稀有金属(英文版) 2019,38(12),1136-1143

Processing maps and microstructural evolution of Al-Cu-Li alloy during hot deformation

Sheng-Li Yang Jian Shen Yong-An Zhang Zhi-Hui Li Xi-Wu Li Shu-Hui Huang Bai-Qing Xiong

State Key Laboratory of Nonferrous Metals and Processes General Research Institute for Nonferrous Metals

作者简介：*Jian Shen e-mail:shenjgrinm@126.com;

收稿日期：17 January 2016

基金：financially supported by the National Program on Key Basic Research Project of China (No.2012CB619504);the National Natural Science Foundation of China (No.51274046);

Processing maps and microstructural evolution of Al-Cu-Li alloy during hot deformation

Sheng-Li Yang Jian Shen Yong-An Zhang Zhi-Hui Li Xi-Wu Li Shu-Hui Huang Bai-Qing Xiong

State Key Laboratory of Nonferrous Metals and Processes General Research Institute for Nonferrous Metals

Abstract：

The hot deformation behavior of Al-Cu-Li alloy was investigated by hot compression tests in the temperature range of 340-500℃ with strain rate of 0.001-10.000 s^-1.Based on the dynamic materials model(DMM),processing maps of the test alloy were developed for optimizing hot processing parameters.The optimum parameters of hot deformation for Al-Cu-Li alloy are at temperature of 400-430℃and strain rate of about 0.100 s^-1,with efficiency of power dissipation of around 30%.The microstructural manifestation of the alloy deformed in instability domains is flow localization,and dynamic softening first occurs in flow localizations structure.In stable domains,dynamic recovery(DRV) and dynamic recrystallization(DRX) are the main microstructural evolution mechanism.DRX is gradually strengthened with the increase in deformation temperature and the decrease in strain rate.During hot deformation,the DRX mechanism of Al-Cu-Li alloy is dominated by continuous DRX(CDRX).A DRX model of Al-Cu-Li alloy is proposed based on the microstructural evolution process of the test alloy.

Keyword：

Al-Cu-Li alloy; Processing map; Dynamic recovery and dynamic recrystallization; Microstructural evolution;

Received： 17 January 2016

1 Introduction

Al-Cu-Li alloys are being extensively used for aerospace structures due to their high stiffness to density ratio and large elastic modulus compared with conventional 2xxx and 7xxx series alloys [ 1] .These extraordinary properties result from the addition of Li [ 2] .For every 1 wt%addition of Li to Al,the density reduced by 3%and the elastic modulus increased by almost 6%,up to the maximum solid solubility of Li in Al,i.e.,4.2 wt%Li [ 3] .AlCu-Li alloy is the third generation of Al-Li alloy,which reduced the content of lithium and added a variety of microalloying elements (Mg,Mn,Zr and Zn,etc.) compared to the previous two generations of Al-Li alloy.Key compositional design principles result in the successful application and commercialization of the Al-Cu-Li alloy [ 4] .Hot deformation process and heat treatment technology are the keys to controlling Al-Cu-Li alloy mechanical properties.

In recent years,the processing map based on dynamic materials model (DMM) is considered to be an important mode]for optimizing the hot working of metals or alloys [ 5, 6, 7] .Stable domains and instability domains can be predicted by using processing maps,constructed from the flow stress as a function of temperature,strain rate and strain.Processing maps have been used to identify the deformation temperature and strain rate corresponding to different microstructural evolution of stable domains and instability domains [ 5, 7, 8, 9] .Generally,dynamic recovery(DRV) and dynamic recrystallization (DRX) are the main micro structure mechanism of the alloy deformed in stable domains.Nayan et al. [ 10] studied the microstructural evolution of AA219 5 Al-Cu-Li alloy during large strain deformation.They proved that DRX grain formed due to the conversion of low angle boundaries (LAB,<15°) to high angle boundaries (HAB,>15°).Xun and Tan [ 11] studied the dynamic and static recrystallization of 8090 Al-Li alloy using electron backscattcr diffraction (EBSD).They concluded that DRV and DRX are ascribed to the rotation of subgrains.Sakai et al. [ 12] reported the continuous DRX (CDRX) of 7475aluminum alloy at different strain rates.It is proposed that DRX grains formed as a result of strain-induced continuous reactions at low strain rate (3×10^-4 s^-1) or grain boundary sliding at high strain rate (3×10^-2s^-1).Sitdikov et al. [ 13] reported DRX mechanism in a coarsegrained 7475 aluminum alloy during hot deformation.The results indicated that DRX grains originated from the microshear bands.In addition,most of the research results show that the microstructural manifestations of the alloy deformed in instability domains are linked to adiabatic shear bands,flow localization and dynamic strain aging [ 14, 15] .And the process of microstructural evolution in instable organization is closely related to DRV and DRX.Li et al. [ 16] studied adiabatic shear band of 7075 aluminum alloy.They proposed that the elongated cells form

subgrains with their misorientation increasing and eventually convert to a DRX structure prior to the formation of adiabatic shear bands.From above discussion,it can be known that DRV and DRX play an important role in the microstructural evolution process of stable domains and instability domains in perse mletallic materials.Generally,DRX mechanism of aluminum alloy mainly has two kinds of discontinuous DRX (DDRX) and CDRX [ 17, 18] .DDRX included nucleation and growth process of grain.But,under certain conditions,the high angle grain boundaries may evolve during high-temperature deformation in ways other than the nucleation and growth of grains at pre-existing boundaries,which is classified as CDRX.However,DRV and DRX mechanism of Al-CuLi alloy with high stacking fault energy is not yet well understood.More work is required at this area to further understand the microstructural evolution process during hot deform ation.

In this work,the effect of strain on processing map of an Al-Cu-Li alloy was investigated based on isothermal compression tests.Optimized hot deformation parameters were established by processing map.Microstructural evolution mechanism of Al-Cu-Li alloy deformed in stable domain and instability domain was discussed.A DRX model was proposed based on the microstructural evolution processing during hot deformation by EBSD and transmission electron microscope (TEM)analysis.The optimal hot processing parameters will provide the important guideline for the optimization of deformation techniques and the improvement of microstructure.

2 Experimental

The chemical composition of the test alloy in this work was2.50Cu,1.58Li,0.30Mn,0.12Zr,0.06 Mg,0.01 Zn,0.05Ti and balance Al (wt%).The cast ingot was homogenized at460℃for 20 h followed by 525℃for 24 h and then cooled down to room temperature in water.Cylindrical specimens with 10 mm in diameter and 15 mm in height were machined for compression tests.The isothermal compression tests were conducted on a Gleeble-1500thermosimulation machine in the temperature range of340-500℃at intervals of 40℃and under five different strain rates (0.001,0.010.0.100,1.000 and 10.000 s^-1).Graphite sheet was used between the specimens for reducing the deformed friction.The specimens were heated with a rate 5℃·s^-1 and held for 3 min at the deformation temperatures in order to establish a uniform temperature prior to deformation.The height reduction of the specimens was 60%by the end of the compression tests.The speeimens were cooled down by water immediately after deformation,so as to retain the microstructure pattern at elevated temperature.The specimens were axially sectioned and mechanically polished.EBSD (JEOL,JSM-7001F scanning electron microscope) investigation was carried out after electrochemically polished using a solution containing 10 ml HClO₄ and 90 ml methanol.TEM films were prepared by the conventional method:0.5-mmthick foils were cut from the deformed specimens.Then,the thick foils were ground into 0.05 mm and disks with3 mm in diameter were punched out from these foils and subsequently two-jet thinned in a solution of 75 m1 HNO₃and 225 mll methanol-cooled to-30℃.TEM observations were conducted on Tecnai G20.

According to the principles of the DMM [ 19] ,processing map can be constructed based on the true stress-strain curve,which was utilized for exploring hot workability of materials.In this model,the mechanical processing and work piece can be viewed as a closed system and power dissipation,respectively.The power dissipation is composed of viscoplastic heat produced by plastic deformation and the power required to cause microstructural changes,such as DRV,DRX and phase transformations.At the given deformation temperature,strain and strain rate,the power dissipation through microstructural changes can be represented by the efficiency of power dissipation (η),given by:

where m is strain rate sensitivity,calculated from m= (lnσ)/?(lnε),whereσis the flow stress and is the strain rate.The power dissipation map is obtained by plotting isoefficiency contour map ofηat different deformation temperatures and strain rates.

Based on the extremum principles of irreversible thermodynamics,an instability criterion was proposed to determine the onset of flow instability,which was derived by [ 5] :

The instability map can be plotted by different values ofξ( ) at different deformation conditions,and the flow instability is predicted whenξ( ) becomes negative.The variation of instability criteriaξ( ) as a function of temperature and strain rates constitutes the instability map.Processing map can be developed by superimposing the instability map over the power dissipation map.Different domains of the processing map may be controlled by different microstructural mechanisms.

3 Results and discussion

3.1 Processing map

It has been demonstrated that the processing map is very useful for optimizing hot workability and controlling microstructure in the material.Figure 1 shows the effects of strain on the processing maps of Al-Cu-Li alloy gencrated in the temperature range of 340-500℃,strain rate range of 0.001-10.000 s^-1 and the strains of 0.3,0.5,0.7and 0.9.In the processing maps,the contours represent the efficiency of power dissipation and the number of the contour indicates the dimension of the efficiency of power dissipation (η).Shadow domains represented instable domains (ξ( )<0).

From Fig.1,it can be easily found that instable domain increases with the increase in strain.As shown in Fig.la,the processing map exhibits two instable domains at a strain of 0.3.The instable Domain I occurs at temperature of lower than 380℃and strain rates from 0.002 to 0.715.Instable DomainⅡlocates at temperature of 380-440℃and high strain rates.The effect of strain on the instable Domain I is slight with strain increasing from 0.3 to0.5.However,the instable DomainⅡmigrates to high temperature with the increase in strain (Fig.1b).As shown in Fig.1c,instability DomainⅡemerges at low temperature and high strain rate.As a result,the instable DomainsⅠandⅡmerge into a big zone at the strain of 0.7.The big instable domain migrates to high temperature and a wider range of strain rate at the strain of 0.9 (Fig.1d).

According to Fig.1,it can be observed that there are some similarities among different strains:The efficiency of power dissipation increases with strain rates decreasing.The minimum and maximum values of efficiency of power dissipation are obtained at low temperature with high strain rate (at temperature of 340℃with strain rate of10.000 s^-1) and at high temperature with low strain rate (at temperature of 500℃with strain rate of 0.001 s^-1),respectively.From Fig.1d.it can be found that there are four peak efficiency domains at strain of 0.9.The first domain is the temperature range of 340-370℃and strain rate range of 0.001-0.002 s^-1.The maximum efficiency of power dissipation in this domain is 32%.But.this domain is inappropriate for the alloy hot processing because the domain is so narrow and close to the instable domain.The second domain is the temperature range of 400-440℃and strain rate range of 0.001-0.056 s^-1 corresponding to a peak efficiency of power dissipation of 32%.The third domain is the temperature range of 440-470℃and strain rate range of 0.001-0.562 s^-1 with a peak efficiency of power dissipation of 30%.The fourth domain is the temperature range of 470-500℃and strain rate range of0.001-0.316 s^-1.The efficiency of power dissipation of this domain increases from 26%to 36%.The variation of the peak efficiency of power dissipation indicates that the hot workability of Al-Cu-Li alloy will be improved with the increase in temperature and decrease in strain rate.Generally,the better workability of deformable material,the higher efficiency of powwer dissipation.Nevertheless,the highest efficiency of power dissipation does not mean better workability [ 20] .Some microstructural evolution of instability structure,such as adiabatic shear bands,flow localizations and dynamic strain aging,can also lead to a high efficiency of power dissipation [ 13] .Only the deformation temperature and strain rate corresponding to the peak efficiency of power dissipation in the stable domains can be considered as the optimum hot working parameters.Therefore,optimal hot processing range from Fig.1d is the temperature range of 400-500℃and strain rate range of0.001-0.100 s^-1.However,the high temperature and low strain rate during hot processing will bring high energy consumption and low processing efficiency.Therefore,it is reasonable to choose the temperature range of 400-430℃and strain rate of about 0.100 s^-1 for practical manufacturing.

3.2 Microstructural evolution

In order to investigate microstructural evolution mechanism and verify the reliability of the hot deformation parameters predicted by processing map,the microstructures of the test alloy deformed under the specific process parameters in instable domains and stable domains were analyzed.

Figure 2 shows EBSD maps at different deformation conditions of instable domains in Fig.1d.Figure 2a,b/c.d/e,f/g,h corresponds to A/B/C/D letters labeled in Fig.1d.Figure 2 a,c,e,g is the inverse pole figure (IPF)maps,and Fig.2 b,d,f,h is the image quality (IQ) maps.

Fig.1 Processing maps of Al-Cu-Li aluminum alloy under different strains:a 0.3,b 0.5,c 0.7 and d 0.9

Fig.2 EBSD maps (a,c,e,g IPF and b,d,f,h IQ maps) of test alloy deformed at different deformation conditions in instable domains:a,b 340℃10.000 s^-1;c,d 340℃,0.100 s^-1;e,f 380℃,0.100 s^-1;g,h 420℃,1.000 s^-1

Microstruetural evolution in the instable domains is corresponding to flow localizations,dynamic strain aging,mechanical twinning or kinking and flow rotation [ 13, 15] ,and it should be avoided during hot deformation.It can be seen from the IPF of Fig.2 that flow localizations are mainly deformation mechanism in instable domains.The generation of flow localizations could be related to the following reasons:A large amount of plastic work produced at high speed is transformed into deformation heat in a short time when the alloy deformed at low temperature and high strain rate,resulting in the rapid increase in local temperature of the alloy.The local temperature of the aluminum alloy can be increased [ 21] ,inducing flow localizations along the maximum shear stress direction,as shown in the IPF maps in Fig.2 (inside the black dotted line).As a result,it is triggered that the dynamic softening first occurs in the instable structure,inducing DRV and DRX non-uniform distribution in the deformed microstructure and making the workability of the alloy worse in turn.Besides,the area of instability structure is gradually reduced with the increase in temperature and the decrease in strain rate.As shown in IQ maps in Fig.2(inside the white dotted line),there are a large number of fine subgrains in flow localizations structure,which proves that DRV is the main microstructural evolution mechanism in flow localizations structure.The increase in the size of subgrains (Fig.2f,h) demonstrates that the higher deformation temperature and lower strain rate are beneficial to promote the dynamic softening in the flow localizations structure.

Figure 3 shows EBSD maps at different deformation conditions of stable domains in Fig.1d.Figure 3a,b/c,d/e,f/g,h corresponds to E/F/G/H letters labeled in Fig.1d.Figure 3 a,c,e,g is IPF maps,and Fig.3b,d,f,h is IQ maps.

Figure 3 shows that the morphology of the deformed grains in stable domains is more homogeneous than those of the instable domain (Fig.2).The efficiency of power dissipation is conforming to microstructural evolution mechanism.Reddy et al. [ 22] concluded that the peak efficiencies of Al-Li alloys for DRV and DRX were about23%and 42%,respectively.As shown in Fig.3a,b(420℃/0.100 s^-1),corresponding to the efficiency of power dissipation of 26%,the shape of the grains still keeps the elongated grains in the deformed micros truetures.Subgrains and partial DRX grains exist almost exclusively at grain boundaries,indicating that DRV and DRX occur simultaneously and DRV is the principal microstructural evolution mechanism.Figure 3c,d (420℃/0.001 s^-1) and e,f (500℃/0.100 s^-1) corresponds to the efficiencies of power dissipation of 32%and 30%,respectively.When the test alloy was deformed at lower strain rate and higher deformation temperature,the characteristic of DRX is more obvious.It can be observed that the size of subgrains and DRX grains increases,and a large number of subgrains and DRX grains occur in the deformed grain interior.DRX becomes the dominant deformation mechanism at high temperatures and low strain rates as shown in Fig.3 g h,corresponding to the efficiency of power dissipation of more than 36%.The elongated grains in the deformed micro structures cannot be found.At higher deformation temperature and lower strain rate,the diffusion abilities of atoms,cross-slip of dislocation and migration of grain boundaries are enhanced,and the available time is extended.All these conditions are beneficial to the nucleation and growth of the DRX.Therefore,both the number and size of the DRX grains increase obviously.

3.3 DRX model of Al-Cu-Li alloy

During hot deformation,DRV and DRX of Al-Cu-Li alloy are closely related to the distribution of original grain boundaries and second-phase particles [ 8] ,which is attributed substantially to the dislocation configuration.Grain boundary distribution map and frequency of misorientation at different height reductions in isothermal deformation (500℃/0.001 s^-1) of the test alloy are shown in Fig.4,and the corresponding TEM images are presented in Fig.5.

Fig.3 EBSD maps (a,c,e,g IPF and b,d,f,h IQ maps) of test alloy deformed at different deformation conditions in instable domains:a,b 420℃,0.100 s^-1;c,d 420℃,0.001 s^-1;e,f 500℃,0.100 s^-1;g,h 500℃,0.001 s^-1

As seen from Fig.4a,there are no obvious DRX grains when the strain is 0.105.The grain boundaries are smooth.The frequency of LAB is only 19.6%.The frequency of LAB sharply increases to 66.5%when the strain increases to 0.405.LAB subgrain structure occurs preferentially adjacent to the original high angle grain boundaries,as shown in Fig.4b.A few of the original grain boundaries are bulged,as indicated by the red arrow in Fig.4b.Some original grain boundary contacts with subgrain boundary toform some trigeminal node,as indicated by the white dashed circle in Fig.4b.The frequency of LAB further increases to 73.6%when the strain increases to 0.693,as shown in Fig.4c.Some recrystallized grains with small size form in the trigeminal node,as indicated by the white dashed circle in Fig.4c.As the strain continues increasing,more and more subgrains gradually could be investigated in the internal deformation of grain with the strain increasing.A small amount of fine DRX grains form in the vicinity of the original grain boundaries and some coarse second-phase particles.Meanwhile,the frequency of LAB decreases from 73.6%to 63.1%when the strain increases from 0.693 to 1.003,as shown in Fig.4d.It is worth noting that the grain gradually changes to lamellar microstructure with the increase in strain,as shown in Fig.4d,e.The lamellar microstructure boundaries are aligned parallel to the compression plane,together with intersecting boundaries which are mainly of LAB.And it tends to reduce the energy by localized boundary migration and finally form DRX grains,as indicated by white dashed line in Fig.4d,e.The frequency of LAB decreases to 55.1%when the strain increases to 1.332.

Fig.4 Effects of strain on grain boundary distribution and frequency of misorientation of test alloy (deformation temperature of 500℃,strain rate of 0.001 s-1):aε=0.105,bε=0.405,cε=0.693,dε=1.003 and eε=1.332 (red lines:2°-5°,green lines:5°-15°,blue lines:above 15°misorientation)

At the initial stage of deformation,a large number of dislocations tangled with each other are induced.Out-oforder dislocation gradually transforms into dislocation nets by polygonizing (annihilation and rearrangement of dislocations),then forming subgrain structure,as shown in Fig.5a,b,causing the increase in LAB frequency (Fig.4a,b,c).The deformed grains gradually turn into lamellar structure with the increase in strain.It is thought that the energy is lowered by localized boundary migration.As shown in Fig.5c,the lamellar microstructure tends to collapse due to the surface tension at the trigeminal node [ 17] .After that,the grains begin to spheroidize and grow by Y-junction migration (Fig.5d),leading to the decrease of LAB frequency (Fig.4d,e).

According to the microstructural evolution process of the test alloy during hot deformation (Figs.4,5),a DRX model of Al-Cu-Li alloy is proposed to describe the microstructural evolution of the test alloy,which is based on the model proposed by Humphreys and Hatherly [ 17] ,as shown in Fig.6.The model could be described as follows.

1.Several initial grains are depicted in Fig.6a.Duringthe deformation process,a large number of straininduced dislocation are rearranged to form subgrain in the initial grain,as shown in Fig.6b.At the same time,the initial grains gradually transform into lamellar structure,as shown in Fig.6c.

2.The boundary of lamellar structure collapses due tothe surface tension at the node point such as A in Fig.6c,where the boundaries perpendicular to the compression direction are pulled by the boundaries non-perpendicular to the compression direction.The critical condition for collapse of the boundaries is when A₁ and A₂ nodes touch.This depends on the length of different direction boundaries and its relative boundary energies (Reference [ 17] gives the related calculation principle.).

Fig.5 TEM images of Al-Cu-Li alloy deformed at different strains (deformation temperature of 500℃,strain rate of 0.001 s^-1):aε=0.405.bε=0.693,cε=1.003 and dε=1.332

Fig.6 DRX model graphs proposed for describing microstructural evolution during deformation:a initial grain structure,b moderate deformation leading to formation of subgrain structure (frequency of LAB increasing),c large deformation leading to formation of lamellar microstructure,d collapse of lamellar boundaries,e spheroidization beginning by Y-junction migration and f DRX grains formed by further spheroidization and growth

3.During hot deformation process,as shown in Fig.6d,e,two new nodes and form and are pulled apart by the boundary tensions when A₁ and A₂ nodes touch.After that,the grain begins further spheroidization and growth by Y-junction migration,forming DRX grain in the deformed grain,as shown in Fig.6f.

As seen from the microstructural evolution (Fig.4) and the proposed DRX model (Fig.6),DRX mechanism of AlCu-Li alloy is mainly CDRX during hot deformation.

4 Conclusion

The hot deformation behavior of Al-Cu-Li alloy was conducted in temperature range of 340-500℃and strain rate range of 0.001-10.000 s^-1.According to the analysis of processing map and micros tructural evolution mechanism,optimal hot processing conditions of Al-Cu-Li alloy are the temperature range of 400-430℃and strain rate of about 0.100 s^-1 for practical manufacturing,with efficiency of power dissipation of around 30%.the deformation mechanism of Al-Cu-Li alloy in the instable domains is flow localizations.Dynamic softening first occurs in flow localizations structure,inducing DRV and DRX non-uniform distribution in the deformed microstructure and making the workability of the alloy worse in turn.In stable domains,DRV and DRX exist simultaneously,and DRX gradually becomes the dominant deformation mechanism with the increase in deformation temperature and decrease in strain rates.DRX mechanism of Al-Cu-Li alloy is mainly CDRX during hot deformation.A DRX model of Al-Cu-Li alloy is proposed based on the microstructural evolution process of the test alloy.In the process of hot deformation,the DRX mechanism of AlCu-Li alloy is dominated by CDRX.

参考文献

[1] Pasang T,Symonds N,Moutsos S,Wanhill RJH,Lynch SP.Low-energy intergranular fracture in Al-Li alloys.Eng Fail Anal.2012;22(6):166.

[2] Qin HL,Zhang H,Wu HQ.The evolution of precipitation and microstructure in friction stir welded 2195-T8 Al-Li alloy.Mater Sci Eng,A.2015;626(3):322.

[3] Heinz A,Haszler A,Keidel C.Moldenhauer S.Benedictus R,Miller WS.Recent development in aluminium alloys for aerospace applications.Mater Sci Eng,A.2000;280(1):102.

[4] Alexopoulos ND,Migklis M,Stylianos A,Myriounis DP.Fatigue behavior of the aeronautical Al-Li(2198)aluminum alloy under constant amplitude loading.Int J Fatigue.2013;56(11):95.

[5] Li HZ,Wang HJ,Liang XP,Liu HT,Liu Y.Zhang XM.Hot deformation and processing map of 2519A aluminum alloy.Mater Sci Eng A.2011;528(3):1548.

[6] Wu HY,Wu CT,Yang JC,Lin MJ.Hot workability analysis of AZ61 Mg alloys with processing maps.Mater Sci Eng A.2014;607(9):261.

[7] Peng XN,Guo HZ,Shi ZF,Qin C,Zhao ZL,Yao ZK.Study on the hot deformation behavior of TC4-DT alloy with equiaxedα+βstarting structure based on processing map.Mater Sci Eng A.2014;605(4):80.

[8] Yin H,Li HY,Su XJ,Huang DS.Processing maps and microstructural evolution of isothermal compressed Al-Cu-Li alloy.Mater Sci Eng A.2013;586(6):115.

[9] Jagan Reddy G,Srinivasan N,Gokhale Amol A,Kashyap BP.Processing map for hot working of spray formed and hot isostatically pressed Al-Li alloy(UL40).J Mater Process Technol.2009;209(18):5964.

[10] Nayan N,Gurao NP,Narayana Murty SVS,Jha Abhay K,Pant B,Sharma SC,George Koshy M.Microstructure and micro-texture evolution during large strain deformation of an aluminium-copper-lithium alloy AA 2195.Mater Des.2015;65(10):862.

[11] Xun Y,Tan MJ.HBSD characterization of 8090 Al-Li alloy during dynamic and static recrystallization.Mater Charact.2004;52(3):187.

[12] Sakai T,Miura H,Goloborodko A.Sitdikov O.Continuous dynamic rccrystallization during the transient severe deformation of aluminum alloy 7475.Acta Mater.2009;57(1):153.

[13] Sitdikov O,Sakai T,Goloborodko A,Miura H.Grain fragmentation in a coarse-grained 7475 Al alloy during hot deformation.Scripta Mater.2004;51(2):175.

[14] Li JQ,Liu J,Cui ZS.Characterization of hot deformation behavior of extruded ZK60 magnesium alloy using 3D processing maps.Mater Des.2014;56(4):889.

[15] Jenab A,Karimi Taheri A.Experimental investigation of the hot deformation behavior of AA7075:development and comparison of flow localization parameter and dynamic material model processing maps.Int J Mech Sci.2014;78(1):97.

[16] Li DH,Yang Y,Xu T,Zheng HQ,Zhu QS,Zhang QM.Observation of the microstructure in the adiabatic shear band of7075 aluminum alloy.Mater Sci Eng A.2010;527(15):3529.

[17] Humphreys FJ,Hatherly M.Recrystallization and Related Annealing Phenomena.Oxford:Pergamon Press;2000.427.

[18] Doherty RD,Hughes DA,Humpherys FJ,Jonas JJ,Juul Jensen D,Kassncr ME,King WE.McNelley TR,McQueen HJ,Rollett AD.Current issues in recrystallization:a review.Mater Sci Eng A.1997;238(2):219.

[19] Prasad YVRK.Processing maps:a status report.J Mater Eng Perform.2003;12(6):638.

[20] Lin YC,Li LT,Xia YC,Jiang YQ.Hot deformation and processing map of a typical Al-Zn-Mg-Cu alloy.J Alloy Compd.2013;550(6):438.

[21] Kapoor R,Shekhawat SK,Samajdar I.Flow localization in an Al-2.5Mg alloy after severe plastic deformation.Mater Sci Eng A.2014;611(9):114.

[22] Reddy GJ,Srinivasan N,Gokhale AA,Kashyap BP.Characterisation of dynamic recovery during hot deformation of spray formed Al-Li alloy(UL40)using processing map approach.Mater Sci Technol.2008;24(6):725.