Evolution of precipitate microstructure during stress aging of an Al–Zn–Mg–Cu alloy

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Evolution of precipitate microstructure during stress aging of an Al–Zn–Mg–Cu alloy Wei Guo, Jiyan Guo, Jinduo Wang, Meng Yang, Hui Li, Xiyu Wen, Jingwu Zhang

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Received date: 20 November 2014 Revised date: 5 March 2015 Accepted date: 15 March 2015 Cite this article as: Wei Guo, Jiyan Guo, Jinduo Wang, Meng Yang, Hui Li, Xiyu Wen, Jingwu Zhang, Evolution of precipitate microstructure during stress aging of an Al–Zn–Mg–Cu alloy, Materials Science & Engineering A, http://dx.doi. org/10.1016/j.msea.2015.03.047 This is a PDF file of an unedited manuscript that has been accepted for publication. As a service to our customers we are providing this early version of the manuscript. The manuscript will undergo copyediting, typesetting, and review of the resulting galley proof before it is published in its final citable form. 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.

Evolution of precipitate microstructure during stress aging of an

Al–Zn–Mg–Cu alloy

Wei Guoa, b , Jiyan Guoc, Jinduo Wanga, b, Meng Yanga, b, Hui Lia, b, Xiyu Wend, Jingwu Zhanga, b, *

a

College of Materials Science and Engineering, Yanshan University, Qinhuangdao, 066004, China

b

State Laboratory of Metastable Materials Science and Technology, Yanshan University, Qinhuangdao,

066004, China

c

College of Mechanical Engineering, Yanshan University, Qinhuangdao, 066004, China

d

Center for Aluminum Technology, College of Engineering, University of Kentucky, 1505 Bull Lea Road,

Lexington, KY 40511, USA

*

Corresponding author: Tel.: +86-0335-8074631; E-mail: [email protected]

Abstract: The effect of elastic tensile stress on the microstructure of 7075 aluminum alloy aged at 433 K

for 1 hour has been investigated in this paper. It was found that double peaks occur in hardness after

various aging treatments. The peak microhardness of the stress-aged specimens reaches 178 HV, much

higher than that of the conventional aged sample. The stress aging exhibits a diverse microstructure: the

GPII zones and various sizes of Ș precipitates are just identified in the stress-free aged sample; the main

hardening Șƍ precipitates, with the highest and lowest degree of dispersion, are formed after 25 and 50

MPa stress-aged respectively; a finer aging precipitate size distribution, a larger grain boundary

precipitate size and spacing and a wider precipitate free zone are represented in the stress-aged specimens.

After the 25 MPa stress aging treatment, a preferential growth orientation of the larger-size MgZn 2 phase

is observed in 7075 aluminum alloy. During aging, the external stress accelerates the growing of the

larger-size MgZn2 phase, promotes the formation of the Șƍ precipitate and inhibits the formation of the Ș

phase. Our study provides a new process to improve the comprehensive properties of Al–Zn–Mg–Cu

alloys.

Keyword: Aluminum alloys; Aging hardening; Precipitation; Microstructure; Electron microscopy

1. Introduction

7xxx series aluminum alloys are precipitation hardening Al–Zn–Mg–(Cu) alloys, with a vast number

of applications in the aerospace and automotive industries. It is generally believed that the precipitation

sequence of this series aluminum alloys is as follows [1, 2]:

Supersaturated solid solution ĺ vacancy-rich clusters ĺ GP zones (GPI and GPII) ĺ metastable Șƍ

phase ĺ stable Ș phase (MgZn 2)

Proper control of the species, size and density of grain boundary precipitate (GBP) and matrix precipitate

(MPt) is crucial for the alloys to obtain the optimum comprehensive properties [3-11]. In the early stage

of the decomposition of the supersaturated solid solution, two types of GP zones (i.e. GPI and GPII) may

occur [3]. GPI zones [3, 4] which are coherent with the aluminum matrix could be formed over a wide

aging temperature range, from room temperature to 413–423 K, with internal ordering of Zn and Al or

Mg on the {001}Al planes. GPII zones [1, 3] which are Zn-rich layers on {111}Al planes, 1–2 atomic

layers in thickness and 3–5 nm in width, are formed after quenching from temperatures above 723 K and

aging at temperatures above 343 K. Generally either GPI or GPII zones can form as precursors to the Șƍ

phase [12-14], a metastable hexagonal phase, coherent or semi-coherent with the aluminum matrix (a =

0.496 nm and c = 1.402 nm) [15, 16]. The GP zones and the Ș' phase are believed to be responsible for

age hardening in 7xxx series aluminum alloys [1, 17]. The equilibrium Ș phase [18] has a hexagonal

structure with lattice parameters a = 0.521 nm and c = 0.860 nm. It is incoherent with the aluminum

matrix and exhibits numerous crystallographic orientation relationships with the matrix [19].

Conventional T6 and over-aging heat treatments [6, 8, 20] can make the aluminum alloys possess

optimum strength and very good stress corrosion resistance respectively, but at the expense of lowering

the other performance. The long aging period and high energy cost, moreover, is another distinct

shortcoming of these two heat treatments. Subsequently developed RRA [9, 10, 21] heat treatment makes

the alloys obtain very good comprehensive properties. Unfortunately its application is limited to the

small-section aluminum alloys due to the very short retrogression time. Some researches [22-27] in the

last 30 years suggest that the precipitation could be modified dramatically during aging under elastic

stress (namely, stress aging). To analyze this precipitate evolution during stress aging might be helpful to

improve the compromise between the mechanical strength and the corrosion resistance: In 1975, Hosford

[23] firstly observed that the stress has noticeable influence on the habit of șƍ precipitate platelets in

Al–Cu alloy. Afterwards Zhu [24] further investigated the effect and discovered that this stress-orienting

effect depends on the applied stress, temperature and alloy composition. On the other hand, Skrotzki et al.

[25] explored the stress aging process of Al–Cu–Mg–Ag alloy and found that the external stress results in

changes in the size and morphology of the GP zones, șƍ and ȍ precipitates. Bakavos [26] and Lin [27]

paid attention to the stress aging treatment of Al–Zn–Mg–Cu alloys recently and the results revealed that

significant interactions take place during the aging stages.

Although the influence of the stress on aging process has been confirmed by various experiments,

there was little research report on the reasons to bring about it. Besides, no detailed HRTEM (high

resolution transmission electron microscopy) investigations were performed on the early aging stage of

aluminum alloys. Most stress aging effort has focused on the Al–Cu and Al–Cu–Mg–Ag alloys. Our early

study [28] has clearly shown that the Al–Zn–Mg–Cu (7075) alloy exhibits pronounced responses to the

stress aging. It is very necessary to find out the interactions between the external stress and the precipitate

evolution during aging. In this paper, aging treatment under various stresses was explored in an

Al–Zn–Mg–Cu (7075) alloy to further determine the influence of the applied stress on the early

precipitation process using TEM (transmission electron microscopy) and HRTEM methods. The

mechanism of the stress aging is also discussed.

2. Materials and methods

The compositions of the experimental 7075 aluminum alloy are showed in Table 1. Elastic stress

aging experiments were carried out in a RDL–50 creep testing machine. External stress was applied on

the specimens along the longest axis, which precisely corresponded to the rolling direction. All the sheets

were solution treated at 763 K for 1.5 h, cold water quenched, then artificially aged at 433 K for 1 hour

with and without the stress. The size of the aging treated samples is 2×6×75 (mm).

Table 1

The value of tensile yield strength (ıs) of solid-soluted specimen at 433 K was 220 MPa. The

maximum tensile stress used in the aging treatment is less than 0.6ıs. Vickers microhardness was carried

out on a FM–700 Hardness Tester with a load of 300 g and a loading time of 10 s. The grain size was

observed in a Zeiss Axiovert 200 MAT optical microscope. The SEM investigations were carried out on a

KYKY–2800B scanning electron microscope equipped with energy dispersive spectrometer (EDS)

operating at 25 kV. X-ray diffraction (XRD) experiments were performed on a Japanese Rigaku

diffractometer using Cu KĮ radiation. Microstructures were characterized by transmission electron

microscopy (TEM) together with high resolution transmission electron microscopy (HRTEM). The TEM

samples were prepared by mechanical grinding and electropolishing in a solution of 30% nitric acid and

70% methanol at 243 K.

3. Results and discussions

There are many kinds of phases in aluminum alloy, which differ in size, distribution and formation

conditions. Various phases have diverse strengthening effects on the aluminum alloy. According to the

precipitate size, the particles in aluminum alloy are divided into three types. The first type is the coarse

insoluble particles. Its size is usually above 1 ȝm. This type of particle is brittle and is harmful to the

fracture toughness of aluminum alloys. The second type is the intermediate particle. Its size is about

0.05–0.5 ȝm. This type of particle can inhibit recrystallization and grain growth during the

homogenization process. The third type is the aging precipitate, and its size is below 0.05 ȝm. This type

of particle is believed to be responsible for age hardening in aluminum alloys.

Fig. 1

The optical microscope and scanning electron microscope images of the 7075 aluminum alloy after

solution treated are showed in Fig. 1. It can be found from Figs. 1a and b that a lot of chain-like coarse

insoluble precipitates were evenly distributed along the rolling direction of the aluminum sheet. And the

specimens had the same grain size of 130 ȝm after solution treated. After aging treatment no dimensional

changes within ±0.04 mm could be detected with a micrometer, indicating a plastic strain of less than 0.1

pct. Fig. 1c shows the high-magnification image of the coarse insoluble precipitates. It reveals from Fig.

1c that two kinds of coarse insoluble precipitates with a total volume fraction of 0.85% can be found in

the aluminum matrix. The SEM observation also reveals that particles (bright precipitates) were of

different sizes. The average size of these coarse insoluble precipitates was about 10 ȝm. The XRD

patterns of the stress(-free) aging treated samples are showed in Fig. 1d. The XRD peaks could be

identified due to Į(Al) solid solution (labeled Al), Al23CuFe4 phase (labeled Al23CuFe4), Al7Cu2Fe phase

(labeled Al7Cu2Fe), and Șƍ and Ș phases (labeled MgZn2). Fig. 1d reveals that the primal aging

precipitates formed in the stress-free and stress aging treatments were both MgZn2 phase.

Energy dispersive spectrometer analysis is carried out to identify the chemical constitutions of the

coarsen insoluble precipitates. The EDS analysis results of precipitate P1 (72.44 at% Al, 14.86 at% Cu,

7.58 at% Fe) and precipitate P2 (62.12 at% Al, 4.16 at% Cu, 25.88 at% Fe) showed in Fig. 1c are

exhibited in Fig. 2, which are close to the stoichiometric constitutions of Al7Cu2Fe phase and Al23CuFe4

phase respectively. These second-phase particles appear during the casting or solidification processes and

can be detected by optical microscopy. In 7075 aluminum alloy the Al23CuFe4 and Al7Cu2Fe

intermetallics are the main second-phase particles and can undergo phase transformation and change their

morphology during ingot homogenization [29]. This result is consistent with the XRD observations (Fig.

1d). The coarse insoluble particles are of large dimension, low plastic deformability, high elastic module

and large non-coherent contact with the aluminum matrix. Proper control of the distribution and the

amount of the coarse insoluble particles is helpful to strengthen the aluminum alloy. Otherwise, it

embrittles the material and results in poor strength. Because the stress aging temperature is much lower

than the dissolution temperature of the coarse insoluble particles, the particles cannot redissolve into the

aluminum matrix during the stress aging. Therefore, the stress aging treatment has little influence on the

size and morphology of the coarse precipitates. Our main objective is to find out the effects of the applied

stress on the smaller-size precipitates (0–0.5 ȝm) during aging.

Fig. 2

3.1. Effects of stress aging process on the intermediate precipitate

Typical TEM microstructures of the experimental 7075 aluminum alloy after solution treated with or

without stress aging are showed in Fig. 3. It reveals that two kinds of intermediate precipitates were

uniformly distributed in the aluminum matrix. One was the dispersed compounds, such as Al3Zr, Al7Cr

and Al18Mg3Cr2. The dispersoids are effective at pinning the grain boundaries and can maintain a fine

grain size. They are formed in the homogenization process. The other was the larger-size MgZn2 phase.

Most larger-size MgZn2 phases are coming from the coarsing MgZn2 phases reserved in the samples after

solution treated. These intermediate precipitates are incoherent with the aluminum matrix. And the size of

the precipitates was about 0.04–0.20 ȝm.

Fig. 3

In order to describe the effects of the stress aging process on the evolution of the intermediate

precipitates, we examined the size, shape and distribution of the intermediate particles. As shown in Fig. 3,

there were no obvious changes in the mean size, shape and distribution of the dispersed compounds after

the stress aging process. It seems that the applied stress has little influence on the development of the

dispersed compounds. This is attributed to the coarsening mechanism of the dispersed compounds. The

dispersoidsƍ coarsening is through a discontinuous mechanism as dislocations or grain boundaries cross

them. This mechanism is only activated at high temperature, much higher than the aging temperature [30].

So, the stress aging process has little influence on the size, shape and distribution of the dispersed

compounds in the 7075 aluminum alloy under the experimental conditions. Similar experimental findings

were obtained by [27].

Compared to the stress-free aged treatment, however, the stress aging process makes some

discrepancies in the size, shape and distribution of the larger-size MgZn2 phase. Finely distributed

precipitates were formed in the grains of both aging treated specimens, while a high degree of precipitate

dispersion was only produced in the stress-aged specimens. The mean size of the rod-like Ș phase was

larger in the stress-aged samples than in the stress-free aged sample. And obviously, the distributions of

precipitate size were broader in the stress-aged conditions. It is concluded that the growing of the

larger-size MgZn2 phase is promoted by the applied tensile stress. Besides, it is important to note that

large rod-like or bar-like MgZn 2 phases, which only developed from solution-treated condition under an

applied stress of 25 MPa (Figs. 3c and d), were distributed either on grain boundaries or intragranularly.

As is reported, the stress aging process usually leads to particular microstructural changes such as a

stress-orienting effect of the plate-like coherent or semi-coherent precipitates in 2000 series aluminum

alloy [23, 24]. And no stress-oriented effect can be found in the studied 7075 aluminum alloy due to the

symmetrical structure of the precipitates in 7000 series aluminum alloy [27, 31]. In our study, however, it

is discovered that the 25 MPa stress aging treatment created a preferential growth orientation along the

MgZn 2 (002) plane of the larger-size MgZn2 phase in 7075 aluminum alloy. The interaction energy for

individual Ș plate is a function of the orientation of the plate. The driving forces for the growth of the Ș

plates are different for Ș plates oriented differently with respect of the direction of the applied stress [24].

This stress-orienting effect is certainly related to the divergences in equilibrium interfacial solute

concentration, which is dependent of the associated elastic strain energy. Only a small number of large

rod-like or bar-like MgZn2 phases were presented in the 25 MPa stress-aged sample, indicating that there

are other factors (except the applied tensile stress) that are associated with the preferentially oriented

growth. A detailed study on this stress-orienting effect would be of interest in the future.

Fig. 4

The effects of the applied stress on the coarsening of the larger-size MgZn 2 phase are significant.

The ultra-high strength properties of Al–Zn–Mg–Cu alloy are greatly determined by interactions between

the dislocations and precipitates, and these interactions are closely relevant with the sizes of precipitates

[32]. As shown in Fig. 4, it can be observed that a number of dislocations were piled up around the

larger-size MgZn 2 phases in the stress-aged specimen. When the loading strains are elastic in the

stress-aging process, creep occurs by diffusion-controlled movement of dislocations in the alloy matrix

[22, 33]. The intermediate particles, under the experimental conditions, retards the movement of the

dislocation. The motion of the dislocation in return accelerates the coarsing of the larger-size MgZn2

phases. Because of the large size (0.04–0.20 ȝm) of the intermediate precipitates, the deformation

mechanism of the 7075 alloy obeys the Orowan way. Due to the effect of the applied tensile stress, the

multiplication of dislocation strengthens the alloy.

3.2. Effects of stress aging process on the aging precipitate

Fig. 5

Fig. 5 shows the evolution of Vickers microhardness in the course of stress aging at 433 K. As can be

seen, double peaks occured after aging treatment under various stresses. The hardness of 7075 aluminum

alloy aged at 433 K for 1 hour was more influenced by a lower applied stress (25 MPa). The peak

microhardness value of the stress-aged specimens reached 178 HV, much higher than that (165 HV) of the

stress-free aged sample. There was no evident change (even a little lower) in hardness after a moderate

stress (50 MPa) aging compared to the conventional aging. The mechanical behavior is controlled by

microstructures. This hardness variance indicates that the effect of the applied stress on the aging

precipitate is somewhat complicated across the whole range of stress aging.

Fig. 6

A comparison of series of TEM and HRTEM micrographs for the aging treated specimens displayed

that noticeable differences emerged in the species, size and density of the aging precipitates.

Representative precipitate microstructures in 7075 alloy after various aging are showed in Fig. 6. The

TEM images exhibit three typical types of precipitate: spherical particles, plate-shaped and equiaxed

precipitates. To identify the second-phase precipitates, we examined the HRTEM images in [011]Al and

[112]Al projection of each sample. As shown in Fig. 7a, after aging treatment, plate-like GPII zones which

were thought to be linked to the vacancy-rich clusters or vacancy-related clusters were identified in the

stress-free aged sample. The GPII zones induce considerable elastic strain into the surrounding aluminum

lattice which makes them appear thicker at lower magnifications. This makes them could not be

distinguished clearly from the Ș' platelets by TEM examination, but could be imaged quite readily by

HRTEM in [110]Al projection. The dark or bright contrast on one side of the layers in Fig. 7a indicates an

associated strain-field. No GPI zones were identified in our observations. Equiaxed, stable hexagonal Ș

phases were found in the conventional aged sample (Figs. 7b and d). From Fig. 7f it can be seen that the

equilibrium Ș phase was apparently incoherent with the aluminum matrix. And the measured c lattice

parameter of the Ș phase was 0.86 nm. Two kinds of image contrasts related to the Ș' precipitates were

noticed in the images of all specimens. One feature was an image of the precipitate with considerable

disorder (Fig. 8a), corresponding to the spherical particles in Fig. 6. The other feature was a set of parallel

line contrasts (Fig. 8d). The measured c parameter of the latter Ș' precipitate along the <111>Al direction

was approximately equal to 6d111Al (=1.403 nm).

Fig. 7

In order to further describe the effects of stress-aging process on the size and distribution of the

aging precipitates, the mean sizes of aging precipitates were measured, as shown in Fig. 9. Microstructure

of the stress-free aged sample was represented coarser precipitate size distribution than that of the

stress-aged specimens, which is in agreement with the observations of Bakavos [26]. Very densely

distributed fine Șƍ precipitates, which contribute more to the hardness [34], were formed in the alloy after

25 (Fig. 6b) and 100 (Fig. 6d) MPa stress-aged respectively. The plate-like GPII zones [3] (Fig. 7a) and

the Ș' precipitates [15] were the main hardening precipitates in the conventional aged sample, while the Ș'

precipitates (Figs. 8a and d) were the only hardening precipitates in the stress-aged specimens. This

makes the 25 and 100 MPa stress-aged specimens achieve the highest hardness. A lower degree of

precipitate dispersion, but a larger mean Ș' precipitate size created the same hardness in the 50 MPa

stress-aged specimen, compared to the stress-free aged sample. It seems that the 25 and 100 MPa stresses

decreased the size of the aging precipitates, while the 50 MPa stress increased the size of the precipitates

during aging. We contribute this anomalous phenomenon of the latter to the less nucleation of the

precipitates (i. e. more supersaturation) during the 50 MPa stress-aging.

Fig. 8

All these microstructural observations have been confirmed by measuring the microhardness (Fig. 5).

Besides, another striking result of our work is the appearance of equiaxed Ș phases [18] (Figs. 7b and d)

with various sizes in the stress-free aged sample, but only very few big Ș phases were observed in the

stress-aged specimens instead. From the TEM and HRTEM results it is concluded that the formation of

the Șƍ precipitate is promoted and the formation of the Ș phase is inhibited by the applied tensile stress

during the early stage of stress aging process. This conclusion means a lot in the strengthing of 7075

aluminum alloy.

Fig. 9

Fig. 10 presents the TEM micrographs of grain boundary precipitates of the stress-free and

stress-aged 7075 aluminum alloy. From Fig. 10, it can be found that two kinds of situations about

precipitates were presented in the grain boundaries of both aging treated specimens: precipitates formed

in the grain boundaries with smaller size (5−10 nm), as shown in Figs. 10a and c; precipitates formed in

the grain boundaries with larger size (30−60 nm), as shown in Figs. 10b and d. In the first condition, the

precipitates were continuously distributed on the grain boundaries in the stress-free aged sample, with a

relatively narrow (0−55 nm) precipitate free zone. While no continuous layer of grain boundary

precipitates was found in the stress-aged specimen. And the width of the precipitate free zone increased to

70 nm after the stress aging treatment. In the second condition, the grain boundary precipitates under the

stress aging process coarsened and concentrated more quickly. And the width of the precipitate free zone

in the stress-aged sample (90 nm) was much larger than that in the stress-free aged sample (70 nm). From

Fig. 10, it is concluded that the stress aging process promotes the growth of the grain boundary

precipitates and increases the width of precipitate free zone.

Fig. 10

The grain boundary precipitates are generally considered as anodic to the aluminum matrix. Grain

boundaries with continuous precipitates are more susceptive to the intergranular corrosion. While

discontinuously distributed grain boundary precipitates can decrease the corrosion susceptibility [35]. The

existence of the wide precipitate free zone is beneficial to the relaxation of stress, and can decrease the

probability of fatigue cracking [36]. On the basis of the above results, it is supposed that the stress aging

treatment could improve the stress corrosion resistance of 7075 aluminum alloy.

Precipitation in multicomponent alloys can take complex pathways depending on the relative

diffusivity of solute atoms and on the relative driving forces involved. As reported, Mg plays a key role in

the formation of GP zones [1, 37] and Zn/Mg ratio has been suggested as the limiting factor in the

formation of the Șƍ and Ș phases [1]. Being in the tensile stress field, the vacancy-solute atom complexes

[22] formed after quenching will segregate at the dislocations or grain boundaries, that leads to the

variance in the migration rates of the solute atoms. G. Sha suggested that Mg segregates strongly to the

grain boundaries in Al–Zn–Mg–Cu alloy processed by equal-channel angular pressing (ECAP) at 473 K

[38]. It is believed that Mg is more influenced by the external stress during aging because of the stronger

Mg-vacancy interaction [14]. So the accompanied changes of the Mg content and Zn/Mg ratio result in

the divergence in the species, size and distribution of the aging precipitates. The finely formed Șƍ

precipitates and the missing GPII zones and Ș phases as well as the wide precipitate free zone after stress

aging are good proofs to this modification. In macro presentation, this discrepancy reflects on the

microhardness. The increasing Mg segregation into the precipitates after stress aging is believed helpful

to inhibit the formation of the Ș phase and is supposed to decrease the stress corrosion susceptibility [39].

It can be seen that the hardness and the stress corrosion resistance are improved using stress aging

treatment. This provides additional insight to the formation and growing of various precipitates in the

early aging stage of Al–Zn–Mg–Cu alloys and is helpful to understand the varied microstructure through

thickness in creep-age forming process. Investigations of precipitation and diffusion processes of

Al–Zn–Mg–Cu aluminum alloys during stress aging are of significant importance from viewpoints of

both basic science and engineering application.

4. Conclusions

In summary, processing by various tensile stress aging treatments of an Al–Zn–Mg–Cu (7075)

aluminum alloy at 433 K for 1 hour has been found to result in remarkable and distinctive changes in

behaviour:

(1) The main coarse insoluble precipitates (Al23CuFe4 and Al7Cu2Fe intermetallics) with a total

volume fraction of 0.85% are distributed along the rolling direction in 7075 aluminum alloy. The elastic

stress aging has little influence on the insoluble particles.

(2) Two kinds of intermediate precipitates (dispersed compounds and larger-size MgZn2 phase) are

uniformly distributed in the aluminum matrix. The applied stress has little effect on the development of

the dispersed compounds, but promotes the growing of the larger-size MgZn 2 phase. A 25 MPa stress

aging treatment creates a preferential growth orientation of the larger-size MgZn2 phase in 7075

aluminum alloy.

(3) Double peaks occur in hardness after stress aging. The peak microhardness of the stress-aged

specimens reaches 178 HV. The GPII zones and the Ș phases with various sizes are just identified in the

stress-free aged sample. Very densely distributed fine Șƍ precipitates are formed in the 25 MPa stress-aged

specimen, while the degree of precipitate dispersion was much lower in the 50 MPa stress-aged condition.

The external stress represents a finer precipitate size distribution and a larger grain boundary precipitate

spacing. The stress aging process accelerates the formation of the Șƍ precipitate, suppresses the formation

of the Ș phase, promotes the growth of the grain boundary precipitates and increases the width of the

precipitate free zone.

Acknowledgments

The authors would like to express their acknowledgment for the financial support from the National

Natural Science Foundation of China (No. 51071140).

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Figure and table captions

Table 1

Compositions of the 7075 aluminum alloy.

Fig. 1. Optical micrographs: (a) long transverse rolling direction; (b) longitudinal rolling direction,

scanning electron micrograph (c) of 7075 aluminum alloy after solution treated. (d) X-ray diffraction

patterns of 7075 aluminum alloy stress-aged at 433 K for 1 hour.

Fig. 2. EDS analysis of the coarse insoluble precipitates of 7075 aluminum alloy after solution treated: (a)

P1 particle and (b) P2 particle.

Fig. 3. (a) Microstructure of 7075 aluminum alloy after solution-treated. Bright field images of 7075

aluminum alloy aged at 433 K for 1 hour under various tensile stresses: (b) 0 MPa; (c), (d) 25 MPa; (e) 50

MPa; (f) 100 MPa.

Fig. 4. Interactions between the dislocations and intermediate precipitates of 7075 aluminum alloy aged at

433 K for 1 hour under a 100MPa tensile stress.

Fig. 5. Age hardening responses of 7075 aluminum alloy stress-aged at 433 K for 1 hour. The error bar is

the standard deviation.

Fig. 6. Bright field images in [112]Al projection from 7075 alloy aged at 433 K for 1 hour under various

tensile stresses: (a) 0 MPa; (b) 25 MPa; (c) 50 MPa; (d) 100 MPa.

Fig. 7. [011]Al (a) and [112]Al (b and d) HRTEM micrographs of 7075 alloy stress-free aged at 433 K for 1

hour: a shows the GPII zones and b and d shows the Ș phases; c and e are the fast Fourier transforms

(FFT) of b and d, respectively; f is the inverse FFT (IFFT) of d.

Fig. 8. [011]Al (a) and [112]Al (d) HRTEM images, and the corresponding FFTs (b and e) and IFFTs (c and

f), showing two kinds of Șƍ precipitates in the 25 MPa stress-aged microstructure of 7075 alloy.

Fig. 9. The mean sizes of aging precipitates in 7075 aluminum alloy after various aging treatments. The

error bar is the standard deviation.

Fig. 10. The TEM micrographs of grain boundary precipitate in the 0 MPa (a, b) and 100 MPa (c, d)

stress-aged samples of 7075 alloy.

Table 1

Compositions of the 7075 aluminum alloy.


Z

Zn

Mg

Cu

Cr

Fe

Si

Mn

Ti

Al

Wt.%

5.85

2.57

1.50

0.21

0.16

0.06

0.05

0.02

Bal




Figure 1

Figure 2

Figure 3

Figure 4

Figure 5

Figure 6

Figure 7

Figure 8

Figure 9

Figure 10