to increase the crack growth rate. And the hydrogen atoms will be carried to grain boundaries together with the dislocations, leading to an increase in the concentration of hydrogen atoms at the grain boundaries. When they are accumulated 14

to a certain critical concentration, the alloy will undergo hydrogen embrittlement along the grain boundaries. Conversely, Neguyen et al. [43] reported that an increase in the size of the precipitate, for example, a change from the GP zone to the semi-coherent η' phase will result in a more uniform sliding mode and further reduce hydrogen transported to the grain boundaries. Thence the deformation of the alloy is relatively uniform in the low-temperature artificial aging (A1A), and the accumulation of the dislocation at the grain boundary is reduced. Therefore the concentration of hydrogen atoms inside the grains becomes relatively uniform; the local concentration of hydrogen atoms is difficult to reach the critical concentration causing hydrogen embrittlement, so that the stress corrosion resistance is significantly improved. Meanwhile, the high proportions of LAGBs can make Al-Zn-Mg-Cu alloy maintain good corrosion resistance. Arafin et al. [44] reported that low-angle boundaries are crack-resistant, while HAGB usually has high energy which is more susceptible to corrosion [32]. Moreover, the precipitation free zones (PFZ) are formed adjacent to the HAGBs [45], which can contribute to the occurrence of intergranular fracture. Previous study shows that LAGBs are beneficial to maintaining high ductility [46]. It can be judged from the above that the LAGBs obtained by the thermomechanical treatment accounts for a larger proportion, which is favorable for SCC and ductility. In addition, the fibrous grain structure formed in deformation can also hinder the stress corrosion crack growth. To sum up, the NTMT here involving artificial low temperature aging has better resistance to stress corrosion.

4.2 Effect of NTMT on mechanical properties of Al-Zn-Mg-Cu alloy Al-Zn-Mg-Cu alloy is a typical age hardening alloy, precipitation hardening plays a dominant role in the hardening response of Al-Zn-Mg-Cu alloys, especially the coherent GP zone and semi-coherent η' phase. For samples in T6 and RRA states, the effect of dislocations on strength is negligible and the size of MPts is similar, so their strength is relatively close. In this study, through the dynamic process of hardening-recovery in SSHR, some of the dislocations may be preserved. The pre-aging (80°C/6h) together with dislocations promotes the precipitation of GP zones or co-clusters, and these uniformly dispersed precipitates interact with dislocations during subsequent deformation, resulting in a more uniform deformation, which is beneficial to improving strength and ductility simultaneously. Secondly, the pre-precipitated fine GP zones or co-clusters can provide nucleation sites for the η' in the final aging process (100°C/6h), reducing the precipitation of coarse η directly on the dislocations, which also increases the strength of the alloy. The precipitate size of C1A sample is smaller than that of RRA state, but larger than that of A1A sample (Fig.3c). The high homogeneity of the distribution of nano-scale precipitates in A1A effectively reduces stress concentrations and provides more effective sites to trap and accumulate the dislocations surrounding to the precipitates, which are benefit to the ductility of the samples. Coupled with the slight recovery, the dislocations in the A1A are decreased to provide more lattice space for sliding dislocations, which also contributes a significant enhancement of ductility. The similar effects were reported by S.K. Panigrahi et al. [47] and Cheng et al [48]. Compared with the 15

RRA sample, the grain size of NTMT samples is slightly reduced (Table 7). Besides the high dislocation density introduced during the NTMT would also account for part of the strengthening effect. One strategy to improve strength can be increasing the Taylor factor. Asymmetrical rolling increased a lot of Brass texture fractions. Compared with other samples, the Taylor factor of ASR-TMT samples is significantly increased. Therefore dislocation strengthening, grain refinement, precipitation hardening and texture configuration are the main reasons why its strength is higher than the T6 state. And the high density of nanoscale η' phases were distributed homogeneously in A1A sample, which also contribute a lot to the increase in strength (Fig.3g). Thus the reason why the strength of the C1A sample is lower than that of the A1A sample is mainly due to the size of the precipitates, the dislocation density, and the texture configuration. Now we will discuss the strengthening mechanism of the A1A sample in more detail. It is widely accepted that the yield strength (σys) of a poly-crystalline metal is related to the critically resolved shear stress (CRSS, expressed as τtot) of the grains and the grain boundary strengthening, as described by [49]: σys=∆σgb+Mτtot

(3)

Where ∆σgb is the strengthening due to grain boundaries and M is an orientation factor (often termed the Taylor factor); ∆τtot is the CRSS of the grains. The strengthening effect of grain boundaries can be evaluated according to the Hall-Petch relationship [50-51]: ∆σgb=kHPd-1/2

(4)

where d is the grain size, kHP is the Hall-Petch coeffizient, the exponents are reported to be in the range from 0 to 1 for different classes of materials, the typically value for kHP in aluminum alloys is about 3 MPa·mm1/2. Apparently, the strength decreases as the grain size increases. For example, the average grain size of the A1A is 9.28 µm, while that of RRA is 11.92µm, as shown in Table 7; the corresponding ∆σgb values are estimated to be 31.14 MPa, and 27.48MPa, respectively. In the present model, various contributions to CRSS of grains will be considered: the intrinsic CRSS, ∆τ0; the contribution due to dislocations, ∆τd; the contribution due to solid solution strengthening, ∆τss; the contribution due to co-clusters ∆τcl. Commonly, the τtot is expressed as follows [52]: τtot=τ0+τss+τd+τcl

(5)

And the contribution due to dislocations is represented as [53] ∆τd=Gbρ1/2

(6)

Where ρ is the dislocation density, and α, G, and b are all constants. Therefore the ∆τd of A1N reaches 52.69MPa; the ∆τd of A1A is 43.57MPa. An estimate of τss can be calculated using the following expression [54]: τss = AC0 2/3 16

(7)

Where A is a constant and C0 is the concentration of the solute in weight percent. However, solid-solution strengthening in the TMT is also low due to the precipitation of η' phase, GP zone and the formation of clusters. Therefore, this study does not make specific calculations. Since the co-clusters in A1A are very small and can be easily sheared by dislocations, the Orowan mechanism is not suitable to explain the strengthening effect caused by co-clusters. Marceau et al. [55] proposed a method for calculating the contribution to the critical resolved shear stress (CRSS) by solute atom clusters (τcl): τcl=µbτ*/Ls

(8)

And τ*=0.9cos3/2(φ/2) · [1-cos5 (φ/2)/6] Ls= (2π/3V)

1/2

(9)

·r

(10)

where φ is the average breaking angle of the atom clusters; µ is the shear modulus, in aluminum alloys this value is taken as 27 GPa; Ls is the average spacing of clusters on the glide plane; V is the volume fraction of the clusters; r is the average radius of the atom clusters; b is the Burgers vector, provided b=0.286 nm. And the breaking angle, φi, of all types of atomic clusters can be expressed as φi=2cos-1(rs/rc) [56], here rs is the radius of atomic clusters on the glide plane; rc is the critical radius, taken as 2.1 nm [55]. Thus the average breaking angle and τcl of the solute atom co-clusters can be estimated in Table 8. TABLE 9 Calculated strength increments from the Zn-Mg co-clusters (∆τcl) Volume fraction

Radius

TMT State (%) A1A(ASR10%, 100 °C/6h)

φ (°)

τcl (MPa)

140

72.2

(nm)

0.49

0.94

In summary, the strength of the A1A samples can be increased remarkably with acceptable ductility by NTMT, resulting from the combination of increased rolling textures, high-density dislocations, nano-sized precipitates and refined grains.

5. Conclusions In this paper, the microstructure and evolution of Al-Zn-Mg-Cu alloy in NTMT sate are studied, and the involved mechanisms are analyzed. The conclusions are as follows: (1) An outstanding combination of strength and ductility can be achieved by the novel thermomechanical treatment, under the condition of TMT (A1A); the alloy obtained a tensile strength exceeding the T6 states while maintaining a high elongation. At the same time, it has excellent resistance to stress corrosion. (2) The increased strength of the NTMT samples results from the synergistic effect of the dislocation strengthening, nano-precipitation, texture configuration and co-cluster strengthening. The 17

elongations of the NTMT samples (A1A) increase with the decreased dislocation density and formation of nano-precipitations. (3) The novel thermomechanical treatment induces grain boundary precipitates to be discontinuous, and can significantly increase the proportion of low-angle grain boundaries, especially in the range of 2~5°, thereby improving corrosion resistance. When the final aging is lowtemperature artificial aging, the dislocation distribution at the grain boundary is changed, which also has a significant effect on the stress corrosion resistance of the alloy. Acknowledgements The authors are grateful for the financial support from the National Natural Science Foundation of China (Grant No. 51011120052 and 50871123) and the double first-class discipline construction program of Hunan Province. References [1] K.H. Rendigs. Aluminium structures used in aerospace-status and prospects. Material Science Forum 1997, 242:11-42. [2] J.C. Williams, E.A. Starke Jr. Progress in structural materials for aerospace systems. Acta Materialia. 51 (2003) 5775-5799. [3] A. Heinz, A. Haszler, C. Keidel, S. Moldenhauer, R. Benedictus, W.S. Miller. Recent development in aluminium alloys for aerospace applications. Materials Science and Engineering: A. 280 (2000) 102-107. [4] D. Najjar, T. Magnin, T.J. Warner. Influence of critical surface defects and localized competition between anodic dissolution and hydrogen effects during stress corrosion cracking of a 7050 aluminum alloy. Materials Science and Engineering: A. 238 (1997) 293-302. [5] 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 Al-Zn-Mg-Cu alloys. Corrosion. 98 (2015) 50-62. [6] H.C. Fang, F.H. Luo, K.H. Chen. Effect of intermetallic phases and recrystallization on the corrosion and fracture behavior of an Al-Zn-Mg-Cu-Zr-Yb-Cr alloy. Materials Science and Engineering: A. 684 (2017) 480-490. [7] B. Li, Q.L. Pan, X. Huang, Z.M. Yin. Microstructures and properties of Al-Zn-Mg-Mn alloy with trace amounts of Sc and Zr. Materials Science and Engineering: A. 621 (2015) 173-181. [8] 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 Al-Zn-Mg-Cu alloys. Journal of Alloys and Compounds. 657 (2016) 717-725. [9] S. Chen, K. Chen, G. Peng, L. Jia, P. Dong. Effect of heat treatment on strength, exfoliation corrosion and electrochemical behaviour of 7085 aluminium alloy. Material and Design. 35 (2012) 93-98. [10] J.R. Davis. Corrosion of Aluminium and Aluminium Alloys. American Society for Metals, Materials Park, OH 1999. [11] J. Thompson, E.S. Tonkins, V.S. Agarwala. A heat treatment for reducing corrosion and stress corrosion cracking susceptibilities in 7XXX aluminium alloys. Material Performance. 26 (1987) 45-52. [12] B.M. Cina. Reducing the susceptibility of alloys particularly aluminium alloys to stress corrosion cracking. US Patents; 1974. [13] S.T. Zhao, C.L. Meng, F.X. Mao, W.P. Hu, G. Gottstein. Influence of severe plastic deformation on dynamic strain aging of ultrafine grained Al–Mg alloys. Acta Materialiaialia. 76 (2014) 54-67. [14] S. Valipour, A.R. Eivani, H.R. Jafarian, S.H. Seyedein, M.R. Aboutalebi. Effect of pre-deformation thermomechanical processing on the development of ultrafine grain structure during equal channel angular extrusion. Materials and Design, 86 (2016) 377-384. [15] M.R. Toroghinejad, F. Ashrafizadeh, R. Jamaati. On the use of accumulative roll bonding process to develop nanostructured aluminum alloy 5083. Materials Science and Engineering A. 561 (2013) 145–151. [16] H.J. Lee, J.K. Han, S. Janakiraman, B. Ahn, M. Kawasaki, T.G. Langdon. Significance of grain refinement on microstructure and mechanical properties of an Al-3%Mg alloy processed by high-pressure torsion. Journal of 18

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Highlights: An excellent combination of strength-plasticity and stress corrosion resistance was achieved in Al-Zn-Mg-Cu alloy via novel thermomechanical treatment. Co-cluster, dislocations, texture configuration and nanoprecipitates synergistically improved mechanical properties. A high proportion of low-angle grain boundaries and the distribution of dislocation obtained by NTMT greatly help improve the alloy's resistance to stress corrosion.

Declaration of Interest Statement We declare that we have no financial and personal relationships with other people or organizations that can inappropriately influence our work; there is no professional or other personal interest of any nature or kind in any product, service and/or company that could be construed as influencing the position presented in, or the review of, the manuscript entitled.