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Microstructure, texture and mechanical properties of Mg-3.0Zn-0.2Ca alloys fabricated by extrusion at various temperatures
Journal of Alloys and Compounds 652 (2015) 122e131
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Microstructure, texture and mechanical properties of Mg-3.0Zn-0.2Ca alloys fabricated by extrusion at various temperatures Cheng-jie Li a, *, Hong-fei Sun a, Xue-wen Li b, Jun-ling Zhang a, Wen-bin Fang a, b, Ze-yi Tan a a b
School of Materials Science and Engineering, Harbin Institute of Technology, Harbin, 150001, China School of Materials Science and Engineering, Harbin University of Science And Technology, Harbin, 150080, China
a r t i c l e i n f o
a b s t r a c t
Article history: Received 3 July 2015 Received in revised form 25 August 2015 Accepted 26 August 2015 Available online 1 September 2015
The Mg-3.0Zn-0.2Ca (wt.%) alloy has been extruded at temperature range of 25e300 C and the resulting microstructure, texture and mechanical properties are systematically investigated. The results show that the grain size monotonically increases with the increasing of the extrusion temperature and the texture intensity increases firstly and decreases subsequently. In addition, a large number of nano-scale precipitates are formed in alloys extruded above 250 C. The weakest basal texture developed in the cold extrusion alloy is related to the deformation twinning, while, the combining effects of activation of multiple deformation mechanisms and dynamic precipitates contribute to developing the weaker basal texture in alloys extruded above 250 C. A sharp basal texture enhances the yield strength of alloy extruded at 150 C greatly at the cost of work hardening rate. The highest elongation is achieved in the alloy extruded at 300 C. The enhanced formability is considered to be associated with the dependence of the weaker basal texture and dynamic precipitates on the strain hardening behavior. Among the multiple strengthening mechanisms, the fine grain strengthening and the solid solution strengthening play a dominant role in the increment of strength for the as-extruded alloys. © 2015 Elsevier B.V. All rights reserved.
Keywords: Magnesium alloy Extrusion Tensile properties Strengthening mechanisms Ductility
1. Introduction The application of wrought magnesium alloy is still limited due to their poor ductility and formability at room temperature [1]. Great deals of efforts have been devoted to improving the roomtemperature ductility of magnesium. Element alloying is an effective method to improve mechanical performance of magnesium alloys in nature, while thermo-mechanical process can enhance the formability of magnesium by grain refinement from the technical point of view. Literature have been documented about the obvious enhancement in mechanical properties of magnesium alloys by element alloying [2,3]. Moreover, the addition of rare-earth (RE) alloying elements (such as Gd, Ce, Y, etc.) to magnesium is believed to be one of the potential ways to enhance the formability of magnesium [4e6]. More random textures can be developed in magnesium alloys containing RE elements during hot-working [7]. However, large amount of expensive alloying elements will
* Corresponding author. E-mail address: [email protected] (C.-j. Li). http://dx.doi.org/10.1016/j.jallcom.2015.08.215 0925-8388/© 2015 Elsevier B.V. All rights reserved.
inevitably increase the materials cost, which is undesirable for industrial application. Another popular method is alloying with element of Li, which decreases the c/a axial ratio of lattice structure and makes non-basal slip systems easily activated at room temperature [8]. Therefore, the high elongation can be obtained in the magnesium-lithium alloys at room temperature [9]. However, the low strength and poor corrosion resistance severely limit their application as the structural components [10,11]. During the past several years, many researchers have devoted to developing a low-cost MgeZneCa alloy for structural and medical applications [12e14]. Such magnesium alloy has high ductility accompanied with high strength and good biocompatibility. In fact, thermo-mechanical processes have been adoped to refine the grain sizes in such magnesium alloys [15,16]. Grain refinement is desirable for their strength and ductility. In addition, a weak basal fiber texture can be developed in hot-extruded MgeZneCa alloy due to the dynamic recrystallization. Both the weakened basal fiber textures and refined grains are favorable to increase roomtemperature ductility of magnesium [15]. At present, many researchers focus on optimizing the extrusion parameters (such as extrusion temperature, ram speed and extrusion ratio) to obtain
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fine grains and random textures [12e17]. Among these variables, extrusion temperature is the most critical parameter that directly affects microstructure, texture evolution and mechanical properties of MgeZneCa alloy. Zhang et al. reported that the excellent mechanical properties in Mg-1.0Zn-0.5Ca alloy can be obtained by hot extrusion because of the grain refinement and the weakened basal texture [15]. They have also proved that the grain refinement method together with the micro-alloying concept was a very effective route for fabricating a high performance MgeZneCa alloy. To date, most of the previous works are concentrated on the hot extrusion behaviors of MgeZneCa alloys. That is, most of these alloys are extruded above 250 C. But, the systematic investigation on the low-temperature (<200 C) extrusion behaviors of MgeZneCa alloys has rarely been reported. Thus, the aim of this study is to investigate the relationship between extrusion temperatures (25e300 C) and microstructure as well as mechanical performance of MgeZneCa alloy. 2. Experimental procedure The as-cast Mg-3.0Zn-0.2Ca (wt.%) alloy used in the experiment was synthesized with pure Mg (99.95 wt.%), Zn powder (99.7 wt.%) and Mg-30 wt.%Ca master alloys. Under protection with a mixed gas atmosphere of CO2 and SF6 (CO2:SF6 ¼ 40:1 vol.%), the Zn powder and MgeCa master alloys were added into the melted pure Mg at 750 C. The melted alloys were isothermally held at 750 C for 30 min to ensure that all the alloying elements dissolved. After that, the melted alloy was poured into a stainless steel mold preheated to 200 C and cooled in the air. The as-cast ingots were solution-treated at 400 C for 10 h, followed by water quenching. The prepared billets were machined into cylindrical bar with the diameter of 52 mm and height of 50 mm for extrusion. Then, the prepared billets were extruded at different temperatures of 25 C, 150 C, 250 C and 300 C. The extrusion ram speed was set as 17 mm/s and the extrusion ratio was 8.3. Finally, the as-extruded Mg alloy bars with a diameter of 18 mm were obtained. The microstructure of as-extruded alloys was observed in an Olympus GX71 optical microscope (OM), and a JEM-2000EX transmission electron microscope (TEM). The grain size was measured using Image-ProPlus 6.0 software. The macro-texture was measured by the Schulz reflection method using X'pert PRO XRD. All the samples for XRD test were cut from center of asextruded alloys and sectioned in parallel to the extrusion direction (ED). To check the repeatability of the results, two experiments were conducted under each set of conditions. The micro-texture was examined by electron backscattering diffraction (EBSD) in Quanta 200FEG emission scanning electron microscope (SEM) equipped with OIM Analysis 5.0 software. The phase identification and analysis were performed on a D/max-rb X-ray diffractometer (XRD) using monochromatic Cu Ka radiation. The tensile tests were carried out by an Instron-5569 standard testing machine at an initial strain rate of 1.67 103 s1 at room temperature. The dogbone tensile specimens with a diameter of 3 mm and gage length of 15 mm were used. All the samples were sectioned in parallel to the extrusion direction (ED). To check the repeatability of the results, three experiments were conducted under each set of conditions. 3. Results and discussion 3.1. Microstructure Fig. 1 shows the microstructures of Mg-3.0Zn-0.2Ca alloys before and after solution treatment. As seen from Fig. 1 (a), the microstructure of as-cast alloy consists of a-Mg matrix and eutectic compounds (marked by white arrows). Most of a-Mg present
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equiaxed grains with an average grain size of 449 mm. After solution treatment, a single phase microstructure can be obtained as shown in Fig. 1 (b), because the net-like interdendritic eutectic structures almost disappear and dissolve into the matrix. The average grain size of solution treated Mg-3.0Zn-0.2Ca alloy is about 453 mm. There is little difference in grain size of the alloys before and after solution treatment. From the SEM micrograph (Fig. 2 (a)), it can be seen clearly that the second phases (in white) exhibit continuous network structure and mainly distribute along grain boundaries and inside the grains. After homogenizing annealing at 400 C for 10 h, no intermetallics can be observed in Fig. 2 (b). This result is in agreement with the result observed in Fig. 1 (b). The result from XRD shows that the as-cast Mg-3.0Zn-0.2Ca alloy contains three phases, a-Mg, MgZn2 and Ca2Mg5Zn13 (Fig. 3 (a)). After solution treatment, only a-Mg can be detected (Fig. 3 (b)). In view of microstructure (Fig. 2 (a)), it can be expected that the primary dendrites should be a-Mg and the white phases may be the MgZn2 and Ca2Mg5Zn13. The EDS analysis indicates that the average Zn/Ca atomic ratio of the white phases (shown by red crosses A, B and C in Fig. 2 (a)) is about 15.8:1 (Fig. 4). The Zn/Ca atomic ratio of Ca2Mg5Zn13 phase is 6.5:1, which is far less than 15.8:1. That is, the white phases contain abundant zinc element. Thus, it can be concluded that the ternary MgeZneCa phases are MgZn2 and Ca2Mg5Zn13 mixed phases. Fig. 5 shows the microstructural characteristics of as-extruded alloys at different extrusion temperatures. Compared with the microstructure of as-cast billet, the grain size of as-extruded alloys is refined remarkably. The as-extruded alloys exhibit uniform and equiaxed grains. It indicates that a complete dynamic recrystallization process can be obtained at all three above-mentioned extrusion temperatures. As the deformation temperature increases from 25 C to 250 C, the average grain size increases from 8.1 mm to 19.5 mm. However, as the deformation temperature further rises to 300 C, the average grain size only increases from 19.5 mm to 20.6 mm. In other words, the increase of grain size is not obvious. Therefore, it can be concluded that the microstructure is sensitive to deformation temperature when the samples are extruded at the lower temperatures (25e250 C). The extrusion temperature not only affects the grain size but also influences the grain size distribution. In the present study, at least over six hundred grains from optical micrographs are analyzed for each sample obtained at different extrusion temperatures. Fig. 6 shows the area-weighted grain size distributions, the average grain size (dm) and the relative grain size dispersion △d/dm of the as-extruded samples. Regarding the microstructural heterogeneity, the relative grain size dispersion △d/dm is used to characterize grain size dispersions according to reference [18], where d is the grain size, △d is the absolute grain size range (△d ¼ dmaxdmin) and dm is the mean grain size. All of the deformed materials show the approximate normal distribution with a single peak in spite of the difference in the peak values. In addition, there is a slight difference in relative grain size dispersion. The most well-distributed grain size is obtained in the cold extruded sample with the smallest △d/dm of 2.47. Further increase in the deformation temperature introduces some grains uneven growth, which leads to microstructure heterogeneity. However, in view of microstructure (Fig. 5) and statistical result (Fig. 6), the △d/ dm is only a slight increase of ~0.2 compared with that of cold extrusion when the extrusions are carried out at the higher temperature range of 150e300 C. It is evident that the extrusion temperature has a less influence on the grain size dispersion than average grain size. To further analyze the microstructure of extruded specimens, TEM tests are performed. Fig. 7 shows TEM images of the alloys extruded at different temperatures. As shown in Fig. 7 (a)e(e), the
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Fig. 1. Optical micrographs of the Mg-3.0Zn-0.2Ca alloy: (a) as-cast; (b) solution treated.
Fig. 2. SEM micrographs of the Mg-3.0Zn-0.2Ca alloy: (a) as-cast; (b) solution treated.
exquiaxed grains with few dislocations recognized as recrystallized structure are observed in alloys extruded below 150 C. It suggests that the dynamic recrystallization occurs during the extrusion process even at room temperature. Such a phenomenon has been observed in AZ31 alloy, high purity Mg and Mg-5.25wt.%Zn-0.6wt.% Ca after severe plastic deformation at room temperature [19e21]. Jotoku et al. suggested that the cross-slip and annihilation of dislocation resulted from the stress concentration at local area due to the pile-up of dislocations, which might cause the recrystallization [20]. In the present study, some dislocations are accumulated at twin boundaries and in the original grains, both of which
Fig. 3. X-ray diffraction patterns of Mg-3.0Zn-0.2Ca alloy: (a) as-cast; (b) solution treated.
are benificial to form some dislocation cells (Fig. 7 (c) and (d)). With increasing plastic deformation, the misorientation between the cells is increased and so the rate of growth of the sub-grain increases. Furthermore, this process results in continuous dynamic recrystallization in some area with high stress concentration. In addition, this process can be accelerated by heating generated from plastic deformation. As indicated by white arrows in Fig. 7 (c) and (d), some submicron twins are also observed, which indicates that deformation twinning plays an important role in the magnesium alloys extruded at lower temperatures (150 C). Because magnesium alloys lack sufficient independent slip systems at lower temperature. Further increase in the extrusion temperature
Fig. 4. EDS analysis of the interdendritic phases marked by red crosses A, B and C in Fig. 2 (a). (For interpretation of the references to colour in this figure legend, the reader is referred to the web version of this article.)
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Fig. 5. Optical micrographs of the as-extruded alloys: (a) 25 C, (b) 150 C, (c) 250 C, (d) 300 C.
(250 C), more slip systems can be activated. Thus, only a few deformation twins with size less than 250 nm in width can be observed in Fig. 7 (f) and (h). As shown in Fig. 7, the extrusion temperature has profound influence on second phases. Very few second phases can be observed in samples extruded below 150 C (as shown in Fig. 7(aee)). Comparatively, a large number of precipitates with
diameter of 5e30 nm dispersed within the grains are observed in samples extruded above 250 C. These second phases formed during the extrusion process are confirmed to be MgZn2 by TEM diffraction patterns. A similar phenomenon has been reported in other Mg alloys at higher extrusion temperatures (300 C) [22e24]. Generally, the high-energy sites (such as dislocations) easily are generated by hot plastic deformation, which are
Fig. 6. Grain size distributions of the as-extruded alloys: (a) 25 C, (b) 150 C, (c) 250 C, (d) 300 C.
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Fig. 7. The TEM micrographs of the as-extruded alloys: (a) the bright field image of alloy extruded at 25 C and (b) corresponding dark field image of (a), (c) bright field image of deformation twins in alloy extruded at 25 C, (d) bright field image of alloy extruded at 150 C and (e) corresponding dark field image of (d), (f) bright field image of alloy extruded at 250 C and (g) corresponding dark field image of (f), (h) bright field image of alloy extruded at 300 C and (i) corresponding dark field image of (h).
conducive to the formation of fine precipitates. These fine precipitates in turn exert a pinning effect on the movement of grain boundaries and suppress dynamic recystallization. However, combining optical and TEM micrographs, the dynamic precipitates are formed in the dynamic recystallized regions. That is, the fine precipitates cannot effectively suppress dynamic recystallization. Comparatively, the dynamic recystallization may stimulate dynamic precipitation (DRXSP). A similar phenomenon of DRXSP has been reported in AT series Mg alloys by Kabir et al. [25]. This may result from the depletion of Zn in the solid solution by DRXSP, which decreases the ratio of gSF/gUSF of basal slip systems (where gSF is stacking fault energy and gUSF is unstable stacking fault energy). The reduction of ratio of gSF/gUSF favors the deformation mechanism by partial dislocations [26]. The alloys extruded at higher temperatures exhibit the larger average grain size and more precipitate volume fraction simultaneously. The fact indicates the grain boundary migration is thermally activated obviously. 3.2. Textures The macro-textures of extruded alloys examined by XRD are shown in Fig. 8 in the terms of (0002) and ð1010Þ pole figures. The as-extruded alloys exhibit a fiber texture with (0002) planes parallel to the extrusion direction. It is the typical texture observed in
Mg alloy after extrusion [27]. However, the evident differences can be found in the distribution of pole intensity and maximum intensity. After cold extrusion, the material exhibits the weakest texture. Because the smallest critical resolved shear stress of dislocations slipping on the basal plane contributes to developing the deformation texture. In addition, the f1012g twinning can be easily operated during cold extrusion due to its low critical resolved shear stress. It gives rise to a lattice rotation of 86.3 , which contributes to the formation of a random texture. The strongest texture is obtained in sample extruded at 150 C with the maximum basal texture intensity of 8.032. Further increase in the extrusion temperature, the texture type has not changed, while the intensity of texture presents a downward trend. Non-basal slip systems can be easily operated in sample extruded above 250 C, which contributes to developing a weaker texture. In addition, the dynamic precipitates are perhaps another factor which leads to the basal texture weakening. According to the SAED patterns as shown in Fig. 7 (g) and (i), the occurrence of continuous rings suggests that these precipitates do not have any preferred orientation with respect to the a-Mg matrix. These nanoscale precipitates can help to pin or inhibit dislocations across the grains and lead to dislocation pile-up, which may act as preferential site for new grains nucleation and growth. In other words, these random distributions of precipitates may play an important role in increasing the
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Fig. 8. The (0002) and ð1010Þ pole figures of as-extruded alloys: (a) 25 C, (b) 150 C, (c) 250 C, (d) 300 C.
orientation between the new grains and parent grains. Thus, the basal texture is further improved. 3.3. Mechanical properties The mechanical properties obtained from tensile tests are shown in Fig. 9. The elongation to failure (d), ultimate tensile strength (UTS) and yield strength (YS) of the as-extruded alloys are summarized in Table 1. In general, the elongations of the as-
extruded alloys increase monotonically with the increasing of the extrusion temperature. The elongations of all samples herein exceed 20%. The lowest d of 21.2% is obtained in the alloy extruded at room temperature, while, the highest d of 36.7% is obtained in the sample extruded at 300 C. The excellent ductility may be associated with the work hardening behavior. The strain hardening curves of the samples using q(ss0.2) plot are shown in Fig. 9. (b), where q ¼ ds/dε, s and ε refer to strain hardening rate, true stress and true strain,
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Fig. 9. The mechanical properties for Mg-3.0Zn-0.2Ca alloys extruded at different temperatures: (a) engineering stressestrain curves, (b) stain-hardening curves using q(ss0.2) plot obtained from true stressestrain curves (as shown in inset).
Table 1 Tensile properties of the alloys extruded at different extrusion temperatures. Temperature/ C
YS/MPa
UTS/MPa
d/%
KⅢ
25 150 250 300
171.6 226.7 146.5 145.5
241.7 242.1 228 240.1
21.2 28.3 29.3 36.7
13.26 9.22 15.59 15.83
respectively. Obviously, these work hardening curves are different in shapes, which indicate that they are corresponding to different deformation mechanisms. A positive slope of strain hardening rate (dq/ds > 0) at the early stage of tensile deformation is observed in the samples extruded above 150 C (T 150 C), which means that twinning-dominated deformation dominates the early stage of strain hardening in tension. After a brief increase in strain hardening rate, a continuous decrease in strain hardening rate is observed in tension. This indicates that multiple slips dominate most parts of strain hardening in tension. The linear region is believed to be related to dynamic recovery (denoted as stage Ⅲ). The slope of stage Ⅲ denoted as KⅢ is used to characterize hardening behavior [28]. As shown, the slope of stage Ⅲ becomes steeper with increase of temperature except 150 C. The hardening behavior is usually associated with the texture, grain size and precipitates. It should be noted that the alloy extruded at room temperature shows a different hardening behavior at the early stage of tensile deformation. Furthermore, the twinning-dominated deformation region is not observed in the cold extrusion sample. That is, with decreasing grain size, twinning becomes increasingly difficult and the non-basal slip systems are easily activated. Both of them account for the early stage of tensile deformation in the cold extrusion sample. However, the highest dislocation density and relatively thick twins may deteriorate the ductility of the cold extrusion sample. Comparatively, the lowest slope of stage Ⅲ, KⅢ is obtained in sample extruded at 150 C, which may result from the strongest basal texture. This is not conducive to slip and needs the deformation twinning to accommodate the deformation. Further increasing the extrusion temperature, the larger grain size and dynamic precipitates are observed as mentioned above, both of which are beneficial to the dislocation accumulation and generate work hardening. Therefore, the ductility of the alloys extruded at higher temperatures is improved. Compared with the ductility, the strength of the as-extruded alloys increases firstly and subsequently decreases with the
increasing of extrusion temperature. The peaks in ultimate tensile strength (UTS) of 242.1 MPa and yield strength (YS) of 226.7 MPa are obtained in the sample extruded at 150 C. The lowest YS of 145.5 MPa is obtained in the sample extruded at 300 C, having the largest grain size of 20.6 mm. It is interesting that the cold extrusion sample with the smallest grain size of 8.1 mm shows the moderate UTS of 241.7 MPa and YS of 171.6 MPa. It seems that the yield stress of as-extruded alloys do against the HallePetch relation. As shown in Fig. 10, the red point A obviously deviates from the fitting of HallePetch relation in form of s0.2¼87.3 þ 312.8d1/2 (as shown in black line). The red point A corresponds to the yield strength of the alloy extruded at 150 C. Therefore, it can be concluded that HallePetch breakdown takes place in present study. Generally, the yield stress of the Mg alloys with hexagonal close packed structure was thought to depend on several factors: the grain size and its distribution, texture, solute atoms and precipitations. In the present case, the complete dynamic recrystallization occurs in all the as-extruded alloys and there is little difference in the relative grain size dispersion of all the as-extruded alloys. These indicate that the effects of residual strain and grain size distribution on yield stress of the Mg alloys are weak. In case of alloys extruded below 150 C, no fine precipitates (precipitates strengthening) appear in the as-extruded alloys. Thus, for the cold extrusion MgeZneCa alloy, the enhancement of YS is mainly attributed to fine grain strengthening (sgb) and solid solution strengthening (sss). The YS of the cold extrusion alloy is expressed as:
Fig. 10. Variations of strengths of as-extruded alloys as function of grain size.
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s0:2 ¼ sMg þ sgb þ sss
(1) sst ¼
where sMg is the yield strength of pure Mg matrix (sMg ¼ 21 MPa [29]), the influence of the grain size on YS can be estimated using HallePetch law [30]:
Dsgb ¼ kd1=2
(2)
where △sgb is the increase in yield stress due to grain refinement, k is the HallePetch coefficient (k z 312.8 MPa (mm)1/2 for MgeZneCa alloy as shown in Fig. 10), as shown in Fig. 6 (a), the average grain size d is 8.1 mm, so the △sgb is about 109.9 MPa. According to solid solution strengthening model of multicomponent alloys proposed by Gypen and Deruyttere [31], the △sss can be described as:
n 1=n 1=n Dsss ¼ kZn cZn þ kCa cCa
(3)
where n is a constant with 1/2 or 2/3, kZn and kCa are the strengthening constants for solutes Zn and Ca, respectively. The cZn and cCa are the concentration of solutes Zn and Ca in at.%, respectively. By the EDS analysis in Table 2, the cZn and cCa of cold extrusion MgeZneCa alloy are 1.15 at.% and 0.14 at.%, respectively. It is assumed that no interaction exists between Zn and Ca and n takes 2/3. The kZn ¼ 578 MPa (at.%)2/3 [29]. Consequently, the enhancement of YS from Zn solid solution strengthening is about 29.4 MPa and the calculative YS is 160.3 MPa. This result is 11.3 MPa lower than the experimental value of 171.6 MPa. The difference between theoretical calculation and experimental value originates from Ca solid solution strengthening. Based on the Equation (3), the strengthening constant for solute Ca is about kCa ¼ 243 MPa (at.%)2/3. At room temperature, the dominant slip system in Mg is basal slip, because the critical resolved shear stress for the basal slip system {0002} < 1120 > is much lower than that for non-basal systems. As shown in Fig. 11, the cold extrusion alloy and hot extrusion alloys exhibit almost similar average Schmid factor (SF) of ~0.20, while, a sharp basal texture is formed in the alloy extruded at 150 C with average Schmid factor of 0.17. Therefore, the texture is also an important factor which should be considered. Among these strengthening mechanisms, the texture and grain size may be the decisive factors affect the yield stress of warm extrusion alloy (150 C). The YS of the warm extrusion alloy can be expressed as:
s0:2 ¼ sMg þ sgb þ sss þ stex þ sother
(4)
where the sMg, sgb and sss have been defined previously. The stex and sother are the increase in yield stress due to texture strengthening and other factors. Based on the Equations (2) and (3), the contributions from sgb and sss to YS of the warm extrusion alloy are 89.2 MPa and 40.4 MPa, respectively. Consequently, the calculative YS is 150.6 MPa. This value is 76.1 MPa lower than the experimental value of 226.7 MPa, which originates from texture strengthening and other factors. Considering the texture strengthening effects, a revised form of HallePetch formula was proposed by Liu et al. [32]. In this formula, the combined effects of grain size and texture on the strength were defined as:
Table 2 The concentrations of solutes of the as-extruded alloys in at.%. Solutes
25 C/at.%
150 C/at.%
250 C/at.%
300 C/at.%
Zn Ca
1.15 0.14
1.13 0.13
0.76 0.13
0.68 0.13
0:3 Ds mt gb
129
(5)
where ss-t is the yield stress which originates from texture strengthening and grain size. mt is the average Schmid factor. As mentioned above, the cold and hot extrusion materials exhibit almost similar average Schmid factor (SF) of ~0.20. In other words, they have the same strengthening effects from the basal slip system. If the average SF of ~0.20 is assumed to be the average level of Schmid factor in the present alloys processed in the current conditions, thus the increment of YS from texture strengthening is assumed to be 0 MPa for the cold and hot extrusion materials. Comparatively, the alloy extruded at 150 C exhibits the lowest average SF factor of 0.17 as-mentioned above. According to the Equation (5), the contributions from texture strengthening to YS of the warm extrusion alloy is 23.6 MPa higher than those of other alloys. Compared to alloy with random orientation (SF ¼ 0.3) [32], the increment of YS from texture strengthening is about 68.2 MPa for the warm extrusion alloy. It should be emphasized that the texture strengthening is relative in this paper. The theoretical calculation indicates that texture strengthening makes important contribution to the increment of strength for warm extrusion alloy. In addition, Zeng et al. observed a remarkable strengthening effect in a tensile strained Mg-0.3Zn-0.1Ca (at.%) sheet alloy after annealing treatment at 80e200 C and they suggested that this phenomenon was caused by the pinning of gliding basal dislocations by solute atoms segregation [33]. As shown in inset of Fig. 7 (e), a similar morphology was observed in the alloy extruded at 150 C. The present alloy (Mg-1.13Zn-0.13Ca (at.%)) has similar chemical compositions as material used by Zeng et al. The extrusion temperature of 150 C is just within 80e200 C. After extrusion, the as-extruded alloy cooled in the air and the slow cooling rate contributed to solid atoms segregation. Furthermore, Zeng et al. suggested that the magnitude of the strength increment was about 8e22 MPa, approximately 9e24% higher than the yield stress of the initial test [33]. As for the warm extrusion alloy, all the factors except solid atoms segregation contribute to YS by about 174.2 MPa. Base on the experiment results, the solid atoms segregation may be another factor to enhance the strength and it may increase the strength increment by about 13.9e41.8 MPa. Comparatively, a large number of precipitates are generated in samples extruded above 250 C and the solute Zn decreases in aMg matrix. Thus, precipitation strengthening is another factor should be considered in samples extruded at higher temperatures. The YS of the hot extrusion (>250 C) alloys can be expressed as:
s0:2 ¼ sMg þ sgb þ sss þ sppt
(6)
where the sppt is the increase in yield stress due to precipitates strengthening. According to the same method, the contributions from sgb and sss to YS of the hot extrusion alloys are 70.8 MPa and 33.6 MPa for 250 C and 68.9 MPa and 32 MPa for 300 C. Consequently, the calculative YS (except precipitates strengthening) are 125.4 MPa for 250 C and 121.9 MPa for 300 C, respectively. These values are 21.1 MPa and 23.6 MPa lower than the experimental values of 146.5 and 145.5 MPa, respectively, which may originate from precipitates strengthening. In general, the quantitative relationship between the YS and precipitates can be estimated by Orowan mechanism. The Orowan equation is given as [34]:
sppt ¼
0:4MGb D pffiffiffiffiffiffiffiffiffiffiffi ln b pL 1 n
(7)
where M is the Taylor factor (M ¼ 5.6 [35]), G the shear modulus
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Fig. 11. Schmid factor for basal slip systems of the as-extruded alloys: (a) 25 C, (b) 150 C, (c) 250 C, (d) 300 C.
(G ¼ 16.6 GMPa [36]), b the Burgers vector (b ¼ 3.21 1010 m for pure Mg [36]), n the Poisson ratio (n ¼ 0.35 [36]), D the mean diameter of precipitate particles and L the effective inter-particle spacing. The corresponding values of D and L are 47.9 nm and 6.1 nm for 250 C and 48.3 nm and 6.5 nm for 300 C. Then the sppt are 336.4 MPa for 250 C and 340.8 MPa for 300 C, respectively. They are much higher than the estimate of 21.1 MPa (250 C) and 23.6 MPa (300 C), or even much higher than the actual yield strengths. It indicates Orowan equation is not suited to the current study. The reason is unclear and needs to be further investigated in future. Supposing no other strengthening mechanisms contributing to the yield strength, thus the estimate of 21.1 MPa (250 C) and 23.6 MPa (300 C) are attributed to precipitates strengthening. The contributions of different strengthening mechanisms to yield strength are summarized in Fig. 12. The theoretical calculation indicates that the fine grain strengthening makes the largest contribution to the enhancement of strength for the as-extruded alloys. The solid solution strengthening is the second contributor to the increment of strength. However, the precipitation strengthening only plays relatively minor role in the flow stress in hot extrusion alloys. The texture strengthening makes more contribution to the increment of YS for the alloy extruded at 150 C than the other alloys.
4. Conclusions The Mg-3.0Zn-0.2Ca (wt.%) alloy has been developed by casting followed by solution treatment and extrusion. Effects of extrusion temperature on the microstructure, texture evolution, mechanical properties and strengthening mechanisms of the as-extruded alloys are systematically studied. The main conclusions are summarized as follows:
Fig. 12. The contributions of different strengthening mechanisms to yield strengths of the as-extruded alloys (sgb: fine grain strengthening, sss: solid solution strengthening, sppt: precipitates strengthening and stex þ sother: texture strengthening and other factors.).
(1) After extrusion, the grain size is greatly refined and the average grain size increases from 8.1 mm to 20.6 mm with the increasing of the extrusion temperature. The smallest and uniform grain size are obtained in alloy extruded at room temperature. In addition, a large number of nano-scale MgZn2 precipitates are formed in alloys extruded above 250 C. (2) The deformation twinning plays an important role in accommodating the plastic deformation in the cold extrusion sample and leads to the basal texture weakening. The sharpest basal texture is obtained in sample extruded at 150 C. Further increase in the extrusion temperature, the intensity of texture presents a downward trend, which may
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be associated with multiple deformation mechanisms and dynamic precipitates. (3) The highest elongation of 36.7% is obtained in alloy extruded at 300 C，which results from the higher work hardening rate. Comparatively, the highest yield strength of 226.7 MPa occurs in alloy extruded at 150 C owing to the sharpest basal texture. (4) The fine grain strengthening and the solid solution strengthening play dominant roles in the increment of strength for the as-extruded alloys. However, the precipitation strengthening only makes relatively minor contribution to the flow stress in hot extrusion samples. Comparatively, the texture strengthening contributes more to the increment of YS for the alloy extruded at 150 C than the other alloys. References [1] C.J. Bettles, M.A. Gibson, Current wrought magnesium alloys: strengths and weaknesses, JOM 57 (2005) 46e49. [2] R.Z. Wu, Z.K. Qu, M.L. Zhang, Reviews on the influences of alloying elements on the microstructure and mechanical properties of MgeLi base alloys, Rev. Adv. Mater. Sci. 24 (2010) 35e43. [3] X. Gao, S.M. Zhu, B.C. Muddle, J.F. Nie, Precipitation-hardened MgeCaeZn alloys with superior creep resistance, Scr. Mater. 53 (2005) 1321e1326. [4] H. Yu, S.H. Park, B.S. You, Y.M. Kim, H.S. Yu, S.S. Park, Effects of extrusion speed on the microstructure and mechanical properties of ZK60 alloys with and without 1wt% cerium addition, Mater. Sci. Eng. A 583 (2013) 25e35. [5] S.H. Park, B.S. You, R.K. Mishra, A.K. Sachdev, Effects of extrusion parameters on the microstructure and mechanical properties of MgeZne(Mn)eCe/Gd alloys, Mater. Sci. Eng. A 598 (2014) 396e406. [6] X.S. Xia, Q. Chen, J.P. Li, D.Y. Shu, C.K. Hu, S.H. Huang, Z.D. Zhao, Characterization of hot deformation behavior of as-extruded MgeGdeYeZneZr alloy, J. Alloys. Comp. 610 (2014) 203e211. € bes, S. Zaefferer, I. Schenstakow, S. Yi, R. Gonzalez-Martinez, On the [7] S. Sandlo role of non-basal deformation mechanisms for the ductility of Mg and MgeY alloys, Acta. Mater. 59 (2011) 429e439. [8] T.L. Zhu, J.F. Sun, C.L. Cui, R.Z. Wu, S. Betsofen, Z. Leng, J.H. Zhang, M.L. Zhang, Influence of Y and Nd on microstructure, texture and anisotropy of Mge5Lie1Al alloy, Mater. Sci. Eng. A 600 (2014) 1e7. [9] X.R. Meng, R.Z. Wu, M.L. Zhang, L.B. Wu, C.L. Cui, Microstructures and properties of superlight MgeLieAleZn wrought alloys, J. Alloys. Comp. 486 (2009) 722e725. [10] J.H. Zhang, L. Zhang, Z. Leng, S.J. Liu, R.Z. Wu, M.L. Zhang, Experimental study on strengthening of MgeLi alloy by introducing long-period stacking ordered structure, Scr. Mater. 68 (2013) 675e678. [11] Y.W. Song, D.Y. Shan, R.S. Chen, E.H. Han, Corrosion characterization of Mge8Li alloy in NaCl solution, Corros. Sci. 51 (2009) 1087e1094. [12] S.W. Xu, K. Oh-ishi, H. Sunohara, S. Kamado, Extruded MgeZneCaeMn alloys with low yield anisotropy, Mater. Sci. Eng. A 558 (2012) 356e365. [13] Z.G. Xu, C. Smith, S. Chen, J. Sankar, Development and microstructural characterizations of MgeZneCa alloys for biomedical applications, Mater. Sci. Eng. B 176 (2011) 1660e1665. [14] J. Hofstetter, M. Becker, E. Martinelli, A.M. Weinberg, B. Mingler, H. Kilian, € ffler, High-Strength Low-Alloy (HSLA) S. Pogatscher, P.J. Uggowitzer, J.F. Lo MgeZneCa alloys with excellent biodegradation performance, JOM 66 (2014) 566e572.
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[15] B.P. Zhang, Y. Wang, L. Geng, C.X. Lu, Effects of calcium on texture and mechanical properties of hot-extruded MgeZneCa alloys, Mater. Sci. Eng. A 539 (2012) 56e60. [16] B.P. Zhang, L. Geng, L.J. Huang, X.X. Zhang, C.C. Dong, Enhanced mechanical properties in fine-grained Mg-1.0Zn-0.5Ca alloys prepared by extrusion at different temperatures, Scr. Mater. 63 (2010) 1024e1327. [17] L.B. Tong, M.Y. Zheng, H. Chang, X.S. Hu, K. Wu, S.W. Xu, S. Kamado, Y. Kojima, Microstructure and mechanical properties of MgeZneCa alloy processed by equal channel angular pressing, Mater. Sci. Eng. A 523 (2009) 289e294. [18] S. Berbeni, V. Favier, M. Berveiller, Micro-macro modelling of the effects of the grain size distribution on the plastic flow stress of heterogeneous materials, Comput. Mater. Sci. 39 (2007) 96e105. [19] Y. Chino, K. Kimura, M. Mabuchi, Twinning behavior and deformation mechanisms of extruded AZ31 Mg alloy, Mater. Sci. Eng. A 486 (2008) 481e488. [20] M. Jotoku, A. Yamamoto, H. Tsubakono, Plastic deformation of high purity magnesium at room temperature, J. Jpn. Inst. Light Met 56 (2006) 711e715. [21] L.B. Tong, M.Y. Zheng, S.W. Xu, X.S. Hu, K. Wu, S. Kamado, G.J. Wang, X.Y. Lv, Room-temperature compressive deformation behavior of MgeZneCa alloy processed by equal channel angular pressing, Mater. Sci. Eng. A 528 (2010) 672e679. [22] E. Ayman, U. Junko, K. Katsuyoshi, Application of rapid solidification powder metallurgy to the fabrication of high-strength, high-ductility MgeAleZneCaeLa alloy through hot extrusion, Acta. Mater. 59 (2011) 273e282. s, M. Maeso, P. Pe rez, P. Adeva, Effect of extrusion temperature on [23] G. Garce superplasticity of PM-WE54, Mater. Sci. Eng. A 462 (2007) 127e131. [24] Z.Y. Zhang, R. Yang, G. Chen, Y.T. Zhao, Y. Shao, Correlation between microstructure and tensile behavior in powder metallurgy ZK60 alloys, Mater. Lett. 89 (2012) 166e168. [25] A.S.H. Kabir, M. Sanjari, J. Su, I.H. Jung, S. Yue, Effect of strain-induced precipitation on dynamic recrystallization in MgeAleSn alloys, Mater. Sci. Eng. A 616 (2014) 252e259. [26] M. Muzyk, Z. Pakiela, K.J. Kurzydlowski, Generalized stacking fault energy in magnesium alloys: density functional theory calculations, Scr. Mater. 66 (2012) 219e222. [27] B. Kim, C.H. Park, H.S. Kim, B.S. You, S.S. Park, Grain refinement and improved tensile properties of Mg-3Al-1Zn alloy processed by low-temperature indirect extrusion, Scr. Mater. 76 (2014) 21e24. [28] U.F. Kocks, H. Mecking, Physics and phenomenology of strain hardening: the FCC case, Prog. Meter. Sci. 48 (2003) 171e273. [29] L. Gao, R.S. Chen, E.H. Han, Solid solution strengthening behaviors in binary MgeY single phase alloys, J. Alloys. Comp. 472 (2009) 234e240. [30] N.J. Petch, The cleavage strength polycrystals, J. Iron Steel Inst. 174 (1953) 25e28. [31] L.A. Gypen, A. Deruyttere, Multi-component intrinsic solid solution softening and hardening, J. Less. Common. Met. 56 (1977) 91e101. [32] D. Liu, Z.Y. Liu, E.D. Wang, Effect of rolling reduction on microstructure, texture, mechanical properties and mechanical anisotropy of AZ31 magnesium alloys, Mater. Sci. Eng. A 612 (2014) 208e213. [33] Z.R. Zeng, Y.M. Zhu, M.Z. Bian, S.W. Xu, C.H.J. Davies, N. Birbilis, J.F. Nie, Annealing strengthening in a dilute MgeZneCa sheet alloy, Scr. Mater. 107 (2015) 127e130. [34] I. Sabirov, M.Y. Murashkin, R.Z. Valiev, Nanostructured aluminium alloys produced by severe plastic deformation: New horizons in development, Mater. Sci. Eng. A 560 (2013) 1e24. [35] A. Akhtar, E. Teghtsoonian, Solid solution strengthening of magnesium single crystals-I alloying behaviour in basal slip, Acta Metall. 17 (1969) 1339e1349. [36] H.J. Frost, M.F. Ashby, Deformation-mechanism Maps, Pergamon Press, Oxford, UK, 1982.
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Journal of Alloys and Compounds journal homepage: http://www.elsevier.com/locate/jalcom
Microstructure, texture and mechanical properties of Mg-3.0Zn-0.2Ca alloys fabricated by extrusion at various temperatures Cheng-jie Li a, *, Hong-fei Sun a, Xue-wen Li b, Jun-ling Zhang a, Wen-bin Fang a, b, Ze-yi Tan a a b
School of Materials Science and Engineering, Harbin Institute of Technology, Harbin, 150001, China School of Materials Science and Engineering, Harbin University of Science And Technology, Harbin, 150080, China
a r t i c l e i n f o
a b s t r a c t
Article history: Received 3 July 2015 Received in revised form 25 August 2015 Accepted 26 August 2015 Available online 1 September 2015
The Mg-3.0Zn-0.2Ca (wt.%) alloy has been extruded at temperature range of 25e300 C and the resulting microstructure, texture and mechanical properties are systematically investigated. The results show that the grain size monotonically increases with the increasing of the extrusion temperature and the texture intensity increases firstly and decreases subsequently. In addition, a large number of nano-scale precipitates are formed in alloys extruded above 250 C. The weakest basal texture developed in the cold extrusion alloy is related to the deformation twinning, while, the combining effects of activation of multiple deformation mechanisms and dynamic precipitates contribute to developing the weaker basal texture in alloys extruded above 250 C. A sharp basal texture enhances the yield strength of alloy extruded at 150 C greatly at the cost of work hardening rate. The highest elongation is achieved in the alloy extruded at 300 C. The enhanced formability is considered to be associated with the dependence of the weaker basal texture and dynamic precipitates on the strain hardening behavior. Among the multiple strengthening mechanisms, the fine grain strengthening and the solid solution strengthening play a dominant role in the increment of strength for the as-extruded alloys. © 2015 Elsevier B.V. All rights reserved.
Keywords: Magnesium alloy Extrusion Tensile properties Strengthening mechanisms Ductility
1. Introduction The application of wrought magnesium alloy is still limited due to their poor ductility and formability at room temperature [1]. Great deals of efforts have been devoted to improving the roomtemperature ductility of magnesium. Element alloying is an effective method to improve mechanical performance of magnesium alloys in nature, while thermo-mechanical process can enhance the formability of magnesium by grain refinement from the technical point of view. Literature have been documented about the obvious enhancement in mechanical properties of magnesium alloys by element alloying [2,3]. Moreover, the addition of rare-earth (RE) alloying elements (such as Gd, Ce, Y, etc.) to magnesium is believed to be one of the potential ways to enhance the formability of magnesium [4e6]. More random textures can be developed in magnesium alloys containing RE elements during hot-working [7]. However, large amount of expensive alloying elements will
* Corresponding author. E-mail address: [email protected] (C.-j. Li). http://dx.doi.org/10.1016/j.jallcom.2015.08.215 0925-8388/© 2015 Elsevier B.V. All rights reserved.
inevitably increase the materials cost, which is undesirable for industrial application. Another popular method is alloying with element of Li, which decreases the c/a axial ratio of lattice structure and makes non-basal slip systems easily activated at room temperature [8]. Therefore, the high elongation can be obtained in the magnesium-lithium alloys at room temperature [9]. However, the low strength and poor corrosion resistance severely limit their application as the structural components [10,11]. During the past several years, many researchers have devoted to developing a low-cost MgeZneCa alloy for structural and medical applications [12e14]. Such magnesium alloy has high ductility accompanied with high strength and good biocompatibility. In fact, thermo-mechanical processes have been adoped to refine the grain sizes in such magnesium alloys [15,16]. Grain refinement is desirable for their strength and ductility. In addition, a weak basal fiber texture can be developed in hot-extruded MgeZneCa alloy due to the dynamic recrystallization. Both the weakened basal fiber textures and refined grains are favorable to increase roomtemperature ductility of magnesium [15]. At present, many researchers focus on optimizing the extrusion parameters (such as extrusion temperature, ram speed and extrusion ratio) to obtain
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fine grains and random textures [12e17]. Among these variables, extrusion temperature is the most critical parameter that directly affects microstructure, texture evolution and mechanical properties of MgeZneCa alloy. Zhang et al. reported that the excellent mechanical properties in Mg-1.0Zn-0.5Ca alloy can be obtained by hot extrusion because of the grain refinement and the weakened basal texture [15]. They have also proved that the grain refinement method together with the micro-alloying concept was a very effective route for fabricating a high performance MgeZneCa alloy. To date, most of the previous works are concentrated on the hot extrusion behaviors of MgeZneCa alloys. That is, most of these alloys are extruded above 250 C. But, the systematic investigation on the low-temperature (<200 C) extrusion behaviors of MgeZneCa alloys has rarely been reported. Thus, the aim of this study is to investigate the relationship between extrusion temperatures (25e300 C) and microstructure as well as mechanical performance of MgeZneCa alloy. 2. Experimental procedure The as-cast Mg-3.0Zn-0.2Ca (wt.%) alloy used in the experiment was synthesized with pure Mg (99.95 wt.%), Zn powder (99.7 wt.%) and Mg-30 wt.%Ca master alloys. Under protection with a mixed gas atmosphere of CO2 and SF6 (CO2:SF6 ¼ 40:1 vol.%), the Zn powder and MgeCa master alloys were added into the melted pure Mg at 750 C. The melted alloys were isothermally held at 750 C for 30 min to ensure that all the alloying elements dissolved. After that, the melted alloy was poured into a stainless steel mold preheated to 200 C and cooled in the air. The as-cast ingots were solution-treated at 400 C for 10 h, followed by water quenching. The prepared billets were machined into cylindrical bar with the diameter of 52 mm and height of 50 mm for extrusion. Then, the prepared billets were extruded at different temperatures of 25 C, 150 C, 250 C and 300 C. The extrusion ram speed was set as 17 mm/s and the extrusion ratio was 8.3. Finally, the as-extruded Mg alloy bars with a diameter of 18 mm were obtained. The microstructure of as-extruded alloys was observed in an Olympus GX71 optical microscope (OM), and a JEM-2000EX transmission electron microscope (TEM). The grain size was measured using Image-ProPlus 6.0 software. The macro-texture was measured by the Schulz reflection method using X'pert PRO XRD. All the samples for XRD test were cut from center of asextruded alloys and sectioned in parallel to the extrusion direction (ED). To check the repeatability of the results, two experiments were conducted under each set of conditions. The micro-texture was examined by electron backscattering diffraction (EBSD) in Quanta 200FEG emission scanning electron microscope (SEM) equipped with OIM Analysis 5.0 software. The phase identification and analysis were performed on a D/max-rb X-ray diffractometer (XRD) using monochromatic Cu Ka radiation. The tensile tests were carried out by an Instron-5569 standard testing machine at an initial strain rate of 1.67 103 s1 at room temperature. The dogbone tensile specimens with a diameter of 3 mm and gage length of 15 mm were used. All the samples were sectioned in parallel to the extrusion direction (ED). To check the repeatability of the results, three experiments were conducted under each set of conditions. 3. Results and discussion 3.1. Microstructure Fig. 1 shows the microstructures of Mg-3.0Zn-0.2Ca alloys before and after solution treatment. As seen from Fig. 1 (a), the microstructure of as-cast alloy consists of a-Mg matrix and eutectic compounds (marked by white arrows). Most of a-Mg present
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equiaxed grains with an average grain size of 449 mm. After solution treatment, a single phase microstructure can be obtained as shown in Fig. 1 (b), because the net-like interdendritic eutectic structures almost disappear and dissolve into the matrix. The average grain size of solution treated Mg-3.0Zn-0.2Ca alloy is about 453 mm. There is little difference in grain size of the alloys before and after solution treatment. From the SEM micrograph (Fig. 2 (a)), it can be seen clearly that the second phases (in white) exhibit continuous network structure and mainly distribute along grain boundaries and inside the grains. After homogenizing annealing at 400 C for 10 h, no intermetallics can be observed in Fig. 2 (b). This result is in agreement with the result observed in Fig. 1 (b). The result from XRD shows that the as-cast Mg-3.0Zn-0.2Ca alloy contains three phases, a-Mg, MgZn2 and Ca2Mg5Zn13 (Fig. 3 (a)). After solution treatment, only a-Mg can be detected (Fig. 3 (b)). In view of microstructure (Fig. 2 (a)), it can be expected that the primary dendrites should be a-Mg and the white phases may be the MgZn2 and Ca2Mg5Zn13. The EDS analysis indicates that the average Zn/Ca atomic ratio of the white phases (shown by red crosses A, B and C in Fig. 2 (a)) is about 15.8:1 (Fig. 4). The Zn/Ca atomic ratio of Ca2Mg5Zn13 phase is 6.5:1, which is far less than 15.8:1. That is, the white phases contain abundant zinc element. Thus, it can be concluded that the ternary MgeZneCa phases are MgZn2 and Ca2Mg5Zn13 mixed phases. Fig. 5 shows the microstructural characteristics of as-extruded alloys at different extrusion temperatures. Compared with the microstructure of as-cast billet, the grain size of as-extruded alloys is refined remarkably. The as-extruded alloys exhibit uniform and equiaxed grains. It indicates that a complete dynamic recrystallization process can be obtained at all three above-mentioned extrusion temperatures. As the deformation temperature increases from 25 C to 250 C, the average grain size increases from 8.1 mm to 19.5 mm. However, as the deformation temperature further rises to 300 C, the average grain size only increases from 19.5 mm to 20.6 mm. In other words, the increase of grain size is not obvious. Therefore, it can be concluded that the microstructure is sensitive to deformation temperature when the samples are extruded at the lower temperatures (25e250 C). The extrusion temperature not only affects the grain size but also influences the grain size distribution. In the present study, at least over six hundred grains from optical micrographs are analyzed for each sample obtained at different extrusion temperatures. Fig. 6 shows the area-weighted grain size distributions, the average grain size (dm) and the relative grain size dispersion △d/dm of the as-extruded samples. Regarding the microstructural heterogeneity, the relative grain size dispersion △d/dm is used to characterize grain size dispersions according to reference [18], where d is the grain size, △d is the absolute grain size range (△d ¼ dmaxdmin) and dm is the mean grain size. All of the deformed materials show the approximate normal distribution with a single peak in spite of the difference in the peak values. In addition, there is a slight difference in relative grain size dispersion. The most well-distributed grain size is obtained in the cold extruded sample with the smallest △d/dm of 2.47. Further increase in the deformation temperature introduces some grains uneven growth, which leads to microstructure heterogeneity. However, in view of microstructure (Fig. 5) and statistical result (Fig. 6), the △d/ dm is only a slight increase of ~0.2 compared with that of cold extrusion when the extrusions are carried out at the higher temperature range of 150e300 C. It is evident that the extrusion temperature has a less influence on the grain size dispersion than average grain size. To further analyze the microstructure of extruded specimens, TEM tests are performed. Fig. 7 shows TEM images of the alloys extruded at different temperatures. As shown in Fig. 7 (a)e(e), the
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Fig. 1. Optical micrographs of the Mg-3.0Zn-0.2Ca alloy: (a) as-cast; (b) solution treated.
Fig. 2. SEM micrographs of the Mg-3.0Zn-0.2Ca alloy: (a) as-cast; (b) solution treated.
exquiaxed grains with few dislocations recognized as recrystallized structure are observed in alloys extruded below 150 C. It suggests that the dynamic recrystallization occurs during the extrusion process even at room temperature. Such a phenomenon has been observed in AZ31 alloy, high purity Mg and Mg-5.25wt.%Zn-0.6wt.% Ca after severe plastic deformation at room temperature [19e21]. Jotoku et al. suggested that the cross-slip and annihilation of dislocation resulted from the stress concentration at local area due to the pile-up of dislocations, which might cause the recrystallization [20]. In the present study, some dislocations are accumulated at twin boundaries and in the original grains, both of which
Fig. 3. X-ray diffraction patterns of Mg-3.0Zn-0.2Ca alloy: (a) as-cast; (b) solution treated.
are benificial to form some dislocation cells (Fig. 7 (c) and (d)). With increasing plastic deformation, the misorientation between the cells is increased and so the rate of growth of the sub-grain increases. Furthermore, this process results in continuous dynamic recrystallization in some area with high stress concentration. In addition, this process can be accelerated by heating generated from plastic deformation. As indicated by white arrows in Fig. 7 (c) and (d), some submicron twins are also observed, which indicates that deformation twinning plays an important role in the magnesium alloys extruded at lower temperatures (150 C). Because magnesium alloys lack sufficient independent slip systems at lower temperature. Further increase in the extrusion temperature
Fig. 4. EDS analysis of the interdendritic phases marked by red crosses A, B and C in Fig. 2 (a). (For interpretation of the references to colour in this figure legend, the reader is referred to the web version of this article.)
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Fig. 5. Optical micrographs of the as-extruded alloys: (a) 25 C, (b) 150 C, (c) 250 C, (d) 300 C.
(250 C), more slip systems can be activated. Thus, only a few deformation twins with size less than 250 nm in width can be observed in Fig. 7 (f) and (h). As shown in Fig. 7, the extrusion temperature has profound influence on second phases. Very few second phases can be observed in samples extruded below 150 C (as shown in Fig. 7(aee)). Comparatively, a large number of precipitates with
diameter of 5e30 nm dispersed within the grains are observed in samples extruded above 250 C. These second phases formed during the extrusion process are confirmed to be MgZn2 by TEM diffraction patterns. A similar phenomenon has been reported in other Mg alloys at higher extrusion temperatures (300 C) [22e24]. Generally, the high-energy sites (such as dislocations) easily are generated by hot plastic deformation, which are
Fig. 6. Grain size distributions of the as-extruded alloys: (a) 25 C, (b) 150 C, (c) 250 C, (d) 300 C.
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Fig. 7. The TEM micrographs of the as-extruded alloys: (a) the bright field image of alloy extruded at 25 C and (b) corresponding dark field image of (a), (c) bright field image of deformation twins in alloy extruded at 25 C, (d) bright field image of alloy extruded at 150 C and (e) corresponding dark field image of (d), (f) bright field image of alloy extruded at 250 C and (g) corresponding dark field image of (f), (h) bright field image of alloy extruded at 300 C and (i) corresponding dark field image of (h).
conducive to the formation of fine precipitates. These fine precipitates in turn exert a pinning effect on the movement of grain boundaries and suppress dynamic recystallization. However, combining optical and TEM micrographs, the dynamic precipitates are formed in the dynamic recystallized regions. That is, the fine precipitates cannot effectively suppress dynamic recystallization. Comparatively, the dynamic recystallization may stimulate dynamic precipitation (DRXSP). A similar phenomenon of DRXSP has been reported in AT series Mg alloys by Kabir et al. [25]. This may result from the depletion of Zn in the solid solution by DRXSP, which decreases the ratio of gSF/gUSF of basal slip systems (where gSF is stacking fault energy and gUSF is unstable stacking fault energy). The reduction of ratio of gSF/gUSF favors the deformation mechanism by partial dislocations [26]. The alloys extruded at higher temperatures exhibit the larger average grain size and more precipitate volume fraction simultaneously. The fact indicates the grain boundary migration is thermally activated obviously. 3.2. Textures The macro-textures of extruded alloys examined by XRD are shown in Fig. 8 in the terms of (0002) and ð1010Þ pole figures. The as-extruded alloys exhibit a fiber texture with (0002) planes parallel to the extrusion direction. It is the typical texture observed in
Mg alloy after extrusion [27]. However, the evident differences can be found in the distribution of pole intensity and maximum intensity. After cold extrusion, the material exhibits the weakest texture. Because the smallest critical resolved shear stress of dislocations slipping on the basal plane contributes to developing the deformation texture. In addition, the f1012g twinning can be easily operated during cold extrusion due to its low critical resolved shear stress. It gives rise to a lattice rotation of 86.3 , which contributes to the formation of a random texture. The strongest texture is obtained in sample extruded at 150 C with the maximum basal texture intensity of 8.032. Further increase in the extrusion temperature, the texture type has not changed, while the intensity of texture presents a downward trend. Non-basal slip systems can be easily operated in sample extruded above 250 C, which contributes to developing a weaker texture. In addition, the dynamic precipitates are perhaps another factor which leads to the basal texture weakening. According to the SAED patterns as shown in Fig. 7 (g) and (i), the occurrence of continuous rings suggests that these precipitates do not have any preferred orientation with respect to the a-Mg matrix. These nanoscale precipitates can help to pin or inhibit dislocations across the grains and lead to dislocation pile-up, which may act as preferential site for new grains nucleation and growth. In other words, these random distributions of precipitates may play an important role in increasing the
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Fig. 8. The (0002) and ð1010Þ pole figures of as-extruded alloys: (a) 25 C, (b) 150 C, (c) 250 C, (d) 300 C.
orientation between the new grains and parent grains. Thus, the basal texture is further improved. 3.3. Mechanical properties The mechanical properties obtained from tensile tests are shown in Fig. 9. The elongation to failure (d), ultimate tensile strength (UTS) and yield strength (YS) of the as-extruded alloys are summarized in Table 1. In general, the elongations of the as-
extruded alloys increase monotonically with the increasing of the extrusion temperature. The elongations of all samples herein exceed 20%. The lowest d of 21.2% is obtained in the alloy extruded at room temperature, while, the highest d of 36.7% is obtained in the sample extruded at 300 C. The excellent ductility may be associated with the work hardening behavior. The strain hardening curves of the samples using q(ss0.2) plot are shown in Fig. 9. (b), where q ¼ ds/dε, s and ε refer to strain hardening rate, true stress and true strain,
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Fig. 9. The mechanical properties for Mg-3.0Zn-0.2Ca alloys extruded at different temperatures: (a) engineering stressestrain curves, (b) stain-hardening curves using q(ss0.2) plot obtained from true stressestrain curves (as shown in inset).
Table 1 Tensile properties of the alloys extruded at different extrusion temperatures. Temperature/ C
YS/MPa
UTS/MPa
d/%
KⅢ
25 150 250 300
171.6 226.7 146.5 145.5
241.7 242.1 228 240.1
21.2 28.3 29.3 36.7
13.26 9.22 15.59 15.83
respectively. Obviously, these work hardening curves are different in shapes, which indicate that they are corresponding to different deformation mechanisms. A positive slope of strain hardening rate (dq/ds > 0) at the early stage of tensile deformation is observed in the samples extruded above 150 C (T 150 C), which means that twinning-dominated deformation dominates the early stage of strain hardening in tension. After a brief increase in strain hardening rate, a continuous decrease in strain hardening rate is observed in tension. This indicates that multiple slips dominate most parts of strain hardening in tension. The linear region is believed to be related to dynamic recovery (denoted as stage Ⅲ). The slope of stage Ⅲ denoted as KⅢ is used to characterize hardening behavior [28]. As shown, the slope of stage Ⅲ becomes steeper with increase of temperature except 150 C. The hardening behavior is usually associated with the texture, grain size and precipitates. It should be noted that the alloy extruded at room temperature shows a different hardening behavior at the early stage of tensile deformation. Furthermore, the twinning-dominated deformation region is not observed in the cold extrusion sample. That is, with decreasing grain size, twinning becomes increasingly difficult and the non-basal slip systems are easily activated. Both of them account for the early stage of tensile deformation in the cold extrusion sample. However, the highest dislocation density and relatively thick twins may deteriorate the ductility of the cold extrusion sample. Comparatively, the lowest slope of stage Ⅲ, KⅢ is obtained in sample extruded at 150 C, which may result from the strongest basal texture. This is not conducive to slip and needs the deformation twinning to accommodate the deformation. Further increasing the extrusion temperature, the larger grain size and dynamic precipitates are observed as mentioned above, both of which are beneficial to the dislocation accumulation and generate work hardening. Therefore, the ductility of the alloys extruded at higher temperatures is improved. Compared with the ductility, the strength of the as-extruded alloys increases firstly and subsequently decreases with the
increasing of extrusion temperature. The peaks in ultimate tensile strength (UTS) of 242.1 MPa and yield strength (YS) of 226.7 MPa are obtained in the sample extruded at 150 C. The lowest YS of 145.5 MPa is obtained in the sample extruded at 300 C, having the largest grain size of 20.6 mm. It is interesting that the cold extrusion sample with the smallest grain size of 8.1 mm shows the moderate UTS of 241.7 MPa and YS of 171.6 MPa. It seems that the yield stress of as-extruded alloys do against the HallePetch relation. As shown in Fig. 10, the red point A obviously deviates from the fitting of HallePetch relation in form of s0.2¼87.3 þ 312.8d1/2 (as shown in black line). The red point A corresponds to the yield strength of the alloy extruded at 150 C. Therefore, it can be concluded that HallePetch breakdown takes place in present study. Generally, the yield stress of the Mg alloys with hexagonal close packed structure was thought to depend on several factors: the grain size and its distribution, texture, solute atoms and precipitations. In the present case, the complete dynamic recrystallization occurs in all the as-extruded alloys and there is little difference in the relative grain size dispersion of all the as-extruded alloys. These indicate that the effects of residual strain and grain size distribution on yield stress of the Mg alloys are weak. In case of alloys extruded below 150 C, no fine precipitates (precipitates strengthening) appear in the as-extruded alloys. Thus, for the cold extrusion MgeZneCa alloy, the enhancement of YS is mainly attributed to fine grain strengthening (sgb) and solid solution strengthening (sss). The YS of the cold extrusion alloy is expressed as:
Fig. 10. Variations of strengths of as-extruded alloys as function of grain size.
C.-j. Li et al. / Journal of Alloys and Compounds 652 (2015) 122e131
s0:2 ¼ sMg þ sgb þ sss
(1) sst ¼
where sMg is the yield strength of pure Mg matrix (sMg ¼ 21 MPa [29]), the influence of the grain size on YS can be estimated using HallePetch law [30]:
Dsgb ¼ kd1=2
(2)
where △sgb is the increase in yield stress due to grain refinement, k is the HallePetch coefficient (k z 312.8 MPa (mm)1/2 for MgeZneCa alloy as shown in Fig. 10), as shown in Fig. 6 (a), the average grain size d is 8.1 mm, so the △sgb is about 109.9 MPa. According to solid solution strengthening model of multicomponent alloys proposed by Gypen and Deruyttere [31], the △sss can be described as:
n 1=n 1=n Dsss ¼ kZn cZn þ kCa cCa
(3)
where n is a constant with 1/2 or 2/3, kZn and kCa are the strengthening constants for solutes Zn and Ca, respectively. The cZn and cCa are the concentration of solutes Zn and Ca in at.%, respectively. By the EDS analysis in Table 2, the cZn and cCa of cold extrusion MgeZneCa alloy are 1.15 at.% and 0.14 at.%, respectively. It is assumed that no interaction exists between Zn and Ca and n takes 2/3. The kZn ¼ 578 MPa (at.%)2/3 [29]. Consequently, the enhancement of YS from Zn solid solution strengthening is about 29.4 MPa and the calculative YS is 160.3 MPa. This result is 11.3 MPa lower than the experimental value of 171.6 MPa. The difference between theoretical calculation and experimental value originates from Ca solid solution strengthening. Based on the Equation (3), the strengthening constant for solute Ca is about kCa ¼ 243 MPa (at.%)2/3. At room temperature, the dominant slip system in Mg is basal slip, because the critical resolved shear stress for the basal slip system {0002} < 1120 > is much lower than that for non-basal systems. As shown in Fig. 11, the cold extrusion alloy and hot extrusion alloys exhibit almost similar average Schmid factor (SF) of ~0.20, while, a sharp basal texture is formed in the alloy extruded at 150 C with average Schmid factor of 0.17. Therefore, the texture is also an important factor which should be considered. Among these strengthening mechanisms, the texture and grain size may be the decisive factors affect the yield stress of warm extrusion alloy (150 C). The YS of the warm extrusion alloy can be expressed as:
s0:2 ¼ sMg þ sgb þ sss þ stex þ sother
(4)
where the sMg, sgb and sss have been defined previously. The stex and sother are the increase in yield stress due to texture strengthening and other factors. Based on the Equations (2) and (3), the contributions from sgb and sss to YS of the warm extrusion alloy are 89.2 MPa and 40.4 MPa, respectively. Consequently, the calculative YS is 150.6 MPa. This value is 76.1 MPa lower than the experimental value of 226.7 MPa, which originates from texture strengthening and other factors. Considering the texture strengthening effects, a revised form of HallePetch formula was proposed by Liu et al. [32]. In this formula, the combined effects of grain size and texture on the strength were defined as:
Table 2 The concentrations of solutes of the as-extruded alloys in at.%. Solutes
25 C/at.%
150 C/at.%
250 C/at.%
300 C/at.%
Zn Ca
1.15 0.14
1.13 0.13
0.76 0.13
0.68 0.13
0:3 Ds mt gb
129
(5)
where ss-t is the yield stress which originates from texture strengthening and grain size. mt is the average Schmid factor. As mentioned above, the cold and hot extrusion materials exhibit almost similar average Schmid factor (SF) of ~0.20. In other words, they have the same strengthening effects from the basal slip system. If the average SF of ~0.20 is assumed to be the average level of Schmid factor in the present alloys processed in the current conditions, thus the increment of YS from texture strengthening is assumed to be 0 MPa for the cold and hot extrusion materials. Comparatively, the alloy extruded at 150 C exhibits the lowest average SF factor of 0.17 as-mentioned above. According to the Equation (5), the contributions from texture strengthening to YS of the warm extrusion alloy is 23.6 MPa higher than those of other alloys. Compared to alloy with random orientation (SF ¼ 0.3) [32], the increment of YS from texture strengthening is about 68.2 MPa for the warm extrusion alloy. It should be emphasized that the texture strengthening is relative in this paper. The theoretical calculation indicates that texture strengthening makes important contribution to the increment of strength for warm extrusion alloy. In addition, Zeng et al. observed a remarkable strengthening effect in a tensile strained Mg-0.3Zn-0.1Ca (at.%) sheet alloy after annealing treatment at 80e200 C and they suggested that this phenomenon was caused by the pinning of gliding basal dislocations by solute atoms segregation [33]. As shown in inset of Fig. 7 (e), a similar morphology was observed in the alloy extruded at 150 C. The present alloy (Mg-1.13Zn-0.13Ca (at.%)) has similar chemical compositions as material used by Zeng et al. The extrusion temperature of 150 C is just within 80e200 C. After extrusion, the as-extruded alloy cooled in the air and the slow cooling rate contributed to solid atoms segregation. Furthermore, Zeng et al. suggested that the magnitude of the strength increment was about 8e22 MPa, approximately 9e24% higher than the yield stress of the initial test [33]. As for the warm extrusion alloy, all the factors except solid atoms segregation contribute to YS by about 174.2 MPa. Base on the experiment results, the solid atoms segregation may be another factor to enhance the strength and it may increase the strength increment by about 13.9e41.8 MPa. Comparatively, a large number of precipitates are generated in samples extruded above 250 C and the solute Zn decreases in aMg matrix. Thus, precipitation strengthening is another factor should be considered in samples extruded at higher temperatures. The YS of the hot extrusion (>250 C) alloys can be expressed as:
s0:2 ¼ sMg þ sgb þ sss þ sppt
(6)
where the sppt is the increase in yield stress due to precipitates strengthening. According to the same method, the contributions from sgb and sss to YS of the hot extrusion alloys are 70.8 MPa and 33.6 MPa for 250 C and 68.9 MPa and 32 MPa for 300 C. Consequently, the calculative YS (except precipitates strengthening) are 125.4 MPa for 250 C and 121.9 MPa for 300 C, respectively. These values are 21.1 MPa and 23.6 MPa lower than the experimental values of 146.5 and 145.5 MPa, respectively, which may originate from precipitates strengthening. In general, the quantitative relationship between the YS and precipitates can be estimated by Orowan mechanism. The Orowan equation is given as [34]:
sppt ¼
0:4MGb D pffiffiffiffiffiffiffiffiffiffiffi ln b pL 1 n
(7)
where M is the Taylor factor (M ¼ 5.6 [35]), G the shear modulus
130
C.-j. Li et al. / Journal of Alloys and Compounds 652 (2015) 122e131
Fig. 11. Schmid factor for basal slip systems of the as-extruded alloys: (a) 25 C, (b) 150 C, (c) 250 C, (d) 300 C.
(G ¼ 16.6 GMPa [36]), b the Burgers vector (b ¼ 3.21 1010 m for pure Mg [36]), n the Poisson ratio (n ¼ 0.35 [36]), D the mean diameter of precipitate particles and L the effective inter-particle spacing. The corresponding values of D and L are 47.9 nm and 6.1 nm for 250 C and 48.3 nm and 6.5 nm for 300 C. Then the sppt are 336.4 MPa for 250 C and 340.8 MPa for 300 C, respectively. They are much higher than the estimate of 21.1 MPa (250 C) and 23.6 MPa (300 C), or even much higher than the actual yield strengths. It indicates Orowan equation is not suited to the current study. The reason is unclear and needs to be further investigated in future. Supposing no other strengthening mechanisms contributing to the yield strength, thus the estimate of 21.1 MPa (250 C) and 23.6 MPa (300 C) are attributed to precipitates strengthening. The contributions of different strengthening mechanisms to yield strength are summarized in Fig. 12. The theoretical calculation indicates that the fine grain strengthening makes the largest contribution to the enhancement of strength for the as-extruded alloys. The solid solution strengthening is the second contributor to the increment of strength. However, the precipitation strengthening only plays relatively minor role in the flow stress in hot extrusion alloys. The texture strengthening makes more contribution to the increment of YS for the alloy extruded at 150 C than the other alloys.
4. Conclusions The Mg-3.0Zn-0.2Ca (wt.%) alloy has been developed by casting followed by solution treatment and extrusion. Effects of extrusion temperature on the microstructure, texture evolution, mechanical properties and strengthening mechanisms of the as-extruded alloys are systematically studied. The main conclusions are summarized as follows:
Fig. 12. The contributions of different strengthening mechanisms to yield strengths of the as-extruded alloys (sgb: fine grain strengthening, sss: solid solution strengthening, sppt: precipitates strengthening and stex þ sother: texture strengthening and other factors.).
(1) After extrusion, the grain size is greatly refined and the average grain size increases from 8.1 mm to 20.6 mm with the increasing of the extrusion temperature. The smallest and uniform grain size are obtained in alloy extruded at room temperature. In addition, a large number of nano-scale MgZn2 precipitates are formed in alloys extruded above 250 C. (2) The deformation twinning plays an important role in accommodating the plastic deformation in the cold extrusion sample and leads to the basal texture weakening. The sharpest basal texture is obtained in sample extruded at 150 C. Further increase in the extrusion temperature, the intensity of texture presents a downward trend, which may
C.-j. Li et al. / Journal of Alloys and Compounds 652 (2015) 122e131
be associated with multiple deformation mechanisms and dynamic precipitates. (3) The highest elongation of 36.7% is obtained in alloy extruded at 300 C，which results from the higher work hardening rate. Comparatively, the highest yield strength of 226.7 MPa occurs in alloy extruded at 150 C owing to the sharpest basal texture. (4) The fine grain strengthening and the solid solution strengthening play dominant roles in the increment of strength for the as-extruded alloys. However, the precipitation strengthening only makes relatively minor contribution to the flow stress in hot extrusion samples. Comparatively, the texture strengthening contributes more to the increment of YS for the alloy extruded at 150 C than the other alloys. References [1] C.J. Bettles, M.A. Gibson, Current wrought magnesium alloys: strengths and weaknesses, JOM 57 (2005) 46e49. [2] R.Z. Wu, Z.K. Qu, M.L. Zhang, Reviews on the influences of alloying elements on the microstructure and mechanical properties of MgeLi base alloys, Rev. Adv. Mater. Sci. 24 (2010) 35e43. [3] X. Gao, S.M. Zhu, B.C. Muddle, J.F. Nie, Precipitation-hardened MgeCaeZn alloys with superior creep resistance, Scr. Mater. 53 (2005) 1321e1326. [4] H. Yu, S.H. Park, B.S. You, Y.M. Kim, H.S. Yu, S.S. Park, Effects of extrusion speed on the microstructure and mechanical properties of ZK60 alloys with and without 1wt% cerium addition, Mater. Sci. Eng. A 583 (2013) 25e35. [5] S.H. Park, B.S. You, R.K. Mishra, A.K. Sachdev, Effects of extrusion parameters on the microstructure and mechanical properties of MgeZne(Mn)eCe/Gd alloys, Mater. Sci. Eng. A 598 (2014) 396e406. [6] X.S. Xia, Q. Chen, J.P. Li, D.Y. Shu, C.K. Hu, S.H. Huang, Z.D. Zhao, Characterization of hot deformation behavior of as-extruded MgeGdeYeZneZr alloy, J. Alloys. Comp. 610 (2014) 203e211. € bes, S. Zaefferer, I. Schenstakow, S. Yi, R. Gonzalez-Martinez, On the [7] S. Sandlo role of non-basal deformation mechanisms for the ductility of Mg and MgeY alloys, Acta. Mater. 59 (2011) 429e439. [8] T.L. Zhu, J.F. Sun, C.L. Cui, R.Z. Wu, S. Betsofen, Z. Leng, J.H. Zhang, M.L. Zhang, Influence of Y and Nd on microstructure, texture and anisotropy of Mge5Lie1Al alloy, Mater. Sci. Eng. A 600 (2014) 1e7. [9] X.R. Meng, R.Z. Wu, M.L. Zhang, L.B. Wu, C.L. Cui, Microstructures and properties of superlight MgeLieAleZn wrought alloys, J. Alloys. Comp. 486 (2009) 722e725. [10] J.H. Zhang, L. Zhang, Z. Leng, S.J. Liu, R.Z. Wu, M.L. Zhang, Experimental study on strengthening of MgeLi alloy by introducing long-period stacking ordered structure, Scr. Mater. 68 (2013) 675e678. [11] Y.W. Song, D.Y. Shan, R.S. Chen, E.H. Han, Corrosion characterization of Mge8Li alloy in NaCl solution, Corros. Sci. 51 (2009) 1087e1094. [12] S.W. Xu, K. Oh-ishi, H. Sunohara, S. Kamado, Extruded MgeZneCaeMn alloys with low yield anisotropy, Mater. Sci. Eng. A 558 (2012) 356e365. [13] Z.G. Xu, C. Smith, S. Chen, J. Sankar, Development and microstructural characterizations of MgeZneCa alloys for biomedical applications, Mater. Sci. Eng. B 176 (2011) 1660e1665. [14] J. Hofstetter, M. Becker, E. Martinelli, A.M. Weinberg, B. Mingler, H. Kilian, € ffler, High-Strength Low-Alloy (HSLA) S. Pogatscher, P.J. Uggowitzer, J.F. Lo MgeZneCa alloys with excellent biodegradation performance, JOM 66 (2014) 566e572.
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