Journal of Materials Science & Technology  2019 , 35 (6): 1198-1209 https://doi.org/10.1016/j.jmst.2018.12.001

Hot deformation behavior and workability characteristic of a fine-grained Mg-8Sn-2Zn-2Al alloy with processing map

Weili Chengabc*, Yang Baic, Shichao Mac, Lifei Wangac, Hongxia Wangac, Hui Yud

a Shanxi key Laboratory of Advanced Magnesium-Based Materials, Taiyuan University of Technology, Taiyuan 030024, China
bKey Laboratory of Interface Science and Engineering in Advanced Materials, Ministry of Education, Taiyuan University of Technology, Taiyuan 030024, China
c School of Materials Science and Engineering, Taiyuan University of Technology, Taiyuan 030024, China
d School of Materials Science and Engineering, Hebei University of Technology, Tianjin 300132, China

Corresponding authors:   * Corresponding author at: Shanxi key Laboratory of Advanced Magnesium-BasedMaterials, Taiyuan University of Technology, Taiyuan 030024, China.E-mail address: chengweili7@126.com (W. Cheng).

Received: 2018-03-30

Revised:  2018-05-7

Accepted:  2018-05-22

Online:  2019-06-20

Copyright:  2019 Editorial board of Journal of Materials Science & Technology Copyright reserved, Editorial board of Journal of Materials Science & Technology

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Abstract

The hot deformation behavior of a fine-grained Mg-8Sn-2Zn-2Al (TZA822, in wt%) alloy was investigated in the temperature range of 150-350 °C and the strain rate of 0.01-10 s-1 employing thermomechanical simulator. In most of the cases, the material showed typical dynamic recrystallization (DRX) features i.e., a signal peak value followed by a gradual decrease or to reach a steady state. The work hardening rate was found to increase with decreasing temperature and increasing strain rate, while strain rates had great effects on work hardening behavior. Meanwhile, the constitutive analysis indicated that cross-slip of dislocations was likely to be the dominant deformation mechanism. In addition, the processing map at the strain of 0.1-0.7 showed two stability domains with high power dissipation efficiencies and the optimum hot working parameters for the studied alloy was determined to be 350 °C/0.01 s-1 and 350 °C/10 s-1, at which continuous DRX (CDRX) and discontinuous DRX (DDRX) as main softening mechanism. The instability regions occurred at 200-250 °C/10 s-1 and the main flow instability mechanism was twinning and/or flow localization bands, which were prone to induce cracks and caused in-consistent mechanical properties of the alloy.

Keywords: Mg-Sn based alloy ; Hot deformation ; Processing map ; Dynamic recrystallization

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Weili Cheng, Yang Bai, Shichao Ma, Lifei Wang, Hongxia Wang, Hui Yu. Hot deformation behavior and workability characteristic of a fine-grained Mg-8Sn-2Zn-2Al alloy with processing map[J]. Journal of Materials Science & Technology, 2019, 35(6): 1198-1209 https://doi.org/10.1016/j.jmst.2018.12.001

1. Introduction

Mg-Sn-based alloys have received considerable attention in last decades due to their excellent creep resistance at elevated temperature, which are mainly attributed to the formation of Mg2Sn particles having high melting temperature (770 °C) [1,2]. Furthermore, a number of previous references suggested that addition of Zn and/or Al could further enhance the strength of extruded Mg-Sn-based alloys significantly. For instance, Sasaki et al. [3] reported that extruded Mg-5Sn-6Zn-2Al (wt%) (TZA562) alloy exhibited superior tensile yield strength (TYS) compared with Mg-Al and Mg-Zn based alloys without Sn alloying. Park et al. [4] also mentioned that extruded Mg-8Sn-1Al-1Zn (wt%) (TAZ811) alloy showed higher strength than that of AZ31 alloy. All these above-mentioned results indicated that grain boundary strengthening (i.e. fine-grained structure) played an important role in the improvement of Mg-Sn-based alloys.

Additionally, fine-grained structure could be obtained by hot extrusion due to the occurrence of dynamic recrystallization (DRX) in Mg alloys with the low stacking fault energy. Generally, DRX is beneficial to improving both workability and mechanical properties of Mg alloys [5], and the nucleation mechanism of DRX could be classified as twin-induced DRX (TDRX), continuous DRX (CDRX) and discontinuous DRX (DDRX) depending on various deformation conditions [6]. However, mostly previous work mainly concentrated on the DRX behavior of coarse-grained Mg-Sn-based alloys (i.e. as-cast or as-homogenized) during hot working process, while ignored DRX behavior of fine-grained Mg alloys [7,8]. Thus, it is very necessary to understand what happened about extruded Mg-Sn alloys with fine-grained structure during hot deformation.

Materials engineers usually get such information by utilization of processing map on the basis of the dynamic materials model, which has been proved to be an effective tool for not only optimizing hot working parameters but also understanding hot deformation behaviors of Mg alloys [9]. Prasad et al. [10] found that the optimum processing parameters of as-cast Mg-3Sn-1Ca alloy were at deformation temperature range of 350-550 °C and strain rate range of 0.0003-0.01 s-1. In addition, Rao et al. [8] studied the hot deformation behavior of as-cast and as-homogenization Mg-2Sn-2Ca alloys and figured out the related plastic deformation mechanisms. The optimum parameter hot working for as-homogenized Mg-8Sn-1Al-1Zn alloy was also investigated in our previous work [1]. However, the hot deformation behaviors of extruded Mg-Sn-based alloys in specialty have rarely reported so far. Furthermore, it has recently been reported that extruded Mg-8Sn-2Zn-2Al (wt%) (TZA822) alloy exhibits superior mechanical properties than commercial Mg alloys under the similar circumstances [11].

Therefore, the aim of the present study is to investigate the hot deformation behaviors of fine-grained TZA822 alloy through constitutive analysis and processing map construction under a wide range of compressive temperatures and strain rates, and the microstructural evolution and DRX mechanism were systematically analyzed to validate the established constitutive equation and processing map. We hope that this kind of research can provide a guideline for post processing of Mg-Sn-based alloys.

2. Experimental

Mg-8Sn-2Zn-2Al (wt%) (TZA822) alloy was made using commercially pure Mg (99.9 wt%), Sn (99.93 wt%), Zn (99.995 wt%) and Al (99.999 wt%), which were melted in an electrical furnace with CO2 and SF6 protection. The cast ingot was homogenized at 320 °C for 3 h and then at 450 °C for 24 h, followed by water-quenching. After pre-heating to 300 °C for 15 min, the machined billets with dimension of 40 mm in diameter and 50 mm in length were extruded at 300 °C and a ram speed of 0.1 mm s-1 with an extrusion ration of 16.

The cylindrical samples with a diameter of 8 mm and a height of 12 mm were machined from the extruded rod with the compression axis parallel to extrusion direction (ED). Compression tests were conducted on a thermomechanical simulator (Gleeble-3500) at the temperature range of 150-350 °C with an interval of 50 °C and in the strain rate range of 0.01-10 s-1, respectively. All samples were compressed to a true strain of 0.7 followed by water-quench in order to maintain deformation microstructure. The microstructural examinations were sectioned in the center parallel to the compression axis, followed by etching using a solution of 3 g picric acid, 10 ml acetic acid, 10 ml H2O and 50 ml alcohol. Microstructures characterization of the specimens were observed via optical microscopy (OM, Leica 2700 M), scanning electron microscopy (SEM, Mira 3XMU) equipped with energy dispersive X-ray spectroscopy (EDS), electron back-scattering diffraction (EBSD) and transmission electron microscopy (TEM, JEM-2100 F). The volume fraction of DRX (VDRX) was calculated by follows: VDRX=SDRX/(SDRX+SunDRX), where, SDRX and SunDRX represented area of DRX and unDRX regions, respectively. The SDRX and SunDRX were measured using three micrographs by the Image-Pro plus 6.0 Software. Specifically, EBSD samples were prepared by mechanical polished and twin-jet electropolishing, and EBSD analysis with a step size of 0.5 μm.

3. Results

3.1. Initial microstructure of TZA822 alloy

Fig. 1 shows the optical and SEM micrographs of fine-grained TZA822 alloy. The alloy exhibited a typical bimodal structure consisting of fine DRXed grains and coarse unDRXed grains elongated along ED (Fig. 1(a, b)). Furthermore, it could be clearly seen that the broken second phase particles in the alloy aligned along the ED in the form of stringers (Fig. 1(b)) as well. Based on EDS analysis (Fig. 1(c)), these intermetallic compounds were identified to be Mg2Sn. Meanwhile, the average grain size of the alloy was confirmed as 3.13 μm, according to the grain size distribution (see Fig. 1(d)).

Fig. 1.   Microstructure of fine-grained TZA822 alloy: (a) OM, (b) SEM, (c) the EDS results of the second phases in (b) and (d) grain size distribution.

3.2. Flow curves

The typical true stress-strain curves for fine-grained TZA822 alloy under different deformation conditions are summarized in Fig. 2. It should be noted that the curves obtained at 150 °C and 0.1-10 s-1 were terminated at relatively low strain levels due to the premature fracture of these samples. Apart from above mentioned cases, other flow stress curves initially increase to a peak value and then decrease or to reach a steady state. Such flow behavior is typical characteristic for hot working accompanied by DRX involved [12]. Especially, at a higher temperature or lower strain rate the samples presented nearly none or very little work hardening followed by flow softening reaching in a steady state. In contrast, the samples demonstrated greatly work hardening followed by flow softening leading to stress decrease or shear fracture.

Fig. 2.   True stress-true strain curves of fine-grained TZA822 alloy at different temperatures with strain rate: (a) $\dot{ε}$ =0.01 s-1; (b) $\dot{ε}$ =0. 1 s-1; (c) $\dot{ε}$ =1 s-1; (d) $\dot{ε}$ =10 s-1.

In general, an enhancement in the flow stress level and the peak stress was observed when increasing the strain rate and/or decreasing the deformation temperature. The variation of peak stress vs. strain rate and deformation temperature is shown in Fig. 3. Clearly, both the strain rate and deformation temperature have a significant influence on the flow behavior. For instance, the peak stress of the alloy decreased from 363 MPa to 75 MPa when temperature increased from 150 °C to 350 °C at a strain rate of 1 s-1. This might because of the increased mobility of grain boundaries, the promoted annihilation and rearrangement of dislocations at a higher temperature [9,13]. Therefore, nucleation and growth rate of DRX grains were boosted, resulting in decrease of flow stress. In addition, the peak stress increased as the strain rate increased at a certain temperature (see Fig. 3(b)), where the peak stress changed from 21 MPa to 123 MPa with increasing strain rate from 0.01 s-1 to 10 s-1 at 350 °C. This is due to insufficient time for DRX or dynamic recovery at very higher strain rate, which leads to larger dislocation densities, namely, a higher strain hardening effect [14]. However, the hot deformation mechanism could not be only determined by relying on the shape of true stress-strain curves. Thus, the constitutive analysis and processing map would be given more details in the following section.

Fig. 3.   Variation of peak stress for fine-grained TZA822 alloy upon different strain rate (a) and deformation temperature (b).

3.3. Constitutive analysis

The constitutive characteristics were studied to reveal the effect of deformation conditions on the flow stress. Based on the present deformation conditions, the hyperbolic equation seems to be suitable to describe the relation among strain rate ε̇, deformation temperature (T) and flow stress (σ) [15], as expressed by Eq. (1).

$\dot{ε}$=A[sinh(ασ)]n exp(- QRT) (1)

where A is the material constant, α is the stress multiplier, n is the stress exponent, Q is the activation energy of hot deformation (kJ mol-1), and R is the gas constant (8.314 J mol-1 K-1). In this study, σ was considered as the peak stress because the steady state was difficult to attain at high strain rate conditions. In the case of the low stress level (ασ<0.8) and high stress level (ασ>0.8), the strain rate could be described in Eqs. (2) and (3), respectively [16,17].

$\dot{ε}$=A1$σ^{{n}_{1}}$σn1 exp($\frac{Q}{RT}$) (2)

$\dot{ε}$=A2 exp(βσ)exp ($\frac{Q}{RT}$) (3)

where A1, A2, n1 and β are the material constants. To simplify these equations, taking natural logarithm on both sided of Eqs. (2) and (3) gives to:

ln$\dot{ε}$=lnA1+n1lnσ - $\frac{Q}{RT}$(4)

ln$\dot{ε}$=lnA2+βσ- $\frac{Q}{RT}$(5)

The linear relation can be obtained in ln$\dot{ε}$-lnσ and ln$\dot{ε}$-σ, which fitted well as shown in Fig. 4(a) and (b), respectively. The values of n1 and β could be derived from slope of the lines in Fig. 4(a) and (b), which were 10.6 and 0.064. Thus, the value of stress multiplier α was determined to be 0.006 according to α=β/n1.

Fig. 4.   Relationship between: (a)ln$\dot{ε}$-lnσ,(b)ln$\dot{ε}$-lnσ-σ, (c) ln$\dot{ε}$-ln[sinh(ασ)] and (d)ln[sinh(ασ)] -1/T.

Taking the natural logarithm on both sides of Eq. (1), would give as:

ln$\dot{ε}$=lnA+nln[sinh(ασ)]- $\frac{Q}{RT}$(6)

Activation energy Q is a measure of the minimum energy required to stimulate dislocation movement by diffusion [13]. It is usually regarded as an important physical parameter to suggest the degree of difficulty of the material deformation under various conditions. For a particular strain rate, Q could be determined by differentiating Eq. (6):

Q=R [lnε̇ln[sinh(ασ)]]T· [ln[sinhsinhασ](1/T)]]ε̇=RNS(7)

where N is the average slope of the lines in the ln$\dot{ε}$-lnsinh(ασ) plots and S is the average slope of the plots of lnsinh(ασ)-1/T at constant strain and temperature. Fig. 4(c and d) shows linear relationships of ln$\dot{ε}$-lnsinh(ασ) and lnsinh(ασ)-1/T, and thus N and S could be determined as 6.97 and 3.27, respectively. So, the average Q of the alloy could be calculated from Eq. (7) as 189.5 kJ mol-1.

In addition, the Zener-Hollomon (Z) parameter is an effective way to describe the combined effects of temperature and strain rate on deformation behavior of materials, which could be calculated by the following equation:

Z=$\dot{ε}$exp($\frac{Q}{RT}$)=A[sinh(ασ)]n (8)

Taking natural logarithms on both sides of Eq. (8), the following Eq. (9) was obtained:

lnZ=lnA+nln[sinh(ασ)] (9)

Fitting lnZ-lnsinh(ασ) plots by the least square method (see Fig. 5(a)), n and ln A in Eq. (9) were determined as 6.45 and 42.59, respectively. Then, the value of A could be obtained as 3.14 × 1018 s-1. The correlation factor of the fitted line was almost 1, suggesting that hyperbolic sine model between the Z parameter and σ is suitable for describing the hot deformation behavior of the studied alloy. Substituting the measured values of A, α, n and Q into the Eq. (1), the constitutive equation was eventually obtained as follow:

$\dot{ε}$=3.14×1018[sinh(0.006σ)] 6.45 exp(-$\frac{189500}{8.314T}$ 1895008.314T) (10)

Fig. 5.   (a) Linear relationship betweenlnZ-ln[sinh(ασ)], (b) the comparison of the calculated and measured peak stress.

In addition, according to the transformation of Eq. (1), σ can be written as follows:

σ= $\frac{1}{α}ln \{(\frac{Z}{A})^{\frac{1}{n}}+[(\frac{Z}{A})^{\frac{2}{n}}+1]^{\frac{1}{2}}\}$ (11)

Substituting the values of α, A and n into Eq. (11), the flow stress of the fine-grained TZA822 alloy could be expressed as follows:

σ=166.7ln$\{(\frac{Z}{3.14 \times})^{\frac{1}{6.5}}+[(\frac{Z}{3.14 \times})^{\frac{2}{6.5}}+1]^{\frac{1}{2}}\}$ (12)

According to above constitutive analysis, Fig. 5(b) shows a comparison between the calculated and the measured peak stress. This means that hyperbolic sine function with correlation coefficient of 0.9913 could describe the hot deformation behavior of fine-grained TZA822 alloy accurately.

3.4. Processing maps

In order to carry out more detailed studies on hot deformation of fine-grained TZA822 alloy, processing maps were developed on the basis of the dynamic materials model (DMM) [18]. In this model, the work piece was thought as a dissipater of power and the characteristics of power dissipation through microstructural changes during deformation were expressed in terms of an efficiency of power dissipation (η) given by:

η=$\frac{2m}{ m+1}$ (13)

where m (=∂(lnσ)/∂(ln$\dot{ε}$)) is the strain rate sensitivity, a function of the deformation temperature and the strain rare. The variation of the η values with different temperatures and strain rates represented a power dissipation map, which showed different domains related to specific microstructure process. Further, a continuum criterion used to define the onset of flow instability was given by Eq. (13) [19]

ξ($\dot{ε}$) = ln[m/m+1]lnε̇+ m≤0 (14)

where ξ($\dot{ε}$) is the instability parameter and flow instability occurred when this value is negative. Subsequently, the instability map could be constructed according to instability parameters (ξ($\dot{ε}$)<0) under different conditions. Finally, the processing map can be plotted by overprinting an instability map and a dissipation map.

Fig. 6 shows the processing maps at the strain of 0.1, 0.3, 0.5 and 0.7 for fine-grained TZA822 alloys. The contour represents the efficiency of power dissipation in terms of percentage and the shaded areas reveal the regimes of flow instability. It demonstrated that the power dissipation at the strain of 0.1-0.7 were similar in terms of contour shape. While the instability region was different at different strains, implying that the instability region of the studied alloy was sensitive to stain. When the strain was 0.1, the instability region consisted of three parts, low temperature-low strain rate region, medium temperature-high strain rate region and high temperature-low strain rate region. With increasing strain, the instability areas reduced from three parts to two parts (i.e. low temperature-high strain rate domain and high temperature-low strain rate domain), as shown in Fig. 6(b and c). Subsequently, when strain increased to 0.7, the instability region transformed into low temperature-high strain rate region (as shown in Fig. 6(d)).

Fig. 6.   Processing map of the fine-grained TZA822 alloy at the strain of 0.1 (a), 0.3 (b), 0.5 (c) and 0.7 (d).

In addition, it is well known that domains with high power efficiency and positive ξ values correspond to the optimum processing conditions since these factors cause good workability with defects-free microstructure (i.e. extensively DRX) [20,21]. However, the instability areas are generally characterized by crack, shear band, twinning and flow localization etc. [22].

4. Discussion

4.1. Flow behavior

Study of the work hardening behavior could provide a great insight into the deformation and restoration characteristics of the material. The work hardening rate (θ=dσdε) derived from selected curves is plotted and given in Fig. 7. As expected, the work hardening rate (θ=dσdε) increased with decreasing temperature and increasing strain rate. This was ascribed to dynamic competition between storage and annihilation of dislocations [23].

Fig. 7.   Variation of the work hardening rate (θ) with different temperatures (a) and strain rate (b).

In all the curves, the resistance to deformation decreased as flow stress increased. Specially, the work hardening behavior could be characterized by a gradual decrease at low stress followed by a fast drop linear decreased at high stress with increasing strain rate (as seen in Fig. 7(b), at 300 °C, 0.01-10 s-1), which implies the strain rate having great influences on the work hardening behavior. Then, θ decreased to zero (stress value close to saturation stress), and further decreased, suggesting that the occurrence of flow softening or the curve reached a steady stage. Generally, the flow softening is related to microstructural mechanisms such as DRX, flow instability and cracking [24], while the steady stage is attributed to a dynamic balance between dislocation generation rate and annihilation rate.

In addition, in a modified Kocks-Mecking model, the dislocation density evolution was thought as the main factor, which governed the deformation behavior in linear stages of work hardening rate curves [25]. In this stage, the instantaneous rate of the work hardening could be described according to (θ=θ0(1-$\frac{σ}{σ_{s}}$)) in which θ0 and σs is the initial work hardening rate and saturation stress at high strains, respectively. In the present study, θ0 and σs are extrapolated for various deformation conditions through fitting linear equations of work hardening. Fig. 8 is the schematic diagram of the saturation stress in stress-strain curves and work hardening rate curves.

Fig. 8.   The Schematic diagram defining and exhibiting how to calculate σs and θ0.

4.2. Deformation mechanism

The average Q of fine-grained TZA822 alloy was obtained to be 189.5 kJ mol-1, which was higher than that for lattice self-diffusion in pure Mg (135 kJ mol-1) or grain boundary diffusion (92 kJ mol-1) [26]. This suggested that the deformation behavior of the alloy was not primarily controlled by the general diffusion mechanism. Since the prismatic slip system could be operated in Mg matrix at a high temperature, the Q value must be large enough to overcome the cross-slip of screw dislocations from basal planes to prismatic planes (160 kJ mol-1) [5]. Thus, the rate controlling mechanism was likely to be cross-slip of dislocations.

In addition, the deformation mechanism could also be identified based on the stress exponent (n) during the deformation process. When the n value ranged from 4 to 6, dislocation climb is the dominant mechanism; while the value of n is greater than 6, the deformation is controlled by cross-slip of screw dislocations [27]. In this study, the calculated n is 6.45, which means that dominant deformation mechanism is controlled by dislocation cross-slipping. And this result is very close to the report by Tahreen et al. [9], who studied Mg-Zn-Mn-Y alloy showing n = 7.3.

The comparison of average Q value of the fine-grained TZA822 alloy and coarse-grained Mg alloys is summarized in Table 1. As indicated, the present fine-grained TZA822 alloy shows a higher Q value compared to previously studied coarse-grained Mg alloys such as Mg-Sn-based alloys, Mg-Al-Zn-based and Mg-Zn-based alloys. This could be ascribed to the presence of Mg2Sn particles in the matrix, causing a significant back stress. It was the long-range stress produced in the Mg matrix by a thermal barrier-liked dispersed stable particles (Mg2Sn) that restricted the dislocation motion effectively [28], resulting in a higher Q value in the studied alloy. Previous reports [1,29] also indicated that the existence of Sn, Al, and Zn elements could slow down the movement of dislocation, leading to higher Q values for basal and non-basal slip systems.

Table 1   Comparison of average Q value of the fine-grained TZA822 alloy and coarse-grained Mg alloys.

AlloyDeformation conditionsStructureQ (kJ/mol)Ref.
Temp. (°C)Stain rate (s-1)
TZA822150-3500.01-10Fine-grained189.5This study
Mg-3Sn-1Ca350-5500.1-10Coarse-grained164[10]
AZ61250-4500.001-1Coarse-grained132[28]
SiCp/AZ91270-4200.001-1Coarse-grained140[5]
Mg-3Zn-0.8Zr250-4000.001-1Coarse-grained124.6[26]

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4.3. Stability domain

The processing map (see Fig. 6(d)) exhibited two domains with higher value of power dissipation: domain I with an efficiency of 36% occurred at the temperature of 350 °C and the strain rate of 10 s-1; domain II with a peak efficiency of 44% was observed within the temperature range of 300-350 °C and strain rate from 0.01 to 1 s-1. Simultaneously, the shape of the flow curves in these two domains exhibited obvious softening and steady stages (see Fig. 2), implying the occurrence of DRX.

Fig. 9 shows micrograph of the samples deformed at 300 °C/0.01 s-1, 350 °C/0.01 s-1 and 10 s-1. It was clear that DRX occurred and the area fraction and size of grains increased as the increment of temperature (see Fig. 9(a and b)). For Mg alloys, previous research results indicated that DRX was the thermally activated process, which was more prone to occur at higher temperature due to the increase of dislocation slide, climb and cross slip as well as grain boundaries migration [28,30]. Therefore, more uniform microstructure was easier observed at a higher temperature, which corresponding to the true stress-strain curves. It should be emphasized that large fraction of DRX is desired for hot deformation because it can not only improve the hot workability but also reconstitute the microstructure [31]. When samples deformed at 350 °C/10 s-1, partially coarse grains (>3.13 μm) were presented in this region, which means that the DRX was not complete. As compared with Fig. 9(b), the average size of grains decreased from 2.65 μm to 1.76 μm with increasing strain rate (see Fig. 9(b and c)). This kind of grain refinement may be attributed to the increase of both dislocation density and cumulative strain energy upon high strain rate, and limited growth time [32]. According to Hall-Petch relationship, fine-grained structure is beneficial to improving the strength of the alloy. Therefore, the optimum processing parameters for fine-grained TZA822 alloy were determined to be 350 °C/0.01 s-1 and 350 °C/10 s-1, respectively.

Fig. 9.   Microstructures of fine-grained TZA822 alloy deformed at (a) 300 °C/0.01 s-1, (b) 350 °C/0.01 s-1 and (c) 350 °C/10 s-1.

Fig. 10 shows the EBSD orientation maps of TZA822 alloy deformed at 350 °C/0.01 s-1 and 350 °C/10 s-1. In short, high angle grain boundary (HAGB, >15°), low angle grain boundary (LAGB, <10°) and the medium angle grain boundary (MAGB, 10 -15°) are in black, green and red, respectively. As can be seen in Fig. 10 (a), some subgrains with LAGBs and several individual HAGBs appear in the interiors of original grains, suggesting the typical characteristics of CDRX [33]. Meanwhile, the cumulative misorientation from the grain interior (point A) to the grain boundary (along the dot line) is plotted in Fig. 10(c), which indicates a large increase in misorientation. Generally, due to the existence of large misorientation gradient, the growing DRX nuclei can accumulate adequate misorientations to form HAGBs [34]. Moreover, it is well known that CDRX can be identified by a misorientation increase from the center to edge of original grains and then subgrain develops [35,36]. Thus, it can be concluded that CDRX mechanism was dominant mechanism at 350 °C/0.01 s-1. Whilst, as to the samples deformed at 350 °C/10 s-1, some serrated and bulging grain boundaries with high angle can be observed (see Fig. 10(b)). These serrated and bulging grain boundaries usually possess large local orientation or strain gradients, which can act as potential sites for nucleation [6,34]. Such grain boundaries caused grain boundary migration indicate that the DRX mechanism belongs to DDRX. Note that, subgrains with LAGBs and individual HAGBs in the interiors of grains can be also clearly observed, suggesting the occurrence of CDRX phenomenon. At the same time, the cumulative misorientation cannot exceed 5°, which implied that the miorientation gradient within grains is relatively steady, as can be seen in Fig. 10(d). This was related to weak effect of progressive subgrains rotation or the accelerating subgrains transformation of LAGBS into HAGBS by means of rotation [33], in other words, possible occurrence of DDRX. Therefore, it can be concluded that both DDRX and CDRX developed simultaneously during hot deformation at 350 °C/10 s-1.

Fig. 10.   The EBSD analysis of TZA822 alloy deformed at 350 °C/0.01 s-1 and 350 °C/10 s-1: (a) and (b) grain boundary maps; (c) and (d) misorientation profiles along the line marked in (a) and (b); (d) and (e) (0001) pole figures.

In addition, mostly DRX grains adopted orientation with basal planes (nearly parallel to CD), which were similar to extrusion texture, as shown in Fig. 10(e). In general, CDRX was controlled by nucleation of new grains owing to high misorientation and development of HAGBs, and these new grains followed the texture of their parent grains [37]. While, in contrast to Fig. 10(e), texture randomization can be found in Fig. 10(f). It may be inferred that DDRX have a positive effect on weakening the texture. Similarly, previous literature also indicated that the new grain nucleation during DDRX acquired a more random texture in hot deformation Mg alloys owing to the increase in activity of no-basal slips [37,38].

In order to reveal the nucleation mechanism under 350 °C/0.01 s-1 and 350 °C/10 s-1 conditions, TEM micrographs are given in Fig. 11. The tangled dislocations produced by strain hardening occurred within the parent grains, resulting in the formation of dislocation network (see Fig. 11(a)). With further deformation, the dislocations within grains were annihilated and recombined to dislocation arrays as shown in Fig. 11(b). Fig. 11(c) indicates the formation of subgrains with LAGBs in the grains, attributed to the cross-slip or climb of dislocations. Finally, new DRX grains with HAGBs were formed by progressing subgrains rotation and absorbing the dislocations [1,6], as shown in Fig. 11(d). Thus, CDRX were confirmed as the dominant nucleation mechanism in the studied alloy at 350 °C/0.01 s-1. Notably, the bulging out part of the grain boundaries can be recognized, as show in Fig. 11(e), suggesting the occurrence of DDRX during hot deformation at 350 °C/10 s-1 [39]. At the same time, the subgrains were also observed in the original grains (see Fig. 11(f)), implying that the nucleation mechanism of subgrain rotation can occur during hot deformation. Subsequently, HAGBs further evolve by subgrains absorbing dislocations in the interior of parent grains [1,33,34]. Thus, DDRX and CDRX developed simultaneously should response to deformation condition of 350 °C/10 s-1.

Fig. 11.   The TEM micrographs of TZA822 alloy deformed at 350 °C/0.01 s-1 exhibiting: (a) dislocation network, (b) dislocations arrays, and (c) subgrains and (d) DRX grain; (e) DDRX and (f) CDRX deformed at 350 °C/10 s-1.

Fig. 12 schematically summaries the two DRX mechanisms discussed above. Route A starts with the dislocation pile-ups (Step A1) and formation of regular dislocation arrays (Step A2). Further straining then causes the development of sub-grains surround by dislocation arrays to new DRX grains separated by HAGBs (Step A3). Route B proceeds by bulging mechanism due to strain-induced local migration of original boundaries (Step B1). This mechanism is aided by the operation of no-basal slips near the original boundaries to form LAGBs that cut off the protrusions (Step B2). Subsequently, the nuclei evolved from protrusions gradually develop to high angle DRX grains according to block mobile dislocations and lattice rotation (Step B3).

Fig. 12.   Schematic summaries the two different DRX mechanisms.

Based on above analysis, on the one hand, domain I (350 °C/10 s-1) was practical when the process involves high strain rate of deformation. On the other hand, domain II (350 °C/0.01 s-1) may be suitable for the development of process strategies to achieve an optimum microstructure. In summary, the optimum parameters for hot working of fine-grained TZA822 alloy were 0.01/10 s-1 at 350 °C.

4.4. Instability domain

Fig. 13 gives the microstructure of the samples deformed at 200-250 °C/10 s-1, which corresponds to the instability domain. Obviously heterogeneous microstructure can be observed, which means the occurrence of flow localization. At high strain rate, the fine DRX grains were formed along original coarse grain boundaries in Mg alloys, which provides a path where slip can easily occur at relatively low strain [40]. Subsequently, stress concentration distributed in the coarse grain area leads to the flow localization bands.

Fig. 13.   Microstructures of TZA822 alloy deformed at: (a) 200 °C/10 s-1, (b) 250 °C/10 s-1.

Meanwhile, it was noteworthy that flow localization bands contained twinning could be observed, as shown in Fig. 13(a). It is generally accepted that twins, which occurred at low temperature in HCP metals are unfavorable to basal slip and it contributed little to ductility but is prone to induce cracks [32]. Thus, both twinning and flow localization bands may cause in-consistent mechanical properties, which should be avoided during hot deformation. Further research will be devoted to reveal the correlation between twinning and flow localization bands in Mg alloys.

5. Conclusions

(1) The flow stress of fine-grained TZA822 alloy showed typical DRX features i.e., a signal peak value followed by a gradual decrease or to reach a steady state. The work hardening rate was found to increase with decreasing temperature and increasing strain rate, while deformation temperatures have great effects on work hardening behavior.

(2) The constitutive equation of fine-grained TZA822 alloy was developed, with average Q and n being 189.5 kJ mol-1 and 6.45, suggesting that the dominant deformation mechanism was likely to be cross-slip of dislocations.

(3) The optimum post-deformation windows for fine-grained TZA822 alloy were determined as 350 °C/0.01 s-1 and 350 °C/10 s-1, at which DRX as main softening mechanism. The former DRX characterized by progressive subgrain rotation (CDRX), while latter one characterized by grain boundary bulging and progressive subgrain rotation (DDRX and CDRX).

(4) The instability regions occurred at 200-250 °C/10 s-1 and the main flow instability mechanism was twinning and flow localization bands, which were prone to induce cracks and caused in-consistent mechanical properties of the alloy.

Acknowledgments

This work was financially supported by the National Natural Science Foundation of China (Nos. 51404166, 51704209 and 51701060), the Shanxi Scholarship Council of China (No. 2014-023), the Scientific and Technological Innovation Programs of Higher Education Institutions in Shanxi (No. 2014017), the Shanxi Province Science Foundation for Youths (No. 2016021063), the Natural Science Foundation of Hebei Province (No. E2016202130), the Research Foundation from Education Department of Hebei (No. QN2015035) and the Outstanding Youth Scholar Science and Technology Innovation Program of Hebei University of Technology (No. 2015002).

The authors have declared that no competing interests exist.


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