Journal of Materials Science & Technology  2019 , 35 (11): 2423-2429 https://doi.org/10.1016/j.jmst.2019.06.008

Orginal Article

Deformation and fracture mechanisms of an annealing-tailored “bimodal” grain-structured Mg alloy

Baojie Wanga, Daokui Xubd*, Liyuan Shengc**, Enhou Hanbd, Jie Suna

aSchool of Environmental and Chemical Engineering, Shenyang Ligong University, Shenyang, 110159, China
bKey Laboratory of Nuclear Materials and Safety Assessment, Institute of Metal Research, Chinese Academy of Sciences, Shenyang, 110016, China
cPeking University, Shenzhen Institute, Shenzhen Key Lab Human Tissue Regenerate & Repair, Shenzhen, 518057, China;
dEnvironmental Corrosion Center, Institute of Metal Research, Chinese Academy of Sciences, Shenyang, 110016, China

Corresponding authors:   *Corresponding author. Key Laboratory of Nuclear Materials and Safety Assess-ment, Institute of Metal Research, Chinese Academy of Sciences, Shenyang, 110016,China.**Corresponding author. E-mail addresses: dkxu@imr.ac.cn (D. Xu), lysheng@yeah.net (L. Sheng).*Corresponding author. Key Laboratory of Nuclear Materials and Safety Assess-ment, Institute of Metal Research, Chinese Academy of Sciences, Shenyang, 110016,China.**Corresponding author. E-mail addresses: dkxu@imr.ac.cn (D. Xu), lysheng@yeah.net (L. Sheng).

Received: 2019-05-17

Revised:  2019-06-11

Accepted:  2019-06-13

Online:  2019-11-05

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

Through investigating and comparing the mechanical behavior of an as-rolled Mg-3%Al-1%Zn (wt%) alloy before and after annealing treatments, it was revealed that the formation of annealing-tailored bimodal grain structure ensured the 330 °C/4 h samples having a good combination of tensile strength and plasticity. Failure analysis demonstrated that for the as-rolled and 330 °C/1 h samples with fine grain structure, their plastic deformation was mainly attributed to basal slips, whereas the deformation mechanism in the bimodal grain-structured samples was dominated by basal slips in fine grains and twinning in coarse grains. For the 330 °C/8 h samples with coarse grain structure, high densities of twins were activated. Meanwhile, basal slips occurred in the twinned and un-twinned areas of coarse grains and could pass through twin boundaries. For differently treated samples, cracking preferentially occurred along slip bands, resulting in their transgranular fractures.

Keywords: Mg alloy ; Grain structure ; Deformation mechanism ; Cracking ; Fracture

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Baojie Wang, Daokui Xu, Liyuan Sheng, Enhou Han, Jie Sun. Deformation and fracture mechanisms of an annealing-tailored “bimodal” grain-structured Mg alloy[J]. Journal of Materials Science & Technology, 2019, 35(11): 2423-2429 https://doi.org/10.1016/j.jmst.2019.06.008

1. Introduction

To improve the potential competitiveness of Mg alloys being as the lightweight structural materials for applications in the fields of automobile and aerospace industries, the basic requirement for their service safety is to have enough mechanical properties in terms of strength and plasticity [[1], [2], [3], [4], [5], [6], [7], [8], [9], [10], [11], [12], [13], [14]]. Besides the alloying and heat treatment, severe plastic deformation (SPD) processing methods such as equal channel angular pressing (ECAP) [15], hot rolling [16], extrusion [[17], [18], [19]] and etc were indispensable for improving mechanical properties of Mg alloys. For the SPD-processed Mg alloys, the main challenge is to form ultrafine/fine grain structure [[20], [21], [22], [23]]. Due to the formation of ultrafine/fine grain structure, Mg alloys could have obviously higher mechanical strengths, but a lower plasticity when compared with those with coarse grain structure [20,22,24]. Recently, researchers reported that through developing a bimodal grain structure in metallic materials, high strength and good ductility can be simultaneously achieved because ultrafine/fine grains provide the strengthening effect and coarse grains enable the strain hardening [[25], [26], [27]]. For Mg alloys, the formation of bimodal grain structures had the similar effect on their mechanical improvement [[28], [29], [30]]. Zha et al. reported that for an as-rolled Mg-9Al-1Zn (wt%) alloy, the formed bimodal grain structure consisted of coarse grains with the size larger than 70 μm and fine recrystallized grains with the size smaller than 5 μm ensured a good combination of strength and plasticity [30]. Since the alloy reported in previous work contained a high content of Al and was performed SPD at different temperatures, its mechanical improvement was not only ascribed to the formed bimodal grain structure, but also related to the crystallographic texture and precipitated Mg17Al12 particles in the matrix [30]. Therefore, in order to disclose the individual role of bimodal grain structure on the mechanical properties of Mg alloys, it is better to investigate a dilute Mg alloy with a lower solute content and being performed a certain SPD process.

For the typical bimodal grain-structured fcc (face-centered cubic structure) metals, the dislocation density in ultrafine/fine grains saturated quickly with increasing strain, whereas coarse grains could provide more space for accommodating the newly generated dislocations during plastic deformation [31]. However, for the hcp (hexagonal close-packed) structured Mg and its alloys, the main deformation mechanisms are basal slip and {10-12} deformation twinning at room temperature because the critical resolved shear stress (CRSS) for prismatic slip is about 100 times higher than that for basal slip at room temperature [[32], [33], [34]]. Additionally, the activation of deformation twinning in Mg alloys becomes much easier with increasing the grain size [[35], [36], [37]]. Meyers et al. developed a constitutive approach and calculated that for hcp-structured metals, the activation of deformation twinning in large grains with the size of ∼100 μm is much easier than that in small grains with the size of ∼3 μm [38]. In the research about the mechanical behavior of a commercial AZ91 alloy, Lahaie et al. reported that the deformation twinning was activated when the grain size was 15 μm, but can hardly be activated when the grain size was 1 μm [33]. Thus, the activation of {10-12} twins in coarse grains must be considered for understanding the effect of bimodal grain structure on the mechanical behavior of Mg alloys.

Based on the description mentioned above, it indicates that although the bimodal grain structure was efficient for maintaining a moderate work hardening and enhancing the ductility [[28], [29], [30], [31]], previous work seldomly concerned about the difference in deformation mechanisms especially the activation of twins in bimodal grain-structured Mg alloys. Moreover, Li et al. reported that for an as-extruded AZ31 alloy during the compression testing, the required activating stress for {10-12} twins in bimodal grains was much lower than that in the equiaxed grains, whilst the coarse grains in the bimodal grain structure can effectively suppress the crack growth [39]. Following this, two questions can be easily proposed: (1) During the tensile testing, what are the dominant deformation mechanisms in fine and coarse grains of the bimodal grain-structured Mg alloys? (2) How does the difference in deformation mechanisms influence mechanical properties and fracture characteristics? In this work, through investigating and comparing the deformation mechanisms and mechanical behavior of an as-extruded Mg-3Al-1Zn (in wt.%) alloy with fine, bimodal and coarse grain structures, these two proposed questions were answered. Additionally, the surfaces and fractures of failed samples were analyzed to disclose the fracture mechanisms of differently grain-structured Mg alloys.

2. Experimental

The investigated material in this study was an as-rolled AZ31 (Mg-2.87Zn-0.95Al (wt%)) plate with a rolling ratio of 20. To avoid the overheating associated with MgZn phases and quick growth of grains, samples were cut from the plate and annealed at 330 °C for 1 h, 4 h and 8 h, respectively. To reveal the grain structure, sample surfaces were polished and etched with an etchant consisted of (25 ml ethanol + 2 g picric acid +5 ml acetic acid +5 ml deionized water). Then, the microstructures of samples were observed with an optical microscope (OM) and the grain sizes were determined on the basis of the linear interception method. By using Image J software, the average surface area fractions of fine and coarse grains in differently annealed samples were measured. To determine the texture evolution during annealing treatment, the intensity contours of {0002} pole figures of differently annealed samples were measured by the Schultz reflection method of X-ray diffractometer (XRD: D/Max 2400) analysis.

Tensile samples with a gauge length of 25 mm, a cross-sectional area of 6 mm (in width) ×3 mm (in thickness) were machined from the plate and their axial direction was parallel to the rolling direction (RD). To reveal the microstructural changes before and after tensile tests, one side of wide surface of tensile samples was etched and then tested on the MTS (858.01 M) testing machine at a strain rate of 1 × 10-3 s-1 at room temperature. For differently annealing-treated conditions, three repeated tensile tests were performed. Afterwards, the overall fracture surfaces and typical fracture characteristics were observed using SEM (XL30-FEG-ESEM). Moreover, the surfaces with the distance less than 1 mm to fracture sites were observed by using OM and SEM to characterize the deformation and fracture mechanisms of differently annealed samples.

3. Results

3.1. Microstructure

Fig. 1 shows the grain structures and {0002} pole figures of differently annealed samples. It can be seen that the as-rolled samples have a fine grain structure and the average grain size is 3 μm (Fig. 1(a)). After annealing treatment, heterogeneous grain growth occurred and the fraction of coarse grains with the size of 20 ∼60 μm increased gradually with prolonging the holding time. For the 330 °C/1 h samples, only a few of coarse grains are observed (Fig. 1(b)). For the 330 °C/4 h samples, a typical bi-modal grain structure forms with the coarse grains almost occupying half of the matrix (Fig. 1(c)). After being annealed at 330 °C for 8 h, the samples have a typical coarse grain structure with the average grain size of 30 μm (Fig. 1(d)). The average surface area fractions of coarse grains in rolled, 330 °C/1 h, 330 °C/4 h and 330 °C/8 h samples were measured to be 0%, 15%, 45% and 100%, respectively. Based on the {0002} pole figures, it can be seen that the as-rolled samples have a typical bi-texture, i.e. strong fiber texture (the c-axis of most grains having a titled angle towards the RD of 0 ∼45°) and weak basal texture (a small amount of grains with their c-axis having a tilted angle of 0 ∼30° with respect to the normal direction (ND) of the plate). After annealing treatment, the intensity of fiber texture decreased but the intensity of basal texture increased gradually with prolonging the annealing time. In the investigation about the correlation of microstructural evolution and formation of basal texture in a coarse-grained Mg-Al alloy during hot rolling, Jin et al. reported that the grains with the c-axis parallel to ND was stable and could consume the grains with other orientations for their growth during dynamic recrystallization [40]. Therefore, the increased intensity of basal texture in annealed samples was mainly ascribed to the formation of coarse grains.

Fig. 1.   Grain structures and {0002} pole figures of (a) as-rolled, (b) 330 °C/1 h, (c) 330 °C/4 h and (d) 330 °C/8 h, respectively. The {0002} pole figures are inserted in (a)-(d).

3.2. Mechanical property growth mechanism

Fig. 2 shows the mechanical comparison between differently annealed samples. It can be seen that the as-rolled samples have the higher tensile strength, but their plasticity is relatively lower. For the 330 °C/1 h samples, the formation of a small fraction of coarse grains causes slight decrease in strength and small increase in plasticity. With the fraction of coarse grains increasing, the strength of 330 °C/4 h samples basically has no change, but their elongation ratio is 2 times higher than that of the as-rolled samples. When the coarse grain structure is formed, the strength and plasticity of 330 °C/8 h samples are the lowest compared with those of other conditions.

Fig. 2.   Comparison of tensile properties between differently annealed samples.

3.3. Failure analysis

To reflect the deformation mechanisms occurred in differently annealed samples, the wide surfaces near to fractures were observed, as shown in Fig. 3. For the as-rolled and 330 °C/1 h samples, high densities of slip bands can be observed (Fig. 3(a) and (b)). For the 330 °C/4 h samples with a bimodal grain structure, the deformation mechanism is dominated by basal slips in fine grains and twinning in coarse grains (Fig. 3(c)). For the 330 °C/8 h samples with coarse grain structure, lots of activated twins and slip bands in both the twinned and un-twinned areas can be observed (Fig. 3(d)). In addition, the microstructural comparison demonstrates that the density of slip bands or twin bands in the bimodal grain-structured samples is much lower than that in fine or coarse grain-structured samples, indicating that the plastic deformation in the gauge section of 330 °C/4 h samples is relatively homogenous. High-magnification SEM images show that for differently annealed samples, the micro-cracks preferentially nucleate along the slip bands (Fig. 3(e) and (f)).

Fig. 3.   Optical microstructures of the surfaces near to fractures for (a) as-rolled, (b) 330 °C/1 h, (c) 330 °C/4 h and (d) 330 °C/8 h; (e, f) SEM images of fine and coarse grains in 330 °C/4 h samples.

3.4. Fractography

The observations to fracture surfaces of differently annealed samples are shown in Fig. 4. Compared with the overall fracture surfaces, it can be seen that the fracture surfaces of as-rolled and 330 °C/8 h samples are quite flat (Fig. 4(a) and (d)), whereas the fracture surfaces of 330 °C/1 h and 330 °C/4 h samples are relatively rough (Fig. 4(b) and (c)). High-magnification images reveals that for the fine and bimodal grain-structured samples, plastic dimples can be obviously observed on fracture surfaces (Fig. 4(e)-(g)), whereas the fracture surface of coarse grain-structured samples is covered with brittle quasi-cleavage facets (Fig. 4(h)). Thus, the plasticity of 330 °C/8 h samples is the lowest. Moreover, for the differently annealed samples, their cracking modes are transgranular.

Fig. 4.   Fractographies of (a) as-rolled, (b) 330 °C/1 h, (c) 330 °C/4 h and (d) 330 °C/8 h samples, respectively; (e-h) high-magnification observations of the squared areas in (a-d), respectively.

4. Discussion

Based on the Hall-Petch equation [41,42], it is reasonable for the as-rolled samples having a relatively higher tensile strength. Since the low dislocation storage efficiency inside ultrafine/fine grains causes the massive loss in plasticity of Mg alloys at room temperature [23], the plasticity of as-rolled samples is relatively poor. After annealing treatment, the increased fraction of coarse grains induced the gradual decrease in tensile strength. However, the bimodal grain-structured samples have a good combination of obviously higher plasticity and maintained strength when compared with fine and coarse grain-structured samples, which should be closely related to deformation mechanisms occurred in fine and coarse grains. To have a good understanding about the difference in deformation and fracture mechanisms between differently annealed samples, a schematic model has been proposed, as shown in Fig. 5.

Fig. 5.   Schematic models for the microstructural evolution at initial (stage I), plastic strain (stage II), cracking (stage III) stages of differently grain-structured samples during tensile process for (a) fine, (b) bimodal and (c) coarse grain structure, respectively.

For pure Mg, it has been widely accepted that the CRSS value for basal slips was between 0.5 and 1 MPa, whereas the CRSS value for the activation of twins varied from 4 to 10 MPa at room temperature [43,44]. Moreover, the extra driving force required for the formation of twin interfaces increased with the grain size decreasing [45]. Therefore, it can be predicted that the CRSS value for twinning activation should decrease with the increase of grain size. For the as-rolled samples with fine grain structure, the twinning can be hardly activated. Thus, their plastic deformation and fracture mechanisms are mainly related to basal slips (Fig. 5(a)). Previous work demonstrated that for the as-forged Mg-Zn-Y-Zr alloy, the {10-12} twins were widely activated in coarse grain structure but can hardly be observed in fine grain structure during fatigue process [37]. Moreover, in coarse grain-structured Mg and its alloys, basal slips occurred simultaneously in twinned and un-twinned areas [35,37]. Nave et al. reported that the theoretical maximum extension of 6.4% along the c-axis can be provided by the {10-12} twinning [46]. Since the formation of coarse grains contributed to the increased intensity of basal texture, the c-axis of most coarse grains should be parallel to the ND direction. In general, the {10-12} twins can be activated when the tension is along the c-axis or compression is perpendicular to the c-axis of grains. Then, for the bimodal grain-structured samples, the plastic deformation and fracture mechanisms are explained as follows (Fig. 5(b)). Due to the slightly larger CRSSs for both basal and non-basal slips in fine grains than those in coarse grains [47,48], basal slips could occur in fine and coarse grains at the initial stage of tensile process. When the tensile strain was increased, the activation of {10-12} twins started in coarse grains and basal slips subsequently occurred in the twinned areas, which could provide extra space for accommodating the accumulated dislocations and then ensured the better plasticity of 330 °C/4 h samples. Based on the discussion mentioned above, the plastic deformation and fracture mechanisms of 330 °C/8 h samples with a coarse grain structure can also be explained, i.e. (1) basal slips occur when the tensile strain is small; (2) the activation of {10-12} twins and basal slips of in the twinned areas can simultaneously occur when the tensile strain is large (Fig. 5(c)). Since the plastic strain at final stage of the tensile process is dominated by basal slips, the cracks in differently annealed samples will preferentially present along slip bands.

5. Conclusion

Through investigating and comparing the deformation mechanisms and mechanical behavior of differently annealed samples, the conclusions can be summarized as follows. The formation of bimodal grain structure is helpful for improving the plasticity and simultaneously maintaining the mechanical strength of wrought Mg alloys, which is mainly ascribed to the coordinated plastic deformation mechanisms occurred in fine and coarse grains.

Acknowledgements

This work was supported financially by the Project from the Strategic New Industry Development Special Foundation of Shenzhen (No. JCYJ20170306141749970), National Natural Science Foundation of China Projects under Grant [Nos. 51871211 and 51701129], the funds of International Joint Laboratory for Light Alloys, the National Key Research and Development Program of China under Grant [No. 2017YFB0702001], Liaoning BaiQianWan Talents Program, the Innovation Fund of Institute of Metal Research (IMR), Chinese Academy of Sciences (CAS).


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