metals

Article Superplastic Tensile Deformation Behavior and Microstructural Evolution of Al–Zn–Mg–Cu Alloy

Guangyu Li 1, Hua Ding 1,*, Jian Wang 1, Ning Zhang 2 and Hongliang Hou 2 1 School of Materials Science and Engineering, Northeastern University, Shenyang 110819, China 2 Beijing Aeronautical Manufacturing Technology Research Institute, Beijing 100024, China * Correspondence: [email protected]; Tel.: +86-139-9987-6262



Received: 30 July 2019; Accepted: 23 August 2019; Published: 28 August 2019


Abstract: The microstructural evolution of the Al–Zn–Mg–Cu alloy during the superplastic deformation process has been studied by high temperature tensile experiment. The superplastic deformation behaviors are investigated under different temperatures of 470 ◦C, 485 ◦C, 500 ◦C, 515 ◦C and 530 C, and different strain rates of 3 10 4 s 1, 1 10 3 s 1, 3 10 2 s 1 and 1 10 2 s 1. ◦ × − − × − − × − − × − − The microstructure observation shows that uniform and equiaxed grains can be obtained by dynamic recrystallization in the initial stage of superplastic deformation. Once the recrystallization process has been finished, the variations of the fraction of high angle boundary, the grain aspect ratio and the Schmid factor are negligible during the superplastic deformation, which shows that the grain boundary sliding and grain rotation are the main deformation mechanisms. The maximum texture intensity decreases compared with the initial microstructure, indicating that grain boundary sliding and grain rotation can weaken the texture, however, the texture intensity increases in the final stage of superplastic deformation, which may be resulted from the stress concentration.

Keywords: Al–Zn–Mg–Cu alloy; microstructure evolution; superplastic

1. Introduction Al–Zn–Mg–Cu alloy, as a traditional high strength engineering material, is widely applied in the airplane and aerospace industry on account of its favorable specific strength, excellent abrasion resistance and stress corrosion resistance [1]. Precipitation hardening is the main strengthening mechanism for this alloy [2,3]. The relatively high contents of alloy composition (Mg, Zn and Cu) in this alloy results in high yield strength and tensile strength, but deteriorates its cold deformation workability in the meantime. It is a problem in the application of this alloy to form components with complex geometry because of a poor cold forming ability of the material. Fortunately, superplastic forming can be used to effectively solve the above problem. Superplastic deformation has a significant industrial application value due to the large elongation and excellent deformation ability under suitable conditions [4]. In addition, both the thermal deformation mechanism and the flow behavior of this alloy during superplastic forming are rather complex. The microstructural evolution and deformation mechanism are different at the various experimental conditions and initial structures, for example grain boundary sliding (GBS) [5], the dynamic recovery (DRV) [6], dynamic recrystallization (DRX) [7], GBS controlled by dislocation movement [8], etc. Hence, a systematic study of microstructural evolution and mechanical behavior of this alloy is of great significance. Numerous researchers have investigated the microstructural evolution and mechanical behavior during superplastic deformation of Al–Zn–Mg–Cu alloys. For instances, Mukhopadhyay [9] utilized the two-step strain rate method to increase the elongation of the Al–Zn–Mg–Cu alloy from 650% to 916% at a temperature of 425 C and in the strain rate ranging from 1.9 10 2 s 1 to 3 10 5 s 1 of the ◦ × − − × − − Al–Zn–Mg–Cu-Zr alloy. Kumar [10] examined the influence of Sc on the superplastic property of the

Metals 2019, 9, 941; doi:10.3390/met9090941 www.mdpi.com/journal/metals Metals 2019, 9, 941 2 of 11

AA 7010 alloy and obtained the unrecrystallized grain structure, then superplastic deformation was applied on the material at a temperature and strain rate combination of 1.9 10 2 s 1 and 475 C to × − − ◦ achieve the recrystallized grain structure and the elongation of 650% was attained. Li [11,12] studied the superplasticity of the 5A70 aluminum alloy with a grain size of 8.48 µm at different temperatures and also investigated the cavity nucleation mechanism in this alloy. The superplasticity of the Al–Zn–Mg–Sc–Zr alloy subjecting to a traditional thermal–mechanical processing was studied at 500 ◦C and 5 10 3 s 1, and it was found that the low angle grain boundaries (LAGBs) gradually transferred × − − into the high angle grain boundaries (HAGBs) and the texture intensity was weakened during the Superplastic deformation process [13]. Sun [14] simulated the continuous DRX of Al–Zn–Mg–Cu alloys during superplastic deformation with the internal-state-variable method and developed a new DRX model of the Al–Zn–Mg–Cu alloy. Mosleh [15] investigated the effect of strain rate and temperature on the superplastic deformation behavior of Ti-4%V-6%Al, Ti-3%Mo-1%V-4%Al and Ti-1.8%Mn-2.5%, the results showed that temperature was a less significant factor than the strain rate. Qin [16] investigated the grain refinement method in a high entropy alloy. Mikhaylovskaya [17] studied the effect of isothermal multi-directional forging on the superplasticity of the Al–Mg-based alloy. Masuda et al. [18] have characterized GBS and grain rotation mechanism with a three-dimensional method in the Al–Zn–Mg–Cu alloy 3D-electron back scattering diffraction (EBSD) tomography during the superplastic deformation. As mentioned above, a number of reports on the superplastic deformation behaviors and microstructural evolution of Al–Zn–Mg–Cu alloy have been published, However, there still lacks systematic investigation on the effect of strain on microstructure evolution of this kind of material. The aim of this work is to study the evolution of microstructure under various strains and study the mechanical properties of the Al–Zn–Mg–Cu alloy at different deformation conditions. Moreover, the superplastic deformation mechanism is also evaluated based on mechanical behavior and microstructure characterization. This work could provide the guidelines for superplastic deformation applications for the Al–Zn–Mg–Cu alloy.

2. Materials and Methods Table1 shows the alloying component of the investigated Al–Zn–Mg–Cu alloy (corresponding to GB/T 3190-2008 (China)). A uniaxial superplastic tensile experiment was carried out on a LETRY-200kN electronic universal mechanical tensile machine (Xi0an Lichuang material testing technology co. Ltd., Xi’an, China). The thickness of the tensile sample was 2 mm, the gauge width was 6 mm and the gauge length was 10 mm. The experimental material was achieved by the following steps. Firstly, the initial plate with a thickness of 25 mm was solution treated at 470 ◦C for 4 h, then water quenched. Secondly, aging treatment was applied on the plate at 350 ◦C for 40 h. Thirdly, a 2 mm thickness plate rolled by a multi-pass was annealed at 300 ◦C for 2 h. Finally, the studied alloy with the average grain size of 10 µm was achieved through annealing in the conditions of 480 ◦C and 30 min in a salt bath furnace. The final thickness of Al–Zn–Mg–Cu was 2 mm. The specimens were sectioned parallel to the rolling direction. The specimens were deformed at different temperatures of 470 ◦C, 485 ◦C, 500 C, 515 C and 530 C, and various strain rates of 3 10 4 s 1, 1 10 3 s 1, 3 10 2 s 1 and ◦ ◦ ◦ × − − × − − × − − 1 10 2 s 1. Temperature measuring instruments were placed in several places in the furnace to keep × − − the temperature constant. The samples were quenched immediately with water after fracture in order to maintain the superplastic deformation microstructure. Samples were taken from the deformation zone in the middle of the specimen and then were mechanically polished by a 0.5 µm diamond paste for EBSD study and the metallographic etchant is Keller’s reagent (190 mL water, 2 mL hydrofluoric acid, 3 mL hydrochloric acid and 5 mL nitric acid) for metallographic observation. The microstructure observations were carried out on a Zeiss ULTRA 55-type field emission scanning electron microscopy (Carl Zeiss, Oberkoichen, Germany) operated at 20 kV. HKL channel 5 software (Oxford Instruments, London, UK) was brought in to process the data of EBSD. Metals 2019, 9, 941 3 of 11

Table 1. Chemical composition of Al–Zn–Mg–Cu alloy in this research (mass fraction/%).

Zn Mg Cu Mn Ti Fe Si Al Metals 2019, 9, x FOR PEER REVIEW 3 of 11 5.96 2.22 1.60 0.4 0.04 0.06 0.03 others 3. Results and Discussion 3. Results and Discussion 3.1. Tensile Deformation Behavior 3.1. Tensile Deformation Behavior The mechanical behavior of the studied alloy at various experiment conditions is presented as shownThe in mechanical Figure 1. Temperature behavior of theand studied strain rate alloy can at a variousffect significantly experiment the conditions flow stress is in presented Figure 1a,b as . shownThe flow in Figure stress1 . increases Temperature with and the strain strain rate rate can increasing a ffect significantly and the thetemperature flow stress decreasing. in Figure1 a,b. The Theobservation flow stress of increases microstructural with the evolution strain rate is increasing under the and experimental the temperature condition decreasing. of 530 The°C and observation an initial ofstrain microstructural rate of 3 × 10 evolution−4 s−1, under is under which the the experimental elongation reaches condition the of maximum 530 ◦C and (1020%). an initial The strain elongations rate of 4 1 3to fracture10− s− ,vary under from which 400% the to elongation 1020% under reaches the temperature the maximum from (1020%). 470 °C The to 530 elongations °C at 3 × to10− fracture4 s−1. The × 4 1 varytrue from stress 400%–strain to 1020% curve under in the the optimum temperature condition from and470 ◦ someC to 530pictures◦C at 3 representing10− s− . Thedifferent true × stress–straindeformation curve extent in during the optimum superplastic condition deformation and some labeled pictures by T0, representing T1, T2, T3, diT4ff areerent shown deformation in Figure extent3c. T0 during is the original superplastic sample, deformation and T1, T2, labeled T3 and by T4 T0, are T1, corresponding T2, T3, T4 are shown to the instrains Figure with1c. T00.41, is the0.69, original1.61 and sample, 2.17. and T1, T2, T3 and T4 are corresponding to the strains with 0.41, 0.69, 1.61 and 2.17.

Figure 1. The curves of flow stress at (a) different strain rates at the same temperature and (b) different temperaturesFigure 1. The at curves the same of flow strain stress rate. at ( c()a) The different stress-true strain strain rates curveat the atsame 530 temperatureC, 3 10 4 sand1 and (b) d sampleifferent ◦ × − − photostemperatures after each at tensilethe same stage strain marked rate. ( withc) The T1, stress T2, T3,-true T4 strain and T5, curve respectively. at 530 °C, 3 × 10−4 s−1 and sample photos after each tensile stage marked with T1, T2, T3, T4 and T5, respectively. The influence of deformation parameters on the peak stress is presented in Figure2a. The peak stress decreasedThe influence with of the deformation temperature parameters increasing on and the with peak the stress strain is rate presented increasing. in Figure In addition, 2a. The ahigh peak temperaturestress decrease cand increase with the the temperature atomic activation increasing energy and [ 19with] and the reduce strain therate internal increasing. force In required addition, for a ofhigh GBS temperature [12]. Meanwhile, can increase the low the strain atomic rate allowed activation enough energy time [19] to and neutralize reduce the the e ff internalect of work force hardeningrequired for and of decreased GBS [12]. the Meanwhile, critical shear the stress low strain of GBS. rate The allow elongationsed enough at fracture time to ofneutralize this studied the alloyeffect underof work various hardening experiment and decrease conditionsd the are critical presented shear in stress Figure of2b. GBS. The The results elongations show that at the fracture elongation of this 1 increasedstudied alloy with under the temperature various experiment increasing, conditions except the are strainpresented rate ofin 0.01Figure s− 2bwhere. The results the elongations show that the elongation increased with the temperature increasing, except the strain rate of 0.01 s−1 where the elongations at fracture suddenly decreased at 530 °C, which might be because the incipient melting leads to the liquid phase, thus the sample fractured easily at the high temperature. With the strain rate decreasing, the elongations at fracture increase obviously from 0.003 s−1 to 0.001 s−1 and changed slightly from 0.001 s−1 to 0.0003 s−1. The improvement for the elongation was limited when the strain rate was low.

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at fracture suddenly decreased at 530 ◦C, which might be because the incipient melting leads to the liquid phase, thus the sample fractured easily at the high temperature. With the strain rate decreasing, 1 1 the elongations at fracture increase obviously from 0.003 s− to 0.001 s− and changed slightly from 1 1 Metals0.001 2019 s− ,to 9, 0.0003x FOR PEER s− . REVIEW The improvement for the elongation was limited when the strain rate was4 of low. 11

-1 1200 -1 25 Strain rate－ 0.0003s Strain rate－ 0.0003s -1 (b) (a) Strain rate－ 0.001s -1 -1 Strain rate－ 0.001s Strain rate－ 0.003s -1 1000 -1 Strain rate－ 0.01s Strain rate－ 0.003s 20 -1 Strain rate－ 0.01s 800

15

600

10 400

Peak stress (Mpa) Peakstress Elongationfracture/% to 200 5

0 470 480 490 500 510 520 530 470 480 490 500 510 520 530 Temperature Temperature Figure 2. ((aa)) The The e effectffect of of deformation deformation parameters on the peak stress stress and ( b) t thehe effectffect of deformation parameters on the elongation.elongation.

3.2. Superplastic C Characteristicharacteristic P Parametersarameters 3.2.1. M-Value 3.2.1. M-Value Strain rate sensitivity index m is significantly effective in studying the superplastic mechanical Strain rate sensitivity index m is significantly effective in studying the superplastic mechanical property, which represents the relationship between flow stress and strain rate during superplastic property, which represents the relationship between flow stress and strain rate during superplastic deformation process and reflects the anti-necking ability of the Al–Zn–Mg–Cu alloy. Backofen [20] deformation process and reflects the anti-necking ability of the Al–Zn–Mg–Cu alloy. Backofen [20] proposed the equation of the relationship between strain rate and flow stress in the superplastic proposed the equation of the relationship between strain rate and flow stress in the superplastic deformation process: deformation process: . m σ = Kε , (1) . m where σ: flow stress; ε: strain rate; m: strain rate= K sensitivity , index and K is a constant related(1 to) deformation conditions and material properties. The natural logarithm and differential on both sides wofhere Equation : flow (1) can stress; be used :to strain calculate rate; them: m-value.strain rate The sensitivity fitting relationship index andfor K i calculatings a constant the related m value to deformation conditions and material properties. The natural logarithm and differential on both sides is shown in Figure3. The m value at a temperature of 470 ◦C, 485 ◦C, 500 ◦C, 515 ◦C and 530 ◦C were of0.33, Equation 0.38, 0.50, (1) 0.62can andbe used 0.63 respectively.to calculate the The m m-value. value ofThe the fitting superplastic relationship material for is calculating generally 0.3–0.9,the m valuethat is, is when shown the in mFigure value 3. is The greater m value than at 0.3, a temperature it is considered of 470 that °C the, 485 studied °C , 500 material °C , 515 ° hasC and favorable 530 °C weresuperplasticity 0.33, 0.38, under0.50, 0.62 the and corresponding 0.63 respectively. deformation The m conditions.value of the Then superplastic the value material is closely is relatedgenerally to 0.3the– contribution0.9, that is, when to GBS. the m value is greater than 0.3, it is considered that the studied material has favorableAccording superplasticity to the calculated under the results, corresponding the value deformation of m increased conditions gradually. The withn the the value deformation is closely related to the contribution to GBS. temperature increasing, and reached 0.63 at 530 ◦C, indicating that the material has the best superplastic deformationAccording ability to the at calculated this temperature. results, the The value m value of m and increase the elongationd gradually at with fracture the deformation significantly temperatureincreased with increasing, the increase and of temperature. reached 0.63 In at addition, 530 °C , theindicating stress exponent that the n material(n = 1/m ) has is one the of best the superplastic characteristicdeformation ability parameter at this to determine temperature. the The deformation m value mechanisms.and the elongation It was at suggested fracture significantlythat the superplastic increase deformationd with the increase mechanism of temperature. might be GBS In whenaddition, the stressthe stress exponent exponent is close n ( ton = 2 1/ [21m].) Theis one stress of the exponent superplastic was about characteristic 1.89 in this parameter work for to the determine investigated the Al–Zn–Mg–Cudeformation mechanisms. alloy, which It stated was suggestedthat the superplastic that the superplastic deformation deformation mechanism mechanism might be GBSmight and be furtherGBS when microstructure the stress exponent validation is close to 2 [21]. The stress exponent was about 1.89 in this work for the investigated Al–Zn–Mg–Cu is required. alloy, which stated that the superplastic deformation mechanism might be GBS and further microstructure validation is required.

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3.5 2.5 (a) (b) 2.0 3.0 1.5 Q = 158.44 kJ/mol 2.5 1.0 0.33 1 = 188.13 kJ/mol

)]/ MPa 0.5 Q 2.0 0.38

MPa 1 /  0.0  0.50 Q = 177.78 kJ/mol ln 1.5 1 -0.5 ℃

470 [sinh( 0.62 -2 -1 485℃ ln 1.0 1 -1.0 Q = 250.54 kJ/mol 1×10 s 500℃ 3×10-3s-1 0.63 -1.5 515℃ -3 -1 0.5 1 1×10 s 530℃ -2.0 3×10-4s-1

0.0 -2.5 -8.5 -8.0 -7.5 -7.0 -6.5 -6.0 -5.5 -5.0 -4.5 1.24 1.26 1.28 1.30 1.32 1.34 1.36 -1 ln ε/s-1 1000/T / K . Figure 3. 3. ((aa)) TheThe relationship relationship between between ln σlnandσ lnandε (bln) theε calculation(b) the calculation of thermal of activation thermal activation energy. 3.2.2.energy Thermal. Activation Energy 3.2.2.The Thermal thermal Activation activation Energy energy contributes to further determine the superplastic deformation mechanism. It can be derived from the Arrhenius equation. The type of Arrhenius equation is able to The thermal activation energy contributes to further determine the superplastic deformation utilize both the large and small deformation: mechanism. It can be derived from the Arrhenius equation. The type of Arrhenius equation is able to utilize both the large and small deformation:.  Q  ε = A[sinh(ασ)] n exp (2) −RT n Q ε = A (sinh[ ασ)] − )(exp (2) where R represents the universal gas constant (8.314 J/(mol K)), TRTrepresents the absolute temperature · (K), A and α represent the material constants, n = 1/m represents the stress index and Q represents where R represents the universal gas constant (8.314 J/(mol·K)), T represents the absolute temperature the activation energy of superplastic deformation (J/mol). Therefore, take the natural logarithm for (K), A and α represent the material constants, n = 1/m represents the stress index and Q represents Equation (2): the activation energy of superplastic. deformation (J/mol). ThereforeQ , take the natural logarithm for ln ε = ln A + n ln[sinh(ασ)] . (3) Equation (2): − RT Since the strain rate is constant, the expression of deformation activationQ energy obtained by differentiating the Equation (3)ε can belnln derived+= nA as:ln[sinh(ασ)] − (3) RT .

Since the strain rate is constant, the expression∂ ln[sinh of(ασ deformation)] activation energy obtained by Q = Rn (4)  1  differentiating the Equation (3) can be derived as: ∂ . T ε

The activation energy Q, evaluated by the slope of ln[sinh(ασ)] 1000/T , is shown in Figure3b. − Under the different initial strain rates of 1 10ln2 s 1, 3(sinh[ ασ10)]3 s 1, 1 10 3 s 1 and 3 10 4 s 1, Q = Rn× − − × − − × − − × − (−4) the corresponding deformation activation energies Q were1 158 kJ/mol, 188 kJ/mol, 178 kJ/mol and ( ) 251 kJ/mol, respectively. These Q values were larger than 142 kJ/mol that is the lattice self-diffusion T ε activation energy of aluminum, which indicates that the superplastic deformation of the investigated ln (sinh[ ασ )] −1000 /T Al–Zn–Mg–CuThe activation alloy energy was dominated Q, evaluated by lattice by the di slopeffusion. of It was noticed that the, is initial shown microstructure in Figure 3b. ofUnder the material the different was not initial completely strain equiaxed rates of in 1 the× 10 initial−2 s−1, stage, 3 × 10 which−3 s−1, resulted 1 × 10−3 ins−1 theand large 3 × deformation 10−4 s−1, the activationcorresponding energy. deformation Thus, the dislocation activation slip energies is an important Q were 158 deformation kJ/mol, 188 mechanism kJ/mol, 178 in thekJ/mol initial and stage. 251 kJ/mol, respectively. These Q values were larger than 142 kJ/mol that is the lattice self-diffusion 3.3.activation Microstructural energy of Evolution aluminum, During which Superplastic indicates Deformationthat the superplastic deformation of the investigated Al–Zn–Mg–Cu alloy was dominated by lattice diffusion. It was noticed that the initial microstructure 3.3.1. Initial Microstructure of the material was not completely equiaxed in the initial stage, which resulted in the large deformationFigure4 illustratesactivation the energy. initial Thus, microstructure the dislocation and slip texture is an ofimportant the Al–Zn–Mg–Cu deformation alloymechanism after the in thermomechanicalthe initial stage. process [22]. The large and small grains coexist according to the phase map as shown in Figure4a at the traverse direction plane. The average grain size was roughly 10 µm, and the average3.3. Microstructur grain aspectal Evolution ratio of D studieduring Superplastic alloy was D 2.7eformation of the as-received sheet. Figure4b shows the distribution of misorientation angle, and the fractions of HAGBs and LAGBs were 75.6% and 24.3%, 3.3.1. Initial Microstructure Metals 2019, 9, x FOR PEER REVIEW 6 of 12

Figure 4 illustrates the initial microstructure and texture of the Al–Zn–Mg–Cu alloy after the thermomechanical process [22]. The large and small grains coexist according to the phase map as shown in Figure 4a at the traverse direction plane. The average grain size was roughly 10 μm, and Metals 2019, 9, 941 6 of 11 the average grain aspect ratio of studied alloy was 2.7 of the as-received sheet. Figure 4b shows the distribution of misorientation angle, and the fractions of HAGBs and LAGBs were 75.6% and 24.3%, respectively.respectively. The maximummaximum texturetexture intensity intensity was was 2.5 2.5 as as shown shown in in Figure Figure4c. 4 Figurec. Figure4d–f 4 representsd–f represents the themicrostructure microstructure of the of normal the normal direction direction (ND) plane,(ND) p rollinglane, rolling direction direction (RD) plane(RD) and plane traverse and directiontraverse direction(TD)(TD) plane, respectively. plane, respectively. It shows that It shows the initial that microstructure the initial microstructure consisted of mostlyconsisted recrystallized of mostly recrystallizedgrains together grains with together some banded with some grains, banded thus, thegrains, initial thus, microstructure the initial microstructure is not an ideal is equiaxednot an ideal for equiaxedsuperplastic for deformation.superplastic deformation.

Figure 4. Microstructure of the initial sheet: (a) Phase map, (b) grain–subgrain boundary map, (c) the maximum texture intensity, (d) the ND (normal direction) plane, (e) RD (rolling direction) plane and (f) TD (traverse direction) plane.

3.3.2. Evolution of Grain Morphology and Size during Superplastic Deformation The variations of grain morphology and size at different strains are shown in Figure5a,d. The variations of the aspect ratio and grain size with true strain are shown in Figure5e,f. The full recrystallization was achieved before the T1 stage as shown in Figure5a. It can be found that the coarse and banded grains were completely replaced by equiaxed and fine grains compared with the initial microstructure during superplastic deformation. As opposed to the initial microstructure, the aspect ratio decreased at the T1 stage and changed slightly from the T1 to T4 stage. The decrease of aspect ratio at the T1 stage was ascribed to the occurrence of DRX before the T1 stage, and the grain aspect ratio had small change from the T1 to T4 stage because of the role of GBS and grain rotation as shown in Figure5e. The average grain size increased during the superplastic deformation from T1 to T4 as shown in Figure5f. The straining could induce grain growth, at the same time, the high temperature also could promote the grain size to increase. The microstructure of the as-received sheet was not in a completely equiaxed state, but during superplastic deformation, it became a completely recrystallized microstructure. Once the equiaxed

grains are formed, this will be beneficial for GBS and grain rotation. The results show that the microstructure transformed from the partly equiaxed morphology to a uniformly and fully equiaxed Figure 4. Microstructure of the initial sheet: (a) Phase map, (b) grain–subgrain boundary map, (c) the structure,maximum which texture is similar intensity to,the (d) the results ND ( innormal other direction works [)23 plane,24]., ( Zhange) RD (rolling [23] found direction) the microstructure plane and transformation(f) TD (traverse from direction) the irregular planeα. ‘-martensite microstructure to equiaxed grains due to the DRX in the Ti-6Al-4V alloy. The equiaxed grains appeared in the interior of the specimen during the superplastic deformation of the Al–Zn–Mg–Cu alloy achieved by the friction stir process [24].

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Metals 2019, 9, 941 7 of 11 specimenspecimen during during the the superplastic superplastic deformation deformation of the Al–Zn––MgMg––CuCu alloyalloy achieved achieved by by thethe frictionfriction stir stir process process [ [2424].].

Figure 5. The evolution of grain morphology and size during superplastic deformation (a–d) respectively FigureFigure 5. 5. The The evolution evolution ofof graingrain morphologymorphology and size during superplastic deformationdeformation ((aa),), ( b(b),), ( c(c) ) for the true strain with 0.41, 0.69, 1.61 and 2.17. (e) The variation of grain aspect ratio with strains and andand ( d(d) )respectively respectively for for the the truetrue strainstrain withwith 0.41,0.41, 0.69, 1.61 and 2.17.. ((ee)) TThehe variationvariation of of grai grainn aspect aspect (f) the variation of grain size with strain. ratioratio with with strains strains and and ( (ff)) the the variationvariation ofof graingrain sizesize with strain.. The influence of superplastic deformation on angle of orientation as shown in Figure6 under TheThe influence influence ofof4 superplasticsuperplastic1 deformationdeformation on angle of orientation asas shownshown inin FigureFigure 6 6 under under 530 ◦C and 3 10− s− . The fraction of HAGBs at strains of 0.41, 0.69, 1.61 and 2.17 could be 530 °C and 3 ×× 10−4−4 s−1−1. The fraction of HAGBs at strains of 0.41, 0.69, 1.61 and 2.17 could be determined determined530 °C and 3 as× 10 98.5%, s . The 98.8%, fraction 98.9% of HAGBs and 98.5%. at strains The of HAGBs 0.41, 0.69, always 1.61 and kept 2.17 a large could fraction be determined during as 98.5%, 98.8%, 98.9% and 98.5%. The HAGBs always kept a large fraction during the process of theas 98. process5%, 98. of8%, superplastic 98.9% and deformation,98.5%. The HAGBs which always verified kept that a the large GBS fraction and grain during rotation the process were the of superplastic deformation, which verified that the GBS and grain rotation were the dominant dominantsuperplastic deformation deformation, mechanisms which verifie in thed superplasticthat the GBS tensile and grainprocess. rotation The maximum were the fraction dominant of deformation mechanisms in the superplastic tensile process. The maximum fraction of misorientation misorientationdeformation mechanisms achieved its in peakthe superplastic above between tensile 45 process.to 50 . The Meanwhile, maximum the fraction average of misorientationmisorientation achieved its peak above between 45° to 50°. Meanwhile,◦ the◦ average misorientation angle increased angleachieve increasedd its peak from above 37.4 betweento 39.6 45°from to 50°. T1 toMeanwhile, T3 and decreased the average to 36.3 misorientationat T4 stage. angle The increase increase ofd from 37.4° to 39.6° from◦ T1 to ◦ T3 and decreased to 36.3° at T4 stage.◦ The increase of average averagefrom 37. misorientation4° to 39.6° from angles T1 resultedto T3 and from decrease the graduallyd to 36.3° uniform at T4 and stage. equiaxed The increase grain structure. of average misorientationmisorientation angles angles resulted resulted fromfrom thethe gradugradually uniform and equiaxed graingrain structure.structure. 0.10 0.10 (a) 0.10 (a) 0.10 (b) (b) 0.08 0.08 0.08 0.08

0.06 0.06 0.06 0.06

0.04 Fraction 0.04 0.04 Fraction

Fraction 0.04 Fraction

0.02 0.02 0.02 0.02

0.00 0.00 0 10 20 30 40 50 60 0.00 0.00 10 20 30 40 50 60 0 10 20 30 40 50 60 Misorientation 10 20 Misorientation30 40 50 60 Misorientation Misorientation 0.10 0.10 0.10 (d) 0.10 (c) (d) (c) 0.08 0.08 0.08 0.08 0.06 0.06 0.06 0.06 0.04 0.04 Fraction Fraction 0.04

0.04 Fraction Fraction

0.02 0.02 0.02 0.02

0.00 0.00 10 20 30 40 50 60 10 20 30 40 50 60 0.00 0.00 10 20 Misorientation30 40 50 60 10 20 Misorientation30 40 50 60 Misorientation Misorientation Figure 6. The misorientation angle at different deformation from T1 to T4 (a), (b), (c) and (d) Figure 6. The misorientation angle at different deformation from T1 to T4 (a–d) respectively for the Figurerespectively 6. The for misorientation the true strains angle with 0.41, at different 0.69, 1.61 deformation and 2.17. from T1 to T4 (a), (b), (c) and (d) truerespectively strains with for the 0.41, true 0.69, strains 1.61 andwith 2.17. 0.41, 0.69, 1.61 and 2.17.

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3.3.3.3.3.3. The The Variation Variation Schmid SchmidFactor Factor andand InverseInverse PolePole FiguresFigures FigureFigure7 a–d7a–d shows shows the the variation variation tendency tendencyof of thethe SchmidSchmid factor at various experiment experiment conditions. conditions. TheThe average average Schmid Schmid factors factors at at deformations deformations of of 0.41, 0.41, 0.69, 0.69, 1.61 1.61 and and 2.17 2.17 were were 0.439, 0.439, 0.445, 0.445, 0.444 0.444 and and 0.45. The0.45. average The average value ofvalue Schmid of Schmid factor factor had little hadchange little change with thewith increase the increase of strain of strain and kept and akep larget a large value duringvalue during the superplastic the superplastic deformation deformation process. process. The large The Schmid large Schmid factor could factor be could beneficial be beneficial for GBS andfor grainGBS rotationand grain [25 rotation,26]. The [25 inverse,26]. The pole inverse figures pole at various figures experimentat various experiment conditions conditions is shown in is Figure shown7e–f. in TheFigure maximum 7e–f. The texture maximum intensity texture of all intensity the surfaces of all were the 1.48, surface 1.25,s were 1.39 and1.48, 1.551.25, at 1.39 diff erent and 1.55 strains at fromdifferent T1 to strains T4, which from wasT1 to less T4, than which the was intensity less than of 2.5the ofintensity the initial of 2.5 microstructure. of the initial microstructure. It is suggested thatIt is the suggested prominent that rolling the prominent deformation rolling texture deformation almost disappearedtexture almost and disappear the randomed and texture the random became dominant.texture became The grain dominant. orientation The grain in the or initialientation structure in the was initial roughly structure oriented was near roughly<111 oriented>. The texture near intensity<111>. The was texture weakened intensity significantly was weakened in all directions significantly especially in all directions in <111> directionespecially during in <111> superplastic direction deformation,during superplastic which indicates deformation, that the which GBS indicates and grain that rotation the GBS was andthe main grain deformation rotation was mechanism. the main Thedeformation results are mechanism. similar to the The works results of [27 are,28 similar]. The maximumto the works grain of orientation[27,28]. The was maximum reduced grainin the orientation was reduced in the AA8090 alloy under the tensile deformation at 530 °C at a 4constant1 AA8090 alloy under the tensile deformation at 530 ◦C at a constant strain rate of 5 10− s− [27]. −4 −1 × Mikhaylovskayastrain rate of 5 ×[ 2810] suggesteds [27]. Mikhaylovskaya that the formation [28] ofsuggested the weak that texture the formation after superplastic of the weak deformation texture ofafter the high-strengthsuperplastic deformation Al–Zn–Mg–Cu of the alloy high is-strength because Al of– theZn– mechanismMg–Cu alloy of is grain-boundary because of the mecha sliding.nism of grain-boundary sliding.

Figure 7. (a–d) Schmid factor map of the Al–Zn–Mg–Cu alloy deformed at different true strains with 0.41,Figure 0.69, 7.1.61 (a–d and) Schmid 2.17. (factore–h) Inverse map of pole the Al figure–Zn– andMg– textureCu alloy intensity deformed during at different superplastic true strains deformation with with0.41, strains 0.69, of 1.61 0.41, and 0.69, 2.17 1.61. (e and– h) 2.17. Inverse pole figure and texture intensity during superplastic deformation with strains of 0.41, 0.69, 1.61 and 2.17.

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Metals 2019, 9, x FOR PEER REVIEW 9 of 11 It can be seen that the texture intensity abnormally increased at the T4 stage, which may be becauseIt can the be local seen stress that concentration the texture intensity was generated. abnormally When increase the stressd at concentration the T4 stage, could which not may be fully be becauserelaxed inthe time local at stress large deformation,concentration then was thegenerated. cavities When were prone the stress to appear concentration in the material could asnot shown be fully in relaxedFigure8 .in time at large deformation, then the cavities were prone to appear in the material as shown in Figure 8.

Figure 8. TheThe cavities occur during the superplastic deformation.deformation. ( (aa)) 3D 3D-electron-electron back scattering diffractiondiffraction ( (EBSD)EBSD) microstructur microstructuree and (b)) metallographic microstructure. In summary, the maximum elongation of 1020% was obtained in the present work under the In summary, the maximum elongation of 1020% was obtained in the present work under the optimum condition. In our study, the alloy was not fully recrystallized microstructure in the initial state, optimum condition. In our study, the alloy was not fully recrystallized microstructure in the initial but transformed into the fully recrystallized microstructure by DRX. In Chen’s work [29,30], the uniform state, but transformed into the fully recrystallized microstructure by DRX. In Chen’s work [29,30], and equiaxed microstructure with a grain size of 10 µm of alloy was prepared for superplastic testing the uniform and equiaxed microstructure with a grain size of 10 μm of alloy was prepared for and the elongation to failure of the material under a strain rate of 10 3s 1 and temperature of 516 C superplastic testing and the elongation to failure of the material under− − a strain rate of 10−3s−1 and◦ was 850%. The elongation of the sample in superplastic testing in the present study is higher (1020%) temperature of 516 °C was 850%. The elongation of the sample in superplastic testing in the present in the similar deformation condition. It should be mentioned that the alloy in our work was not a fully study is higher (1020%) in the similar deformation condition. It should be mentioned that the alloy recrystallized microstructure in the initial state, but transformed into a fully recrystallized one by DRX in our work was not a fully recrystallized microstructure in the initial state, but transformed into a during superplastic deformation. Therefore, it is suggested that good superplasticity could be achieved fully recrystallized one by DRX during superplastic deformation. Therefore, it is suggested that good in the non-ideal initial microstructure of the Al–Zn–Mg–Cu alloy and the elongation might be even superplasticity could be achieved in the non-ideal initial microstructure of the Al–Zn–Mg–Cu alloy higher compared with the homogenous microstructure of the material, as the equiaxed grains could be and the elongation might be even higher compared with the homogenous microstructure of the obtained and grain refining could be realized during superplastic deformation. This means that the material, as the equiaxed grains could be obtained and grain refining could be realized during preparation processes for the Al–Zn–Mg–Cu superplastic alloy could be simplified, which contributes superplastic deformation. This means that the preparation processes for the Al–Zn–Mg–Cu to the preparation of the material conveniently for superplastic forming and to the improvement of superplastic alloy could be simplified, which contributes to the preparation of the material production efficiency. On the other hand, though, low temperature and high speed superplasticity can conveniently for superplastic forming and to the improvement of production efficiency. On the other be obtained from the ultrafine Al 7075 alloy prepared by severe plastic deformation such as friction stir hand, though, low temperature and high speed superplasticity can be obtained from the ultrafine Al processing [24], the method is only suitable for a small batch production. Thus, the preparation method 7075 alloy prepared by severe plastic deformation such as friction stir processing [24], the method is in our study can be applied for large parts of superplastic forming, which has greater industrial value. only suitable for a small batch production. Thus, the preparation method in our study can be applied for4.Conclusions large parts of superplastic forming, which has greater industrial value.

4. ConclusionIn this study, a systematic investigation on the effect of strain on microstructure evolution of the Al–Zn–Mg–Cu alloy was performed. The superplastic deformation behaviors were investigated In this study, a systematic investigation on the effect of strain on microstructure evolution of the under different temperatures of 470 ◦C, 485 ◦C, 500 ◦C, 515 ◦C and 530 ◦C, and different strain rates of Al–Zn–4Mg–1Cu alloy3 was1 performed2 1. The superplastic2 1 deformation behaviors were investigated 3 10− s− , 1 10− s− , 3 10− s− and 1 10− s− and the results were summarized as follows: under× different× temperatures× of 470 °C, 485 °C,× 500 °C, 515 °C and 530 °C , and different strain rates (1) The maximum elongation of 1020% in the Al–Zn–Mg–Cu alloy was attained under the condition −4 −1 −3 −1 −2 −1 −2 −1 of 3 × 10 s , 1 × 10 s4 , 31 × 10 s and 1 × 10 s and the results were summarized as follows: of 530 ◦C and 3 10− s− , showing the excellent superplastic deformation capability of the material. (1) The maximum× elongation of 1020% in the Al–Zn–Mg–Cu alloy was attained under the The peak stress decreased with deformation temperature increasing and with strain rate decreasing. condition of 530 °C and 3 × 10−4 s−1, showing the excellent superplastic deformation capability of the The elongation at fracture increased with the increase of the deformation temperature and with the material. The peak stress decreased with deformation temperature increasing and with strain rate decrease of strain rate. decreasing. The elongation at fracture increased with the increase of the deformation temperature and with the decrease of strain rate.

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(2) The activation energy and strain rate index values were 250.5 kJ/mol and 0.63 in the optimum loading conditions, respectively, which indicates that the superplastic deformation mechanism was GBS accommodated by lattice diffusion in the initial stage. (3) The Al–Zn–Mg–Cu alloy had some banded grains in the initial sample, but the homogeneous and equiaxed structure was obtained by DRX in the initial stage of superplastic deformation. The Schmid factor, the grain aspect ratio and the fraction of HAGBS changed slightly during the superplastic process, revealing that the GBS and grain rotation were the main deformation mechanisms. Meanwhile, the grains grew obviously during superplastic deformation once the recrystallization process was finished. (4) The texture intensity increased in the final stage of superplastic deformation, indicating that local stress concentration was generated, which could result in cavity nucleation.

Author Contributions: Conceptualization, H.D. and H.H.; methodology, N.Z.; software, G.L. and J.W.; validation, H.D., G.L. and H.H.; formal analysis, G.L.; investigation, J.W.; resources, H.D.; data curation, G.L.; writing—original draft preparation, G.L.; writing—review and editing, G.L. and H.D.; visualization, H.D.; supervision, H.H.; project administration, H.D.; funding acquisition, H.H. Funding: This research was funded by National Science Foundation of China, grant number 51334006. Conflicts of Interest: The authors declare no conflict of interest.

References

1. Zhang, Y.; Jin, S.; Trimby, P.; Liao, X.; Murashkin, M.Y.; Valiev, R.Z.; Sha, G. Strengthening mechanisms in an ultrafine-grained AlZnMgCu alloy processed by high pressure torsion at different temperatures. Mater. Sci. Eng. A 2019, 752, 223–232. [CrossRef] 2. Chen, Y.; Ding, H.; Cai, Z.; Zhao, J.; Li, J. Effect of initial base metal temper on microstructure and mechanical properties of friction stir processed Al-7B04 alloy. Mater. Sci. Eng. A 2016, 650, 396–403. [CrossRef] 3. Marlaud, T.; Deschamps, A.; Bley, F.; Lefebvre, W.; Baroux, B. Influence of alloy composition and heat treatment on precipitate composition in Al–Zn–Mg–Cu alloys. Acta Mater. 2010, 58, 248–260. [CrossRef] 4. Liu, Z.; Li, P.; Xiong, L.; Liu, T.; He, L. High-temperature tensile deformation behavior and microstructure evolution of Ti55 titanium alloy. Mater. Sci. Eng. A 2017, 680, 259–269. [CrossRef] 5. Wang, Y.N.; Huang, J.C. Comparison of grain boundary sliding in fine grained Mg and Al alloys during superplastic deformation. Scripta Mater. 2003, 48, 1117–1122. [CrossRef] 6. McQueen, H.J. Development of dynamic recrystallization theory. Mater. Sci. Eng. A 2004, 387–389, 203–208. [CrossRef] 7. Zhang, X.; Cao, L.; Zhao, Y.; Chen, Y.; Tian, X.; Deng, J. Superplastic behavior and deformation mechanism of Ti600 alloy. Mater. Sci. Eng. A 2013, 560, 700–704. [CrossRef] 8. Wei, X.; Lee, D.; Shim, S.; Chen, X.; Kysar, J.W. Plane-strain bulge test for nanocrystalline copper thin films. Scripta Mater. 2007, 57, 541–544. [CrossRef] 9. Mukhopadhyay, A.K.; Kumar, A.; Raveendra, S.; Samajdar, I. Development of grain structure during superplastic deformation of an Al–Zn–Mg–Cu–Zr alloy containing Sc. Scripta Mater. 2011, 64, 386–389. [CrossRef] 10. Kumar, A.; Mukhopadhyay, A.K.; Prasad, K.S. Superplastic behaviour of Al–Zn–Mg–Cu–Zr alloy AA7010 containing Sc. Mater. Sci. Eng. A 2010, 527, 854–857. [CrossRef] 11. Li, S.; Jin, S.; Huang, Z. Cavity Behavior of Fine-Grained 5A70 Aluminum Alloy during Superplastic Formation. Metals 2018, 8, 1065. [CrossRef] 12. Li, S.; Huang, Z.; Jin, S. Superplastic Behavioral Characteristics of Fine-Grained 5A70 Aluminum Alloy. Metals 2019, 9, 62. [CrossRef] 13. Duan, Y.; Xu, G.; Zhou, L.; Xiao, D.; Deng, Y.; Yin, Z.; Peng, B.; Pan, Q.; Wang, Y.; Lu, L. Achieving high superplasticity of a traditional thermal–mechanical processed non-superplastic Al–Zn–Mg alloy sheet by low Sc additions. J. Alloys Compd. 2015, 638, 364–373. [CrossRef] 14. Sun, Z.C.; Wu, H.L.; Cao, J.; Yin, Z.K. Modeling of continuous dynamic recrystallization of Al-Zn-Cu-Mg alloy during hot deformation based on the internal-state-variable (ISV) method. Int. J. Plast. 2018, 106, 73–87. [CrossRef] Metals 2019, 9, 941 11 of 11

15. Mosleh, A.O.; Mikhaylovskaya, A.V.; Kotov, A.D.; AbuShanab, W.S.; Moustafa, E.B.; Portnoy, V.K. Experimental Investigation of the Effect of Temperature and Strain Rate on the Superplastic Deformation Behavior of Ti-Based Alloys in the (α + β) Temperature Field. Metals 2018, 8, 819. [CrossRef] 16. Qin, G.; Chen, R.; Liaw, P.K.; Gao, Y.; Li, X.; Zheng, H.; Wang, L.; Su, Y.; Guo, J.; Fu, H. A novel face-centered-cubic high-entropy alloy strengthened by nanoscale precipitates. Scripta Mater. 2019, 172, 51–55. [CrossRef] 17. Mikhaylovskaya, A.V.; Kotov, A.D.; Kishchik, M.S.; Prosviryakov, A.S.; Portnoy, V.K. The Effect of Isothermal Multi-Directional Forging on the Grain Structure, Superplasticity, and Mechanical Properties of the Conventional Al–Mg-Based Alloy. Metals 2019, 9, 33. [CrossRef] 18. Masuda, H.; Tobe, H.; Hara, T.; Sato, E. Three-dimensional characterization of superplastic grain boundary sliding inside Al–Zn–Mg–Cu alloy sheet. Scripta Mater. 2019, 164, 82–85. [CrossRef] 19. Roy, S.; Suwas, S. Enhanced superplasticity for (α + β)-hot rolled Ti–6Al–4V–0.1B alloy by means of dynamic globularization. Mater. Des. 2014, 58, 52–64. [CrossRef] 20. Backofen, W.A. Superplasticity in an Al-Zn Alloy. Trans. ASM. 1964, 57, 980–990. 21. Chen, B.; Tian, X.L.; Li, X.L.; Lu, C. Hot Deformation Behavior and Processing Maps of 2099 Al-Li Alloy. J. Mater. Eng. Perform. 2014, 23, 1929–1935. [CrossRef] 22. Wert, J.A.; Paton, N.E.; Hamilton, C.H.; Mahoney, M.W. Grain refinement in 7075 aluminum by thermomechanical processing. Metall. Trans. A 1981, 12, 1267–1276. [CrossRef] 23. Zhang, W.; Ding, H.; Cai, M.; Yang, W.; Li, J. Ultra-grain refinement and enhanced low-temperature superplasticity in a friction stir-processed Ti-6Al-4V alloy. Mater. Sci. Eng. A 2018, 727, 90–96. [CrossRef] 24. Orozco-Caballero, A.; Álvarez-Leal, M.; Hidalgo-Manrique, P.; Cepeda-Jiménez, C.M.; Ruano, O.A.; Carreño, F. Grain size versus microstructural stability in the high strain rate superplastic response of a severely friction stir processed Al–Zn–Mg–Cu alloy. Mater. Sci. Eng. A 2017, 680, 329–337. [CrossRef] 25. Hémery, S.; Tromas, C.; Villechaise, P. Slip-stimulated grain boundary sliding in Ti-6Al-4 V at room temperature. Materialia 2019, 5, 100189. [CrossRef] 26. Watanabe, H.; Fukusumi, M.; Somekawa, H.; Mukai, T. Texture and mechanical properties of a superplastically deformed Mg–Al–Zn alloy sheet. Scripta Mater. 2009, 61, 883–886. [CrossRef] 27. Bate, P.; Ridley, N.; Zhang, B. Mechanical behaviour and microstructural evolution in superplastic Al–Li–Mg–Cu–Zr AA8090. Acta Mater. 2007, 55, 4995–5006. [CrossRef] 28. Mikhaylovskaya, A.V.; Kotov, A.D.; Pozdniakov, A.V.; Portnoy, V.K. A high-strength aluminium-based alloy with advanced superplasticity. J. Alloys Compd. 2014, 599, 139–144. [CrossRef] 29. Chen, C.L.; Tan, M.J. Effect of grain boundary character distribution (GBCD) on the cavitation behaviour during superplastic deformation of Al 7475. Mater. Sci. Eng. A 2002, 338, 243–252. [CrossRef] 30. Chen, C.L.; Tan, M.J. Cavity growth and filament formation of superplastically deformed Al 7475 Alloy. Mater. Sci. Eng. A 2001, 298, 235–244. [CrossRef]

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