CHINESE JOURNAL OF MECHANICAL ENGINEERING Vol. 26, No. 3, 2013

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DOI: 10.3901/CJME.2013.03.437, available online at www.springerlink.com; www.cjmenet.com; www.cjmenet.com.cn

Deformation Behavior and Microstructure Evolution of Wrought Magnesium Alloys

WANG Shouren1, *, SONG Linghui1, KANG Sukbong2, CHO Jaehyung2, and WANG Yingzi3

1 School of Mechanical Engineering, University of Jinan, Ji’nan 250022, China 2 Korea Institute of Materials Science, Changwon 641010, Korea

3 School of Materials Science and Engineering, University of Jinan, Ji’nan 250022, China

Received September 7, 2012; revised December 12, 2012; accepted February 25, 2013

Abstract: There are many researches on the deformation behavior of wrought magnesium alloys, such as AZ31, AZ80, AZ91, and ZK60 magnesium alloys at different temperatures and strain rates, but few of them focuses on the deformation behavior of AZ41M and ZK60M alloys, especially under the twin-roll casting (TRC) state. Meanwhile, the existing researches only focus on the grain refinement law of the magnesium alloys under deformation conditions, the deformation mechanism has not been revealed yet. The hot compression behavior of AZ41M and ZK60M magnesium alloys under the temperature and strain rate ranges of 250–400 °C and 0.001–1 s–1 are studied by thermal simulation methods using Gleeble 1500 machine and virtual simulation using finite element analysis software. Simulation results show that sine hyperbolic law is the most suitable flow stress model for wider deformation conditions. The most reasonable selected deformation conditions of ZK60M alloy is 350 ºC/0.1 s–1 for TRC and 350 ºC/1 s–1 for conventional casting (CC), while AZ41M alloy is 300 ºC/0.01 s–1 for TRC and 350 ºC/0.1 s–1 for CC. Deformation behavior and dynamic recrystallization (DRX) mechanism of them are analyzed at the same deformation conditions. The microstructures of AZ41M and ZK60M alloys are observed at different deformed conditions by optical microscopy (OM) and electron back scatter diffraction (EBSD) and it reveals the flow behavior and deformation mechanism of them. Working harden and work soften contribute to the activation of basal, non-basal slip systems which promote DRX. The proposed research reveals the deformation behavior and mechanism of the AZ41M and ZK 60M magnesium alloys and concludes their optimized deformation parameters and processes and provides a theory basis for their manufacturing and application. Key words: twin rolling cast, deformation behavior, dynamic recrystallization, hot compression

1 Introduction

Magnesium alloys offer great potential application for aircraft, automotive and electronic industries due to low density and high specific strength[1–2]. They have good castability, machinability, damping capacity, shielding capability, and recyclability, which are superior to those of the competing metallic materials such as steels and Al alloys[3]. Especially, Wrought magnesium alloys exhibit higher strength, improved ductility and enhanced mechanical property advantages[4–6]. Wrought magnesium alloys are produced in a range of product forms. Extruded magnesium bars, sections and tubes are widely used in aerospace, nuclear, automotive and other engineering applications. Rolled magnesium sheet and magnesium plate are produced for seawater battery, photoengraving(PE), space, production tooling and jigging, automotive and information and communication technologies(ICT)

* Corresponding author. E-mail: sherman0158@tom.com

This project is supported by National Natural Science Foundation of China (Grant No. U1134101), and Shandong Provincial Natural Science Foundation of China (Grant No. ZR2011EMM003)

© Chinese Mechanical Engineering Society and Springer-Verlag Berlin Heidelberg 2013

applications. Magnesium forgings are used primarily for aerospace applications[7].

However, magnesium alloys have poor formability and limited ductility at room temperature because of their intrinsic characterization of the hcp structure. Activation of non-basal slips system of {1010} 1120 prismatic and {1011} 1120 pyramids at elevated temperature lead to distinctly improved formability[8–10]. Some forming technologies, such as twin-toll casting(TRC)[11] and hot compression(HC)[12], have been widely used at room or elevated temperature for wrought magnesium alloys. TRC is a promising technology compared with conventional casting(CC) because of its advantages of combining casting and hot rolling into a single operation. On the one hand, TRC is especially attractive for Mg alloys which can produce high quality Mg sheet products and provides much faster solidification rates(102–103 Ks) than CC. The relatively fast solidifcation rate of TRC has beneficial effects on the resultant microstructure, such as reducing segregation, refining microstructural features, and extending solid solubilities. On the other hand, TRC process can produce relatively thin Mg strips less than 10 mm in thickness compared to the conventional ingot

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casting producing Mg slabs of over 300 mm in thickness. Two kinds of magnesium alloys (Mg–Al–Zn, AZ series

and Mg–Zn–Zr, ZK series) are the common wrought magnesium alloys. AZ series magnesium alloys (additions of aluminum) are widely used in power generation, defence, and ICT. While ZK series magnesium alloys (non-addition of aluminum) holding high strength and ductility are used in aerospace applications including helicopter gearbox parts and wheels[13–14]. The common addition of Zn in two kinds of magnesium alloys result in increases fluidity of melt, weak grain refiner, precipitation hardening, strength and corrosion improvement at ambient temperature. The solid solution strengthening effect of Zn is much higher than Al at higher concentrations. Zr is the most effective grain refiner and incompatible with Si, Al and Mn, which can removes Fe, Al and Si from matrix. Mg–4.1Al–1.2Zn– 0.5Mn–0.5Ca alloy (AZ41M) is one of modification of base material AZ41 by adding Mn and Ca elements. And, Mg–6.2Zn–0.48Zr alloy (ZK60M) is one of modification of base material ZK60 by adding more Zn and Zr elements. Due to addition Ca and Mn in AZ41 alloy, the ignition point of magnesium alloys increase greatly[15]. Furthermore, elevated temperature properties, corrosion resistance and grain refinement are improved[13]. And, due to more Zn and Zr addition in ZK60 alloy, some purposes about grain refinement, properties improvement and impurities removal of Mg alloys are easily attained.

There are many researchers who are studying the deformation behavior of wrought magnesium alloys. For example, the deformation behavior of AZ31[16–17], AZ80[18], AZ91[19] and ZK60[20] magnesium alloys was studied at different temperatures and strain rates. However, few articles are studying the deformation behavior of AZ41M and ZK60M alloys under TRC state. In the present works, AZ41M and ZK60M alloy were formed by TRC or CC processes, subsequently deformed by the plane strain HC. The corresponding flow stress curves were recorded. The optimized deformation parameters and processes were concluded, and the hot compression behaviors of Mg alloys were analyzed by finite element software DEFORMTM 3D, and experiments were conducted to validate the simulation results.

2 Finite Element Simulation and

Experimental Procedure

2.1 Finite element modeling DEFORM-3D and Unigraphics (UG) software are used

to analyze the metal flow characteristic based on finite element method (FEM). UG software holds a kind of strong three-dimensional solid modeling capabilities. The three-dimensional model of the HCed processing was created through parametric modeling approach which is shown in Figs. 1(a) and 1(b). The FEM modeling of un-HCed and HCed samples are shown in Figs. 1(c) and 1(d). The die has a rectangular deform region and two

sidewalls (thick arrow). The other directions have zero restriction (narrow arrow). The FEM analyses adopt the following four assumptions: (1) both of the punch and die were rigid bodies; (2) the forming process was isothermal and the temperature of the model was 300 °C; (3) the constitutive relation equation was applied as sine hyperbolic law:

[sinh( )] exp ,n QART

ε ασ

(1)

where A and α are materials constants; n, stress exponent, is the slope of the sinh( ) exp( )Q RTασ ε line; R is the gas constant (8.314 J(mol • K)); Q is the activation energy (QKRn, n1m); ε is the strain rate and T is the adiabatic deformation temperature. This description for the Zener-Hollomon parameter can be suitable for every kind of deformed conditions; and (4) the constants, such as A, n, and α, were used in the description in Ref. [21].

(a) Hot compression die (b) HCed processing modeling

(c) Un-HCed sample (d) HCed sample

Fig. 1. Three-dimensional model of the HCed processing

2.2 Experimental procedure Materials used in the present work are AZ41M and

ZK60M alloys (“M” indicated as modified), which chemical compositions were shown in Table 1. The alloys were melted to 720 °C under an inert atmosphere of CO2 and SF6 mixture. After melting, the molten metal was flowed down into the preheated tundish with a temperature of 680–700 °C, followed by twin roll strip casting. The roller diameter was Φ300 mm. Roll gap was set at 3.2 mm and the roll speed was 5–6 rmin during TRC process. The dimension of TRC strip was 100 mm wide, 3.2 mm thick

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and 10 m long. The as-casted strip and CC ingots were machined to size as 50 mm20 mm3 mm. The plane strain hot compression tests of sheet sample were done using Gleeble 1500 machine. Deformation processes parameters of HC are shown in Table 2. Schematic diagram of TRC-HCed processing is shown in Fig. 2. The photo of TRC strip and HCed samples are shown in Fig. 3. The specimens were quenched using nitrogen gas immediately after hot compression to prevent the microstructures from changing after deformation.

Table 1. Chemical composition of magnesium alloys wt%

Alloy Al Zn Mn Ca Fe Si Zr Mg AZ41M ZK60M

4.2 1.2 0.5 0.5 0.01 0.002 – Bal. – 6.2 – – – – 0.48 Bal.

Table 2. Strain rate ε of AZ41M and ZK60M alloys s–1

Temperature TºC ZK60M (TRC, CC) AZ41M (TRC, CC) 300 – 0.1 – 0.001 0.01 0.1 – 350 0.01 0.1 1 – – 0.1 1

Fig. 2. Schematic diagram

of TRC-HCed processing Optical microstructures were observed after mechanical

polishing with polycrystalline diamond suspension glycol based solution. The grain structure was revealed by subsequent etching in 7 s with a solution of picric acid (1 g), acetic acid (1 mL), distilled water (2 mL), and ethanol (20 mL). All optical microstructures (OM) were observed along the transverse direction of the rolling sheets for TRC sample. Microstructural evolution and grain boundary distribution measurements were performed by electron back scatter diffraction (EBSD). The EBSD specimens were prepared using mechanical polishing and electro polishing. The mechanical polishing was done using 2000# sand paper. The electro polishing was done at 11 V for 120 s in solution of 400 mL butyl glycol, 80 mL ethanol, and 40 mL perchloric acid (HClO4) under 15 °C in order to remove surface strain. EBSD data was acquired using JEOLJSM-7001F field emission scanning electron

microscope (FESEM) and HKL Channel 5 EBSD software (HKL Technology, Denmark). The FESEM was operated at 20 kV in high vacuum mode and the specimen was tilted at 70°.

(a) TRC strip

(b) HCed samples

Fig. 3. Photo of TRC strip and HCed samples

3 Results and Discussions 3.1 Dynamic model and finite element analysis

The relationship among true stress σ, true strain ε, strain rate ε and deformation temperature T can be described as

1 2( , ) ( ).f T fσ ε ε (2)

According to BASARAN, et al[22], energy dissipations

can be divided into two terms as intrinsic dissipation and thermal dissipation. The intrinsic dissipation consists of flow deformation and structural evolution dissipation. Then, it is possible to use the following statement:

0 0d d ,P J G ε σσ ε σ ε ε σ

(3)

where P is the total energy dissipation; G is the content dissipation related to kinetic energy; J is the co-content dissipation related to potential energy.

The total energy dissipation can be divided by σ ε curve where content G represents the area under the curve and co-content J represents the area above that curve which is shown in Fig. 4. The content G and co-content J are related to a partitioning parameter, called the strain rate sensitivity (m):

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,

(ln )(ln ) T

JmG ε

ε σ σσ ε ε

. (4)

Fig. 4. Schematic representation of constitutive relation of material system as energy dissipation including nonlinear and liner energy dissipation

If the flow softening mechanism takes up a dominant

position holding a steady state, then strain rate sensitivity falls into the range of 0m1. If no strain hardening occurs, the lower limit m0 represents strain rate independent deformation for which no power is dissipated within the deformed material[23]. The upper limit m=1 means plastic deformation like a viscous fluid. So the co-content J can be denoted as

0 d .1

mJm

ε σ ε σε

(5)

Eq. (5) is based on constitutive relation abided by the

materials, ,mCσ ε when m1,

max .2

J σε

(6)

The efficiency of power dissipation is defined as

ηJJmax and can be calculated as

max

2 .1

J mJ m

η

(7)

According to Arrhenuis equation, the Zener-Hollomon

parameter is defined as[24–25]

[sinh( )] exp .n QZ ART

ασ ε

(8)

According to the linear relationship in ln lnε σ and

lnε σ , it can be calculated that the value of α. It is noted that σ value is the stress corresponding to true strain as 0.5.

The logarithmic transformation for Eq. (8) is as follows:

ln ln ln[sinh( )],Q A nRT

ε ασ (9)

1lnln[sinh( )] ,Anεασ

(10)

where 1ln .A QA

n n RT

The parameters, such as A, n, Q, and α are different with alloys kinds and deformation temperatures. They are calculated and simulated by FEM. The parameters are shown in Table 3 and Table 4.

Table 3. Parameters of ZK60

Temperature ToC

Materials constant A

Stress exponent

n

Materials constant

α

Activation energy

Q(kJ•mol–1) 25 760 3.567 330 0.006 448 121 100 760 4.258 920 0.006 667 135 200 760 4.289 050 0.006 936 136 300 760 4.439 611 0.007 413 156 350 760 4.524 804 0.006 414 166 400 760 3.665 756 0.008 888 125

Table 4. Parameters of AZ41

Temperature ToC

Materials constant A

Stress exponent

n

Materials constant α

Activation energy

Q(kJ•mol–1) 25 760 2.112 65 0.004 964 125 100 760 2.438 05 0.005 129 131 200 760 2.489 11 0.005 333 135 300 760 2.667 68 0.005 624 148 350 760 2.778 45 0.005 616 151 400 760 3.509 09 0.009 974 161

According to Eqs. (8)–(10) and Tables 3 and 4, three kind of typical flow stress curves are simulated by FEM which are shown in Fig. 5. Each route is obtained under different deformation conditions: (1) deformation temperature as 25–200 °C, stain rate as 0.1–1 000 s–1; (2) deformation temperature as 200–300 °C, stain rate as 0.01–0.1 s–1; deformation temperature as 300–400 °C, stain rate as 0.001–0.1 s–1. The flow behavior and deformation mechanism described in the figure represent as work hardening (curve I), work hardening dynamic recovery (curve II) and work softening (curve III, work hardening dynamic recovery dynamic recrystallization).

Fig. 5. Classic strain-stress curves

of wrought magnesium alloy

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Fig. 6 shows the stress distribution of ZK60M and AZ41M magnesium alloy under hot compression condition of 350 °C/0.1 s–1. It is shown that the max stress region occurs under the bottom of the punch. However, other regions (red arrows) also show bigger stress characterization. Some fractured cracks even occur in those regions (Fig. 7).

Fig. 6. Stress distribution of magnesium alloy under 350 °C/0.1 s–1 (MPa)

Fig. 7. Cracks in region A of Fig. 6(a)

3.2 Characterization of flow curves

Fig. 8 shows the stress–strain curves of the TRC and CC ZK60M alloy under different hot compression conditions. It is shown that the general characteristics of the flow stress curves seem to be nearly similar except of the deformation conditions of 350 °C/0.01 s–1. For 350 °C/0.01 s–1, the flow stress exhibits a sharp increase at the initial stage of strain and then slowly increases up to a transient equilibrium until stress arrives at the maximum value. Subsequently stress decreases and then drives to a steady value with increasing strain. Flow stress follows classical power law for TRC and CC. These curve features represent working softening and typical dynamic recrystallization (DRX). In another case,

slight work hardening is advanced at moderate strain rate while distinct working hardening occurs at high strain rate for TRC. Flow stresses do not change extensively at strains between 0.3 and 0.6. For CC, all of the stress-strain curves exhibit a quasi-steady state under the same deformed conditions. It is indicated that slight work softening is activated.

Fig. 8. True strain-true stress curves of ZK60M alloys

Fig. 9 shows the stress-strain curves of the TRCed and

CCed AZ41M alloy under different hot compression conditions. The characteristics of the flow stress curves are different from each other. Therein, the stress decreases quickly with increasing strain for 350 °C/1 s–1 and 300 °C/0.1 s–1. The values of peak stresses (σp) of TRC processes under different hot compression conditions are shown in Table 5.

Fig. 9. True strain-true stress curves of AZ41M alloy

Table 5. Peak stress and steady stress under TRC state

Peak stress and steady stress

AZ41M ZK60M 1 s–1 0.1 s–1 0.01 s–1 0.001 s–1 1 s–1 0.1 s–1 0.01 s–1

σp/MPa 300 °C – 280.7 209.2 140.2 – 295.8 – 350 °C 241.2 240 – – 272.2 116.2

σs/MPa 300 °C – 180.1 117.6 – 288.4 – 350 °C 200.1 – – 252.3 80.5

Note: – Test was not done; : Test was done but the sample failure under the deformation condition

Obviously, values of σp increase with deformation

temperature decreasing and decrease with strain rates decreasing no matter AZ41M or ZK60M alloys. However, the maximum of σp of ZK60M is bigger than that of

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AZ41M alloys. At the lower temperatures (T300 °C) and higher strain rate ( 0.1ε s–1), the flow curves of AZ41M alloys achieve to peak values and then descend quickly. However, the flow curves of ZK60M ascend all along with the increasing of strain under the same deformation conditions. At lower strain rates ( 0.01ε s–1), the flow stress increases to a peak values and then decreases to a steady state with the increasing of true strain, aiming at both AZ41M and ZK60M alloys. Such flow behavior is characteristic of hot working accompanied by DRX. The values of steady stress (σs) are also shown in Table 5. At moderate deformation conditions (T300 °C, 0.01 s–1 ε0.1 s–1), the flow curves achieve to a peak value and then decline towards to a plateau. However, the descending degree of AZ41M is higher than that of ZK60M alloys. It is indicated that AZ41M alloys is easier to deform than ZK60M alloys at the same deformation conditions.

The σp and σs values of CC are shown in Table 6. Values of σp and σs for ZK60M alloy are bigger than that of AZ41M alloy regardless CC or TRC. For CCed AZ41M alloy, the most values of σp and σs are smaller than that of TRC. However, it is interesting to find that σp and σs values of CC are bigger than that of TRC at high deformation temperature. For CCed ZK60M alloy, the data of CC are bigger than that of TRC expect of 350 °C/0.1 s–1. From Fig. 8 and Fig. 9, it is shown that AZ41M alloy under CC state are easier to get flow softening state than that of TRCed under low deformation temperature. Even through at larger strain rate, it is difficult to find the fracture failure occurred at CC state. In generally, the flow curves of CC are easier to attain the steady state than that of TRC. However, for ZK60M alloy, under moderate deformation conditions, working softening is easier for CC than TRC.

Table 6. Peak stress and steady stress under CC state

Peak stress and steady stress

AZ41M ZK60M 1 s–1 0.1 s–1 0.01 s–1 0.001 s–1 1 s–1 0.1 s–1 0.01 s–1

σp/MPa 300 °C – 268.5 180.2 110.2 – 302.7 – 350 °C 258.2 250.1 – – 266.1 127.9

σs/MPa 300 °C – 240.5 140.2 85.6 120.5 – – 350 °C 255.1 245.2 – – 131.2 226.1 87.5

Note: –: Test was not done; : Test was done but the sample failure under the deformation condition

3.3 Deformation behaviors

The flow behavior and deformation mechanism of AZ41M and ZK60M alloy at TRC and CC state as function of various deformation temperatures and strain rates were summarized in Table 7. At the lower temperatures (T 300 °C) and higher strain rate ( 0.1ε s–1), AZ41M alloys exhibit severe work hardening followed by fracture failure occurred at TRC state. Curves decline severely presenting load released (300 °C/0.1 s–1 and 350 °C /1 s–1). More grain boundaries are found in TRC ingot. Dislocations collecting are easy around the boundaries with increasing the strain rate. Once the stress achieves to a critical value, the sample

will be ruptured. There is not enough time to obtain copious dynamic recovery (DRV) and DRX at the higher strain rates[26]. However, the flow curves of ZK60M alloy have no distinct peak values under the same deformation conditions. Work hardening phenomenon is continually carrying on. With regard to CC, the work hardening mechanism plays a main role at the commencement. Once the true strain exceeds a critical value, the work hardening will be restricted and the work softening mechanism will take up a dominant position. At the lower strain rates ( 0.01ε s–1), there exhibits distinct work softening behaviors of AZ41M and ZK60M alloys. At the moderate deformation conditions (T300 °C, 0.01 s–1 ε0.1 s–1), the work hardening mechanism still playing a dominant role at the beginning of deformation and then hardening and softening reach balance when strain reaches a certain critical degree. At low strain rate of 0.001 s–1, the flow stress value reaches to a stable stress at all deformation temperature. The general curves of CC and TRC have the character of dynamic recrystallization, while most curves of CC have the obvious fluctuation. In conclusion, at high temperature and low strain rate, in despite of AZ41M or ZK60M alloys and no matter CC or TRC, the alloys follow the power-law behavior, exp( ).nA Q RTε σ While the power-law breaks down at high strain rate and low temperature. Such flow stress behaviors are typical characteristics of hot working which accompanied by DRX[21,27]. At the moderate deformation conditions (T300 °C, 0.01 s–1 ε 0.1 s–1), the exponential law

exp( )exp( ),A Q RTε βσ is applicable. For various strain rates and deformation temperatures, hyperbolic sine function, ( ) exp( ) [sinh( )] ,nf Q RT Aσ ε ασ is applicable unless the sample fractures.

Table 7. Overview of flow behavior and deformation

mechanism of the alloys as function of strain rate and temperature

State AZ41M ZK60M

1 s–1 0.1 s–1 0.01 s–1 0.001 s–1 1 s–1 0.1 s–1 0.01 s–1

TRC 300 °C ▲* ▲ ◆ ◆ ▲* ★ ◆ 350 °C ▲ ◆ ◆* ◆* ★ ★ ◆*

CC 300 °C ▲* ▲ ◆ ◆ ▲* ★ ◆ 350 °C ★ ★ ◆* ◆* ★ ◆ ◆*

Note: ★: Work hardening; ◆: Work softening; ▲: Fracture; *: Predicted result

DRX mechanisms corresponding to different strain rates

and deformation temperatures are shown in Table 8. At low temperatures and large strain rates, basal slip(BS) and mechanical twinning(MT) are readily activated. Basal dislocations with a Burgers vector accumulate near twin boundaries resulting in occurrence of some shear bands (SB)[25, 28]. Once the internal stress exceeds the limit, high angle grain boundary containing a high density of dislocation will begin to active and deviate to each other, which cause the critical stress concentration and finally result in fracture failure(FF). If the internal stress does not

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exceed the fracture limit, only has exceeded locally the critical resolved shear stress, the non-basal slip will be activated. For ZK60 alloys, however, at the commencement and at the same deformation conditions compared to AZ41M alloys, cross-slip(CS) plays an important role in activation of DRV accompanied by MT and BS. And then, twins induce dynamic recrystallization(TDRX) and low temperature dynamic recrystallization(LTDRX). They contribute to softening behavior and become the dominant deformation mechanism with strain increasing. At low temperatures and low strain rates, basal slip, twinning and local Friedel-Escaig cross-slip(FECS) are activated. A conventional DRX, dis-continuous DRX(DDRX), arising from nucleation and growth, play a leading role in these deformation conditions. Basal and non-basal slips activation causing LTDRX become the second deformation mode. At intermediate temperatures and moderate strain rate, dislocation rearrangements by CS and climbing on non-basal plane generate a low-angle boundary(LAB) network in the vicinity of original boundaries. Continuous absorption of dislocations in the LABs results in continuous dynamic recrystallization(CDRX). However, the stacking fault energy(SFE) of ZK60 alloys is low, non-basal slip is readily activated. Friedel-Escaig mechanism controls both plastic flow and nucleation of DRX[29]. At high temperatures and low stress rates, microscopic strain localization at slip lines causes formation of bulges of grain boundaries(BGB) or grain boundary migration(GBM) which leads to nucleation of DDRX grains. At moderate and higher strains, fast dislocation climbing leads to the formation of LABs which trap moving dislocations and gradually convert into true high-angle boundaries(HAB)[30]. Another mode is that CS mainly induced the occurrence of DDRX due to activation energy which corresponded to self-diffusion.

Table 8. Overview of DRX mechanisms

State AZ41M ZK60M

1 s–1 0.1 s–1 0.01 s–1 0.001 s–1 1 s–1 0.1 s–1 0.01 s–1

TRC 300 °C ▽ ▼ ○ ● ▼ ▼ ○ 350 °C ◇ ☆ □ ■ ◇ ★ □

CC 300 °C ▼ ☆ ○ ● ▼ ▼ ○ 350 °C ◆ ★ □ ■ ◇ ★ □

Note: ○: BS + FECS – MT – DRX + DDRX, under low T and low ε ; ●: BS + NS – DRX + LDRX, under low T and low ε ; □: CS – BGB – DDRX, under high T and low ε ; ■: CS – GBM – DDRX, under high T and low ε ; ▽: BS+MT – SB – FF, under low T and high ε ; ▼: CS – DRV+ MT+BS – TDRX+ LTDRX, under low T and high ε ; ☆: CS+DC – CDRX, under moderate T and moderate ε ; ★: NS – FECS – DRX, under moderate T and moderate ε ; ◇: DC – LAB – HAB, under high T and high ε ; ◆: CS – DDRX, under high T and high ε .

3.4 Microstructural observation

The flow stress behavior distinctly influences the microstructural evolution[31]. Fig. 10 shows the microstructural characteristic of AZ41M and ZK60M alloys

at 300 °C/0.1 s–1 under different initial states.

Fig. 10. Optical microscopy of AZ41M and ZK60M alloys at deformation conditions as 300 °C/0.1 s–1

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From Fig. 10, it is clear that many twins are observed in CC-HCed alloys. Some dynamic recrystallized grains are studded in twin planes. However, for TRC-HCed state, the microstructure is not clear under 300 °C/0.1 s–1. Some shear bands and deformed microstructures are observed. Fig. 11 shows the microstructural characteristic of AZ41M and ZK60M alloys at 350 °C/0.1 s–1. There exhibit distinct structural variations in microstructure of TRC samples compared to CC samples. Equiaxed DRX grains and some stretched grains are usually observed in CC state of AZ41M alloys (Fig. 11(c)). However, numerous deformed grains and a small fraction of newly developed grains are observed in CC state of ZK60 alloys (Fig. 11 (a)). For TRC, microstructure is very different from CC. Large numbers of twins are found in TRC-HCed ZK60M alloys (arrows in Fig. 11(b)). Recrystallized small grains are found along the initial grain boundaries. Many deformation structure or shear bands are found in TRC-HCed alloys (black arrows in Fig. 11(d)). Some twins were also found between two shear bands[32].

Fig. 12 shows the EBSD orientation maps of ZK60M and AZ41M alloys under 300 °C/0.1 s–1 deformation state. Orientation color code is given for the hexagonal system. The similar color of the grains denotes the smaller misorientation angle. From Fig. 12(a), there exists large fraction of coarse grains and a few of fine grains such as twins. The size of coarse grains is about 120 μm, while that of the fine grains is about 56 μm. The average grain size is 82 μm. Twins favor to occur in the grains possessing {0001} (red color) orientations. The main twin systems are extension twins such as {1012}86.3 1210 . The activation of them causes a misorientation of 86.3º between the twinned and the original part of crystal. From Fig. 12(b), the homogeneity of the grains increases. However, overall pattern quality or bands contrast decrease (occurrence of some bleak regions). This is because of poor degree of lattice perfection caused by grain boundaries, twin boundaries and higher dislocation density. The average grain sizes of TRC-ZK60M samples are 35 μm. From Fig. 12(c), deformed grains, DRXed grains and twins are observed in CC-AZ41M sample. The grain sizes of fine DRXed grains even achieve to 5–8 μm. This is because of dislocation rearrangement and pile-up caused DRX. From Fig. 12(d), the homogeneity of the grains increase all the more, the DRXed grains turn finer (2–4 μm). Although the black regions reduced, on the whole, EBSD orientation maps are still not clear.

Fig. 13 shows the EBSD orientation maps of ZK60M and AZ41M alloys under 350 °C/0.1 s–1 deformation state. Compared to Fig. 12, amounts of recrystalled grains in Fig. 13 are more than those in Fig. 12. With the increases of the temperature, DRX is easily taken place. This is consistent with the result tested by optical microscopy (Fig. 10 and Fig. 11).

Fig. 11. Optical microscopy of ZK 60 and AZ41M alloys at deformation conditions as 350 °C /0.1 s–1

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Fig. 12. EBSD of ZK 60 and AZ41M alloys at deformation conditions as 300 °C/0.1 s–1

Fig. 13. EBSD of ZK 60 and AZ41M alloys at deformation conditions as 350 °C/0.1 s–1

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With the increase of plastic deformation the grain size distribution is inclined to polarize and grains distinctly become finer. According to our experiments of hot compression at 350 °C (reduction area of 50% and strain rate of 0.1 s–1), most regions are still deformed and recrystallized grains are limited to grain boundaries or shear bands. The volume fractions of recrystallized, substructured and deformed grains collected form EBSD data are shown in Fig. 14. For CC-ZK60M/ 300 °C/0.1 s–1, they are 12%, 10% and 78%, respectively. For TRC-ZK60M/300 °C/0.1 s–1, they are 6%, 4%, and 90%, respectively. However, with increases of the temperature, dislocations rearrange forming sub-grains which have low angle grain boundaries (LAGBs), LAGBs continuously evolved into high angle grain boundaries (HAGBs) indicating occurrence of DRX and static recrystallization (SRX). High proportions of HAGBs are consistent with high deformation temperatures and moderate strain rates, whereas LAGBs are relational to severe deformed conditions.

Fig. 14. Recrystallized, substructured and deformed percent of wrought magnesium alloy

4 Conclusions (1) The hot compression behavior of AZ41M and

ZK60M alloys magnesium alloy is investigated in the temperature and strain rate ranges of 250–400 ºC and 0.001–1 s–1, respectively. The response of the material is evaluated in the twin rolling cast and conventional cast state.

(2) Obvious work hardening is revealed at temperature

lower than 300 ºC and strain rates higher than 0.1 s–1 for TRC not only ZK60M but also AZ41M alloys. However, for CC, work hardening including parts of work softening dominates the main deformation style. The flow softening occurred at high temperatures and low strain rates and the power law is followed. Nevertheless, the work hardening plays a main role at low temperatures and high strain rates.

(3) The flow stress model under different deformation conditions (strain, strain rate and deformation temperature) is put forward based on finite element methods. It has been proved suitable for wider applications at intermediate deformation conditions of AZ41M and ZK60M alloy under TRC-HC and CC-HC processing. It is clear that sine hyperbolic law is the most suitable for wider deformation conditions.

(4) Microstructure investigations reveal the flow behavior and deformation mechanism. Working harden and work soften contribute to the activation of basal, non-basal slip systems which promote DRX. DRX are dominant in not only microstructure but also basal texture evolution.

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Biographical notes WANG Shouren, born in 1966, is currently a professor at School of Mechanical Engineering, University of Jinan, China. He received his PhD degree from Shandong University, China, in 2007. His research interests include light metal, metal matrix composites and tribology. Tel: +86-531-82765476; E-mail: sherman0158@tom.com SONG Linghui, born in 1986, is currently a master candidate at School of Mechanical Engineering, University of Jinan, China. E-mail: chuyun_s@163.com KANG Sukbong, born in 1952, is currently a professor at Korea Institute of Materials Science, Korea. He is famous in the research the field of wrought magnesium alloys, special in twin roll casting technology. E-mail: sbkang@kims.re.kr CHO Jaehyung, born in 1974, is currently an associate professor at Korea Institute of Materials Science, Korea. He received his PhD degree from Seoul University, Korea. E-mail: JHCHO@kims.re.kr WANG Yingzi, born in 1976, is currently a laboratory technician at School of Materials Science and Engineering, University of Jinan, China. E-mail: mse_wangyz@ujn.edu.cn