Effect of Sc and Zr additions on grain stability and superplasticity of the simple thermal–mechanical processed Al–Zn–Mg alloy sheet

Author’s Accepted Manuscript Effect of Sc and Zr additions on grain stability and superplasticity of the simple thermal-mechanical processed Al-Zn-Mg alloy sheet Y.L. Duan, G.F Xu, X.Y. Peng, Y. Deng, Z. Li, Z.M. Yin www.elsevier.com/locate/msea

PII: DOI: Reference:

S0921-5093(15)30389-0 http://dx.doi.org/10.1016/j.msea.2015.09.049 MSA32775

To appear in: Materials Science & Engineering A Received date: 11 August 2015 Revised date: 5 September 2015 Accepted date: 11 September 2015 Cite this article as: Y.L. Duan, G.F Xu, X.Y. Peng, Y. Deng, Z. Li and Z.M. Yin, Effect of Sc and Zr additions on grain stability and superplasticity of the simple thermal-mechanical processed Al-Zn-Mg alloy sheet, Materials Science & Engineering A, http://dx.doi.org/10.1016/j.msea.2015.09.049 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.

Effect of Sc and Zr additions on grain stability and superplasticity of the simple thermal-mechanical processed Al-Zn-Mg alloy sheet Y.L. Duana, G.F. Xua, b, , X.Y. Penga, Y. Denga, b, Z. Lia, b, Z.M. Yina, b a

b

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

Key Laboratory of Nonferrous Materials Science and Engineering of Ministry of Education, Central South

University, Changsha 410083, China Abstract: Effect of scandium and zirconium on grain stability and superplastic ductility in the simple

thermal-mechanical processed Al-Zn-Mg alloys was investigated. Tensile testing revealed that the Al-Zn-Mg alloy without Sc and Zr additions showed no superplasticity because of the larger grain size (>10 μm) and the poor stability of the microcrystalline structure during superplastic deformation. However, the Al-Zn-Mg-0.25Sc-0.10Zr (wt.%) alloy exhibited excellent superplastic (elongations of ≥ 500 %) at a wide temperature range of 450 ~ 550 ℃ and high strain rate range of 5×10−3 ~ 5×10−2 s−1, and the maximum elongation of ~1523 % was achieved at 500 ℃ and 1×10−2 s−1. Electron back scatter diffraction analysis and transmission electron microscopy results showed that superior superplastic ductility of the Al-Zn-Mg-0.25Sc-0.10Zr alloy can be ascribed to the complete transformation of low angle grain boundaries to high angle grain boundaries due to the occurrence of dynamic recrystallization and the presence of coherent Al3ScxZr1-x particles that effectively impede the growth of the grains during superplastic deformation. Besides, strong β-fiber rolling textures gradually weakened, and random textures were predominant in the superplastic deformed alloy. Analyses on the superplastic data revealed that the average strain rate sensitivity and the average activation energy of the Al-Zn-Mg-0.25Sc-0.10Zr alloy were ~0.37 and ~84.5 kJ/mol–1, respectively. All results indicated that the main superplastic deformation mechanism was grain boundary sliding.

Keywords: Superplasticity; Al-Zn-Mg alloy; low/high angle grain boundary; Al3ScxZr1−x particles; Deformation mechanism

1. Introduction Superplasticity refers to the ability of material to pull out to a high tensile elongation without 

Corresponding author, Tel: +86-731-88877217. E-mail address: [email protected]. 1

the development of necking within the gauge length. In general, the fine grain size (typically less than 10μm) and the thermal stability of the fine microstructure at high temperatures are two major prerequisites for achieving structural superplasticity [1]. Superplastic forming (SF) is widely used to fabricate complex parts with metallic sheets. However, the expansion of SF into the fabrication of high-volume components in commercial applications is currently limited, because the production of superplastic aluminum alloys is relatively expensive due to the complex thermomechanical processing is typically necessary to make these materials superplastic. Therefore, to advance SF into production oriented industries, some complicated processing technologies have been used in 2xxx, 5xxx and 7xxx series aluminum alloys to reduce the grain size , such as equal channel angular pressing (ECAP) [2-5], friction stir processing (FSP) [6-9], high-pressure torsion (HPT) [10], multiaxial alternative forging (MAF) [11], and accumulative roll bonding (ARB) [12]. Although, the above processes may be readily utilized to achieve ultra-fine grains and superplasticity, extra cost and additional technologies are required, and they are not suitable for engineering production. What is more, FSP and ECAP can only refine the grains significantly in a small area of materials but not the entire area. Those barriers restrict the expansion application of those technologies on the commercial SF. Furthermore, to stabilize the fine grains, suitable alloying elements (eg. Zr or Sc [13-15]) were added to overcome the poor stability of the microcrystalline structure of aluminum alloys at high temperatures. In the recent study involving a Zr-containing Al-Zn-Mg-Cu alloy AA 7449, it was demonstrated that the alloy exhibited high values of total elongation using appropriate combinations of strain rates and temperatures [16]. As well known, besides, in Al alloys containing both Zr and Sc,

2

the Al3ScxZr1-x dispersoids have proven more effective in pinning grain and subgrain boundaries to inhibit recrystallization grain growth than Al3Sc or Al3Zr [17]. They have the potential to be used in the commercial SF. However, it should be pointed out that all the superplastic investigations of the Al-Zn-Mg alloy containing Sc or/and Zr were conducted using some complicated processing technologies, i.e., FSP, ECAP…… For superplastic forming production, the aim is to improve the structural performance of the workpiece and reduce the fabrication cost. If high strength commercial Al-Zn-Mg-Sc-Zr alloy processed by a simple thermal-mechanical processing that only includes homogenization, rolling and solution treatment is superplastic deformed directly, it would significantly shorten the fabrication period of superplastic forming and reduce the fabrication cost. However, few researches on the superplastic of the Al-Zn-Mg-Sc-Zr alloy processed by a simple thermal-mechanical processing have

been reported so far. In this work, the superplastic behavior of two kinds of commercial Al-Zn-Mg alloys with and without Sc and Zr additions subjected to a simple thermal-mechanical processing is to be investigated. The purposes of this study are: (i) to evaluate the superplastic behavior of two kinds of Al-Zn-Mg alloys with and without Sc and Zr additions at various temperatures and strain rates; (ii) through calculating the coefficient of strain rate sensitivity and the activation energy to clarify the superplasticity deformation mechanism; (iii) last, but most important, to further analyse the influence of scandium and zirconium on grain stability and superplastic ductility in Al-Zn-Mg alloys.

2. Experimental

3

Two kinds of Al-Zn-Mg alloys with and without Sc and Zr additions were applied for comparative research. Semi-continuous ingots, with 172 mm in diameter, were provided by Northeast Light Alloy Co. Ltd. Table 1 shows their chemical compositions. The process history of the alloys is as follows: semi-continuous ingot was homogenized at 350 ℃ for 8 h firstly and then at 470 ℃ for 12 h. The homogenized ingot was hot rolled to 7 mm thick plates and then cold rolled to 2 mm sheets. Finally, it was subjected to solution treatment at 470 ℃ for 1 h, followed by water quenching, and then aged at 120 ℃ for 12 h. Table 1 Chemical composition of studied alloys (in wt.%).

Alloy

Zn

Mg

Sc

Zr

Cu

Mn

Si

Fe

Al

Al-Zn-Mg

5.39

1.91

-

-

0.34

0.33

0.08

0.16

Bal.

Al-Zn-Mg-0.25Sc-0.10Zr

5.41

1.90

0.25

0.10

0.33

0.32

0.11

0.18

Bal.

Tensile specimens with a gauge length of 6 mm and a gauge width of 4 mm were machined from the final aged sheet; and the gauge length was parallel to the rolling direction. Tensile tests were carried out in the temperature interval 450 ~ 550 ℃ and at strain rates ranging from 1×10−3 s−1 to 1×10−1 s−1. Each sample was held at a testing temperature for about 20 min prior to applying the load in order to reach a thermal equilibrium. The temperature stability during tests was better than 1 ℃ and the maximum gradient along the specimen did not exceed 3 ℃. Tensile tests with a constant crosshead speed were conducted by using an Instron 8032 electro-fluid servo-fatigue tester. The failed specimens were subjected to scanning electron microscopy (SEM). The microstructures of the undeformed specimen and the specimens which were pulled to different strains were examined by transmission electron microscopy (TEM), electron back scattered

4

diffraction(EBSD) technique. SEM observations were carried out on a FEI Sirion 200 field emission gun scanning electron microscope, operating at 20 kV. Thin foils for transmission electron microscope (TEM) observations were sectioned the center part of the tensile samples on the longitudinal - transverse (L-T) section. The foils were prepared by twin-jet electropolishing at 20 V in a solution of 30% nitric acid and 70% methanol solution cooled to −30 ℃ and observed on a TECNAIG2 20 electron microscope and a JEM-3010 electron microscope respectively, both with an acceleration voltage of 200 kV. EBSD analyses were performed using a FEI Sirion 200 field emission gun scanning electron microscope equipped with EBSD system. Usually scans were taken at magnification 200× and beam parameters 25 kV and 18 nA. Typical scan area was 250 μm × 250 μm with step size equal to 0.50 μm. EBSD data was subsequently analyzed by the OIM Analysis 5.0 software. Moreover, misorientations below 2° were not measured in order to avoid spurious boundaries, i.e. the orientation uncertainty or “orientation noise”. This limit was used for all samples in order to provide consistent quantitative data. Thus, a low-angle grain boundary (LAB) was defined by a misorientation between 2° and 15°, and a high-angle grain boundary (HAB) was defined by misorientation >15°. HABs and LABs were shown as black and white lines on the EBSD maps, respectively.

3. Experimental results 3.1 Elongation to failure Fig. 1 shows the variations of elongation with an initial strain rate under different deformation temperatures for two studied alloys. It is mentioned that, with the increase of deformation

5

temperatures or strain rates, the elongations increase firstly and then decrease. However, the Al-Zn-Mg alloy never shows superplastic behavior and all elongations to failure are less than 200% under all testing conditions. 220

200

Elongation (%)

180

(b) 1600

450℃ 475℃ 500℃ 525℃ 550℃

450℃ 475℃ 500℃ 525℃ 550℃

1400 1200

Elongation (%)

(a)

160

140

1000 800 600 400

120

200

100 -4 10

10

-3

-2

10

10

-1

10

0

10

-4

10

-3

10

-2

10

-1

10

0

-1

-1

Strain Rate (s )

Strain Rate (s )

Fig. 1 Variation of elongation with initial strain rate at various testing temperatures for two studied alloys:

(a) Al-Zn-Mg, (b) Al-Zn-Mg-0.25Sc-0.10Zr

Fig. 2 Macrograph of Al-Zn-Mg-0.25Sc-0.10Zr alloy samples deformed to failure at different testing conditions.

In contrast, the Al-Zn-Mg alloy with Sc and Zr additions exhibits superplasticity (elongation ≥500%) in a wide temperature range of 450 ~ 550 ℃ with higher strain rates of 5×10−3 ~ 5×10−2 s-1. For the Al-Zn-Mg-Sc-Zr alloy, besides, maximum elongation of ~1523% is obtained at a temperature

6

of 500 ℃ with an initial strain rate of 1×10−2 s-1. Fig. 2 shows macrographs of Al-Zn-Mg-Sc-Zr alloy samples deformed to failure at different testing conditions. The Al-Zn-Mg-Sc-Zr alloy exhibits superior superplasticity. Moreover, it can be noted that relatively uniform deformation occurs within the gauge zone, and there is no obvious necking. 3.2 Superplastic behaviors of Al-Zn-Mg-Sc-Zr alloy 3.2.1 Stress-strain relationship Fig. 3 shows the true stress - true strain (σ-ε) curves for the Al-Zn-Mg-0.25Sc-0.10Zr alloy at typical deformation conditions (at 1 × 10-2 s-1 or 500 ℃). It can be found that the values of the true stress rise rapidly and then hold constant or decrease to some extent after attaining the peak values with the increase of the true strains at different deformation conditions, showing a dynamic flow softening. The reason is that dislocation density increases dramatically leading to the rapid increase of stress at the beginning of the deformation, However, as deformation proceeds, dynamic softening arising from a recrystallization process may occur, which can offset or partially offset the effect of work hardening, and the true stress may be kept unchanged or fall to some extend with increasing deformation.. Moreover, the peak values of true stress decrease with the increase of deformation temperatures or the decrease of the strain rates, and can be observed at the strain of ~0.7. When the strain value is more than 1 at temperatures ranging from 425 to 550 ℃, the stress oscillation appears, which means that there is strain hardening that was reported in FSP 7075 Al deformed at 1×10−3 s−1 [18].

7

(a)

(b) 30

50 500℃

=1×10 s

-2 -1

=1×10 s

-1 -1

40

450℃ 475℃ 500℃ 525℃ 550℃

25

=5×10 s

-2 -1

=1×10 s

Ture Stress (MPa)

Ture Stress (MPa)

-2 -1

=5×10 s

-3 -1

30

=1×10 s

-3 -1

20

10

0 0.0

20

15

10

5

0.5

1.0

1.5

2.0

2.5

0 0.0

3.0

0.5

1.0

1.5

2.0

2.5

3.0

Ture Strain

Ture Strain

Fig. 3 Effect of (a) strain rate and (b) temperature on true stress - true strain curves for

the Al-Zn-Mg-0.25Sc-0.10Zr alloy.

3.2.2 Strain rate sensitivity parameter Strain rate sensitivity, m, is an important parameter in superplastic deformation and can be used to characterize the capacity of an alloy in resisting necking spread. Similarly to the creep deformation of alloys (e.g. reported by Langdon [19]), superplastic deformation is usually described by an equation for power-law creep of the form:

D Gb  b     A 0     RT  d   G  

p

1

m

 Q  exp     RT 

(1)

where A and p are empirical constants, G is the shear modulus, b is the Burgers vector, R is the gas constant, T is the absolute temperature, D0 is the frequency factor, m is the strain rate sensitivity exponent (1/m is the stress exponent n, n=1/m.) and Q is the activation energy of an appropriate diffusion process. For any selected material and testing conditions, the values of m and Q in Eq. (1) can describe the deformation mechanism. The strain rate sensitivity parameter, m, is defined as

m

 log 

(2)



 log 

T

The flow stress (σ) values are obtained by taking the data at a true strain of ~0.7 (the peak

8

values) under different superplastic deformation conditions. Fig. 4 shows the variation of flow stress 

(σ) with an initial strain rate (  ) for the Al-Zn-Mg-Sc-Zr tensile specimens. It is interesting to note that the stress (σ) decreases with increasing temperatures in the Al-Zn-Mg-Sc-Zr alloys, which is consistent with the results in Fig. 3(b). According to the Eq. (2) and Fig. 4, it can be concluded that

Flow Stress (MPa)

the average value of strain rate sensitivity (m) parameters is ~0.37. 10

2

10

1

0.37

450℃ 475℃ 500℃ 525℃ 550℃ 10

0

10

-4

10

-3

10

-2

10

-1

10

0

-1

Strain Rate (s ) Fig. 4 Variation of logσ as a function of log  for various testing temperatures in the Al-Zn-Mg-0.25Sc-0.10Zr alloy. 

3.2.3 Activation energy Superplastic deformation is a diffusion controlled process. To further elucidate the temperature dependence of the strain rate in the Al-Zn-Mg-0.25Sc-0.10Zr tensile test, the calculation of Q value is assessed for all the experimental conditions of strain rate and temperature. Neglecting the G1-1/m/T term in Eq. (1) the apparent activation energy Q may be evaluated as

Q

R ln m 1 T

 

(3) 



Fig. 5 shows lnσ vs. 1/T for various strain rate levels in the Al-Zn-Mg-Sc-Zr alloy. The activation energy at various strain rates could be calculated from Eq. (3) and Fig. 5, which is documented in Table 2. Moreover, the average value of the activation energy is approximately 84.5

9

kJ/mol, which is close to that for grain boundary self-diffusion of Al (84 kJ/mol=QGB) [20].

5

In(σ) (MPa)

4

3

2

=1×10 s

-1 -1

=5×10 s

-2 -1

=1×10 s

-2 -1

1

=5×10 s

-3 -1

=1×10 s

-3 -1

0 0.00120

0.00125

0.00130

0.00135

0.00140

-1

1/T (K )

Fig. 5 lnσ vs. 1/T for various strain rate levels in the Al-Zn-Mg-0.25Sc-0.10Zr alloy.

Table 2 The activation energy (Q) for various strain rate levels in the Al-Zn-Mg-0.25Sc-0.10Zr alloy. Strain rate [s-1]

1×10−3

5×10−3

1×10−2

5×10−2

1×10−1

Activation energy [KJ/mol]

85.0

84.9

84.3

84.2

83.8

3.3 EBSD analysis 3.3.1 Al-Zn-Mg alloy Fig. 6 shows EBSD analyses of the Al-Zn-Mg alloy deformed at different conditions. The color of each grain is coded by its crystal orientation based on [001] inverse pole figure. Fig. 6(a) shows that the microstructure consists mainly of large grains. The average grain size is more than 10 μm. The misorientation angle histogram of the untested sample shown in Fig. 6(b) indicates that the alloy mainly consists of HABs, and the fraction of HABs is 95.8%. The theoretical distribution of grain boundary misorientation angles for a random grain assembly predicted by Mackenzie [21] is also shown by a black solid line in Fig. 6(b). Fig. 6(c) (d) show EBSD maps of the Al-Zn-Mg alloy

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deformed at 500 ℃ and strain rates of 5×10−3 s-1~1×10−2 s-1 after interrupting the tensile test at a strain value of 0.69 (corresponding to 100% elongation), respectively. Compared with the microstructures of the sample before deformation (Fig.6(a)), it can be inferred from Fig.6(c)(d) that the grains are gradually elongated in the tensile direction and some small equiaxed grains are generated near old deformed grains, indicating that recrystallization occurs. Besides, the grain growth is very obvious, indicating that the Al-Zn-Mg alloy shows very poor stability of the microcrystalline structure during superplastic deformation. The grain sizes are still larger than 10 μm. However, the major prerequisites for achieving structural superplasticity are a fine grain size (typically less than 10 μm) and the thermal stability of the fine microstructure at high temperatures [1]. Therefore, the Al-Zn-Mg alloy exhibits no superplasticity because of the larger grain size (>10 μm) and the poor stability of the microcrystalline structure during superplastic deformation. 3.3.2 Al-Zn-Mg-0.25Sc-0.10Zr alloy Fig. 7 shows the EBSD images of the Al-Zn-Mg-0.25Sc-0.10Zr alloy deformed at 500 ℃and 1 × 10-2 s-1 after interrupting the tensile test at different strain, and the corresponding misorientation angle distribution are shown in Fig. 8. Fig. 7(a) shows the initial microstructures of the Al-Zn-Mg-0.25Sc-0.10Zr alloy. Fig. 7(a) shows that the microstructure consists mainly of LABs. The new grains entirely outlined by HABs form near old deformed grains. The partially recrystallized microstructure with a uniform and equiaxed grain distribution is detected, and the average grain size is determined to be 2.82 μm. The misorientation angle histogram of the untested sample shown in Fig. 8(a) indicates that the number fraction of HABs is 27.3%, indicating that the alloy mainly consists of subgrains.

11

(b)

(a)

0.06

fHABs=95.8% 0.05

Frequency

0.04

0.03

0.02

0.01

0.00 10

20

30

40

50

60

Misorientation Angle (degree)

(d)

(c)

TD

TD

RD

RD

Fig. 6 EBSD analyses of the Al-Zn-Mg alloy deformed at different conditions: (a) the untested alloy(ε=0);

(b) misorientation angle distribution of the untested alloy, solid line represents Mackenzie [21] distribution; (c) 500 ℃, 5 × 10-3 s-1, ε=0.69; (d) 500 ℃, 1 × 10-2 s-1, ε=0.69.

With the increase of deformation degree (ε), besides, it can be inferred from Fig.7(b)~(d) and Fig. 8(b)~(d) that the grains gradually transfer into equiaxed grains and the fraction of HABs significantly increases, indicating that dynamic recrystallization occurs. The occurrence of dynamic recrystallization during superplastic deformation leads to complete transformation of LABs into

12

HABs [22]. When the amount of deformation up to a strain value of 2.40 at 500 ℃ and 1× 10-2 s-1, the grains completely transfer into uniform equiaxed grains (Fig. 7(d)) having an average size of 4.79 μm, and the misorientation angle histogram (Fig. 8(d)) is close to random misorientation distribution predicted by Mackenzie [21]. Same time, the fraction of HABs reaches 98.8%. This means the alloy is completely recrystallized under these conditions.

(a)

(b)

TD

TD

RD

RD (d)

(c)

TD

TD RD

RD

Fig. 7 EBSD analyses of the Al-Zn-Mg-0.25Sc-0.10Zr alloy deformed at 500 ℃and 1 × 10-2 s-1 after interrupting the tensile test at different strain: (a) ε=0, (b) ε=0.69; (c) ε=1.10; (d) ε= 2.40.

13

(b)

(a)

0.18

0.30

fHABs=27.3%

fHABs=43.5%

0.16

0.25 0.14 0.12

Frequency

Frequency

0.20

0.15

0.10 0.08 0.06

0.10

0.04

0.05 0.02 0.00

0.00 10

20

30

40

50

10

60

20

30

40

50

60

50

60

Misorientation Angle (degree)

Misorientation Angle (degree)

(c)

(d)

0.06

0.05

fHABs=75.9%

fHABs=98.8%

0.05

0.04

Frequency

Frequency

0.04

0.03

0.03

0.02

0.02 0.01

0.01

0.00 10

20

30

40

50

0.00

60

10

Misorientation Angle (degree)

20

30

40

Misorientation Angle (degree)

Fig. 8 Misorientation angle distribution of the Al-Zn-Mg-0.25Sc-0.10Zr alloy deformed at 500 ℃and 1 × 10-2 s-1 after interrupting the tensile test at different strain: (a) ε=0; (b) ε=0.69; (c) ε=1.10; (d) ε= 2.40.

EBSD images of the Al-Zn-Mg-0.25Sc-0.10Zr alloy deformed at different conditions after interrupting the tensile test at a strain value of 1.10 (corresponding to 200% elongation) are presented in Fig. 9. Fig. 9 shows that dynamic recrystallization occurs during superplastic deformation, and the recrystallized grains are generated. Furthermore, grain boundary character distributions data and average grain sizes of the samples in Fig. 7 and Fig. 9 are listed in Table 3. Compared with Fig. 7 and Fig. 9, it can be observed that the recrystallized grains gradually coarsen with increasing the deformation temperatures or decreasing the strain rate or increasing the deformation degree. However, The final grain structure has an average recrystallized grain size of

14

less than 5 μm under all the superplastic deformation conditions studied in this work, revealing that the Al-Zn-Mg-0.25Sc-0.10Zr alloy has stronger ability to inhibit grain growth than other alloys [12, 13]. This is an important reason why the studied alloy is highly superplastic in a wide temperature range.

(a)

(b)

TD

TD RD

RD (c)

(d)

TD

TD RD

RD

Fig. 9 EBSD analyses of the Al-Zn-Mg-0.25Sc-0.10Zr alloy deformed at different conditions after interrupting the tensile test at a strain value of 1.10: (a) 475 ℃ and 1 × 10-2 s-1, (b) 525 ℃ and 1 × 10-2 s-1; (c) 500 ℃ and 1 × 10-1 s-1; (d) 500 ℃ and 1 × 10-3 s-1.

15

Table 3 Grain boundary character distributions and grain sizes of the Al-Zn-Mg-0.25Sc-0.10Zr alloy

at deformed condition

Deformed condition

HABs(%)

Grain size (μm)

untested

27.3

2.82

475 ℃,  =1×10−2 s-1 and ε=1.10

74.5

3.53

65.7

4.48

43.5

3.67

75.9

4.46

98.8

4.79

77.3

3.84

71.8

3.96





500 ℃,  =1×10−3 s-1 and ε=1.10 

500 ℃,  =1×10−2 s-1 and ε=0.69 

500 ℃,  =1×10−2 s-1 and ε=1.10 

500 ℃,  =1×10−2 s-1 and ε=2.40 

500 ℃,  =1×10−1 s-1 and ε=1.10 

525 ℃,  =1×10−2 s-1 and ε=1.10

3.4 Microtexture evolution Fig. 10 shows the microtexture evolution of the Al-Zn-Mg-0.25Sc-0.10Zr alloy during superplastic deformation at the optimal deformation parameters (500℃ and 1×10−2 s-1). Fig. 10(a) shows that the texture of the alloy is composed of a weak Cube component {001}<100> and a strong β-fiber rolling texture, which runs from the Copper orientation {211}<111> through the S orientation {123}<634> to the Brass orientation {011}<211>. Moreover, the intensity of the overall texture decreases and β-fiber texture gradually weakens and transfers into a weaker Goss component {011}<100> and random texture with increasing deformation, as shown in Fig. 10(b)~(d). As is widely known, the slip system for aluminum is (111)[110] and the Taylor factors (TF) of the β-fiber rolling texture consisting of Copper, Brass and S components are larger than that of the cube

16

component. Therefore, the disappearance of β-fiber rolling texture and the randomization of the grain orientation can accelerate more homogeneous grain boundary sliding, leading to a higher superplasticity. (a)

(b)

(c)

(d)

Fig. 10 ODFs representative sections of the Al-Zn-Mg-0.25Sc-0.10Zr alloy deformed at 500 ℃and 1 × 10-2 s-1 with different strain: (a) ε=0, (b)ε=0.69; (c)ε=1.10; (d) ε= 2.40.

3.5 Al3ScxZr1-x particles Fig. 11 shows the transmission electron micrograph obtained from the gage portion of the

17

Al-Zn-Mg-0.25Sc-0.10Zr alloy sample tested at 500 ℃ and 1×10−2 s-1 after interrupting the tensile test at a strain value of 1.10 (corresponding to 200% elongation). It can be observed that lots of nano-scaled particles with Ashby-Brown contrast can be observed. Combined with their superstructure reflections like {001} and {1-10} in the [110]Al projection (detailed in the inset Fig. 11(a)), it can be confirmed that these particles are Al3ScxZr1-x particles, which are coherent with the Al matrix [23]. Those coherent particles strongly pin dislocations and grain boundaries, as arrowed in Fig. 11(a) and (b), leading to the controlled growth of the micro-scaled grains, verified in Section 3.3.2.2, and high superplasticity.

(a)

(b)

Al3ScxZr1-x Al3ScxZr1-x pins dislocation

Al3ScxZr1-x

Al3ScxZr1-x pins grain boundaries

Fig. 11 TEM microstructures of the Al-Zn-Mg-0.25Sc-0.10Zr alloy deformed at 500 ℃ and 1 × 10-2 s-1 after interrupting the tensile test at ε=1.10.

To further investigate those particles, several dark field TEM images were recorded for the Al-Zn-Mg-0.25Sc-0.10Zr alloy deformed at 500 ℃ and 1 × 10-2 s-1 with different strain under selected conditions. These {100} dark field images are shown in Fig. 12. Fig. 12(a)(b) show the dark field images of the untested alloy. In this condition, the Al-Zn-Mg-0.25Sc-0.10Zr alloy contains

18

sphere Al3ScxZr1-x precipitates having sizes within the range of 20~30 nm. After the alloy deformed at 500 ℃ and 1 × 10-2 s-1 with the strain value of 2.40 (corresponding to 1000% elongation), it can be seen that Al3ScxZr1-x particles almost keep an unchangeable size (20~30 nm). Therefore, it is concluded that the Al-Zn-Mg-Sc-Zr alloy exhibits superior grain stability because of the presence of Al3ScxZr1-x precipitates that do not exhibit noticeable coarsening during the process of deformation.

(a)

(b)

(c)

(d)

Fig. 12 Dark field images of the Al-Zn-Mg-0.25Sc-0.10Zr alloy deformed at 500 ℃and 1 × 10-2 s-1 with different strain: (a)(b) ε=0, (c)(d)ε=2.40.

19

3.6 Surface topography of deformed specimens (a)

(b)

cavitations

(c)

(d)

fibers

Fig. 13 SEM images of tensile specimens deformed to failure at 500 ℃ and 1 × 10-2 s-1: (a) cavitations, (b) grain

boundary sliding, and (c)(d) the fracture morphology.

To further verify the superplastic deformation mechanism of the Al-Zn-Mg-0.25Sc-0.10Zr alloy, SEM observations of the fracture position of the samples failed at 500 ℃ and 1 × 10-2 s-1 were performed and the results are shown in Fig. 13. Remarkable cavitation is found on the sample surface. Besides, most cavities grow mainly along the tensile direction and have an irregular and jagged shape suggesting the operation of a plasticity-controlled cavity growth mechanism [24], as can be observed in Fig. 13(a). The volume fraction of the cavitations is much less than the other reports [13, 18, 25]. This is an important factor that the Al-Zn-Mg-0.25Sc-0.10Zr alloy has very high superplasticity. Fig. 13(b) shows that the surface microstructure indicates that extensive GBS along

20

with somewhat elongated grains and equiaxed grains occurred during superplastic deformation (Fig. 13(b)). A similar phenomenon was reported in FSP 7075Al-T651 rolled plates deformed at strain rate range of 1×10−5~1×10−2 s−1 and at temperature range of 200~350 ℃ [26]. Observed from Fig.13(c), the fracture mechanism of the Al-Zn-Mg-0.25Sc-0.10Zr alloy is intergranular fracture. Furthermore, some fibers (whisker) are detected among sliding grains. Fiber (whisker) formation is generally thought to be the evidence of the existence of liquid phases along grain boundaries at high temperatures [27]. However, further study is required to elucidate the origin of the fibers. But it should be pointed out that the existence of fiber (whisker) could be an indirect evidence for GBS.

4、Discussion 4.1 Superplasticity deformation mechanism It is generally accepted that grain boundary sliding (GBS) is the predominant deformation mechanism during superplastic flow for most of fine-grained materials and is characterized by a strain rate sensitivity of ~0.5. However, a previous study showed that ‘‘Rachinger GBS” might be a plausible mode of deformation for FSP Al-Zn-Mg-Sc alloy with a strain rate sensitivity of ~0.33 [12]. Furthermore, Islamgaliev et al. suggested that the GBS is an important deformation process in 1421 aluminum alloy, though the measured strain rate sensitivity of ~0.2-0.3 is lower than the anticipated value for conventional superplasticity [28, 29]. For the Al-Zn-Mg-0.25Sc-0.10Zr alloy, the large elongations (Fig. 1(b)) and the average strain rate sensitivity of ~0.37 (Fig. 4) are observed in the strain rate range of 1 × 10-3 ~ 1 × 10-1s-1 for the temperatures ranging from 450 to 550 ℃, which imply that GBS plays an important role in the process of superplastic deformation of the Al-Zn-Mg-0.25Sc-0.10Zr alloy. What’s more, the average value of the activation energy is about 84.5 kJ/mol (Fig. 5), which are close to that for grain boundary self-diffusion of Al alloy (QGB), according to the GBS models of Mukherjee [30] and Ball

21

and Hutchinson [31]. Additionally, SEM examinations further revealed the distinct evidence of the GBS on the surface of the tensile specimens superplastic deformed at 500 ℃ and 1 × 10-2 s-1 (Fig. 13(b)). Therefore, those results mean that GBS is the main superplastic deformation mechanism for the Al-Zn-Mg-0.25Sc-0.10Zr alloy. 4.2 The transformation of LABs to HABs The initial microstructure of the Al-Zn-Mg-0.25Sc-0.10Zr alloy contains basically a high amount of LABs (Fig. 7(a)). However, LABs are generally believed to be not suitable for GBS owing to their low orientation differences, while it is well known that random (disordered) HABs are desirable for superplastic flow by GBS [32, 33]. Dynamic recrystallization during superplastic deformation leads to the transformation of LABs to HABs. Several mechanisms describing dynamic recrystallization involve either the rotation [34] or switching [35] of subgrains, or both. Subgrain rotation, in turn, can result from boundary sliding, or from dislocation slip on selected sets of slip systems within each grain. This leads to a gradual increase in the subgrain boundary angle and eventually into HABs. The same discussion can be applied to the transformation of LABs to HABs during superplastic deformation of the Al-Zn-Mg-0.25Sc-0.10Zr alloy. As discussed earlier, the initial structure of the current alloy consists of many subgrains characterized by a high ratio of LABs (Fig. 7(a)). Those subgrains are composed of unstable dislocation cells. During the preheat in each test, dislocations in the dislocation cell wall quickly interact and annihilate with each other; this results in the formation of stable subgrains. Upon mechanical loading, dislocations with a favored Schmid factor begin to move across subgrains on selected sets of slip systems, resulting in an increase in the sub-boundary angle. These moving dislocations are expected to be pinned by Al3ScxZr1-x particles at the sub-boundary (Fig. 11). As this process continues, subgrain boundaries eventually convert to the true HABs. It was found that the fraction of HABs in the tested specimen deformed with ε=2.40 (1000% elongation) at 500 ℃ and 1×10−2 s-1 (Fig. 8(d)) is as high as 98.8%, significantly higher than the 27.3% in untested specimen (Fig. 8(a)). Such a microstructure is a typical structure facilitating the occurrence of GBS in the initial superplastic stage. 4.3 The role of Al3ScxZr1-x particles in superplastic deformation 22

The transformation of LABs to HABs usually resulted in a significant grain growth, which reduced the superplastic elongation. Considering the existence of a high density of fine Al3ScxZr1-x particles dispersed both at the grain boundaries and within the grain interiors (Fig.11), the grain growth at high temperatures in the Al-Zn-Mg-0.25Sc-0.10Zr is slow. The Al3ScxZr1-x dispersoids have been proven to be most effective in pinning grain and subgrain boundaries inhibiting recrystallization grain growth during mechanical and thermal processing [36]. Moreover, comparing Fig. 12(b) with Fig. 12(d), it should be noted that the Al3ScxZr1-x particles did not exhibit noticeable coarsening during the process of deformation. Such the fine and stable microstructure is suitable for superplastic deformation at these temperatures. It is well documented that when the GBS is the dominant mechanism of superplastic deformation in the superplastic materials, the sliding needs to be accommodated by other mechanisms, such as dislocation slip. When the GBS cannot be well accommodated, stress concentration at certain sites, such as the grain triple junctions, and boundary ledges, may cause the development of cavitation at the sites [37]. Moreover, grain growth occurring during superplastic flow that leads to flow hardening can also cause additional cavity nucleation in a continuous manner [38].

Cavities

were

observed

continuously

during

deformation

in

the

present

Al-Zn-Mg-0.25Sc-0.10Zr alloy (Fig. 13(a)); but the volume fraction of cavities are smaller than the previous reports, which can be attributed to the stable fine-grained structure and enhanced uniformity of microstructure. That micro structural stability of the present material might be ascribed to the existence of numerous fine Al3ScxZr1-x dispersoids that effectively impede the growth of the grains to suppress cavity formation during mechanical and thermal processing. Furthermore, additional interfacial energy of the Al3ScxZr1-x particles can contribute to the stability of the cavity. This is also very important to achieve excellent superplastic properties in Al-Zn-Mg-0.10Sc-0.10Zr alloy.

5. Conclusion The Al-Zn-Mg alloy showed no superplasticity because of the larger grain size (>10 μm) and

23

the poor stability of the microcrystalline structure during superplastic deformation. However, the Al-Zn-Mg-0.25Sc-0.10Zr alloy exhibited excellent superplastic (elongations of ≥ 500%) at a wide temperature range of 450 ~ 550 ℃ and high strain rate range of 5×10−3 ~ 5×10−2 s−1. The maximum elongation of ~1523 % was achieved at 500 ℃ and 1×10−2 s−1. Enhanced superplasticity of the Al-Zn-Mg-0.25Sc-0.10Zr alloy was attributed to the complete transformation of LABs to HABs due to the occurrence of dynamic recrystallization and the presence of coherent Al3ScxZr1-x particles that effectively impede the growth of the grains. Besides, β-fiber rolling textures weakened, and random textures were predominant in the superplastic deformed alloy. On the other hand, the average strain rate sensitivity of ~0.37 was observed during deformation, and the activation energy of ~84.5 kJ/mol was very close to that for grain boundary self-diffusion of Al, indicating that GBS was the controlling deformation mechanism.

Acknowledgments This work was financially supported by the National General Pre-research Project of China (51312010402) and the China Postdoctoral Science Foundation (2014M552149).

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