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Achieving high strength in indirectly-extruded binary Mg–Ca alloy containing Guinier–Preston zones
Journal of Alloys and Compounds 630 (2015) 272–276
Contents lists available at ScienceDirect
Journal of Alloys and Compounds journal homepage: www.elsevier.com/locate/jalcom
Letter
Achieving high strength in indirectly-extruded binary Mg–Ca alloy containing Guinier–Preston zones Hucheng Pan, Gaowu Qin ⇑, Yuping Ren, Liqing Wang, Shineng Sun, Xiangying Meng Key Laboratory for Anisotropy and Texture of Materials (Ministry of Education), Northeastern University, Shenyang 110819, China
a r t i c l e
i n f o
Article history: Received 28 October 2014 Received in revised form 10 January 2015 Accepted 13 January 2015 Available online 19 January 2015 Keywords: Magnesium alloys Mg–Ca Indirect extrusion G.P. zone Mechanical properties
a b s t r a c t We developed a high strength Mg–1 wt.% Ca binary alloy by an integrated method comprising of conventional casting, homogenization and indirect extrusion. The as-extruded Mg–Ca alloy exhibits superior mechanical properties compared with the currently reported results, possessing a tensile yield strength of 310 MPa and an ultimate tensile strength of 330 MPa, which is attributed to the combined effects of fine dynamically recrystallized grains and nanoscale Mg2Ca heterogeneous phases. It is worthy to emphasize that a high density of Guinier–Preston (G.P.) zones appear and are expected to play important roles in precipitation hardening and pinning the grain growth of a-Mg matrix. Ó 2015 Elsevier B.V. All rights reserved.
1. Introduction Magnesium alloys have received increasing interests as structural materials due to unique advantages such as low density and specific strength, and thus they have been widely utilized in vehicles to achieve a lightweight design [1–3]. Besides that, Mg alloys, with proper mechanical and biocompatible properties, have great potentials as load-bearing degradable biomaterials, for example, artificial human bones [4,5]. Among various candidates, it has been confirmed that the corrosion resistance of binary Mg–Ca alloys (Ca < 2 wt.%) is the mostly comparable with pure Mg metal because of the anodic role of Mg2Ca phase in the Mg–Ca binary alloys [6], which thus results in a lower degradation rate than the other Mg-based biomaterials [7]. Moreover, the released Ca component is beneficial for the bone regeneration. With these advantages, Mg–Ca binary alloys are expected to be one of the most promising candidates for high-performance biomaterials [7]. Recent studies have been focused on improving mechanical properties of Mg–Ca binary alloys [6–8]. For example, Seong et al. [6] improved the strength of binary Mg–Ca alloy through a microstructure refinement via a high-ratio differential speed rolling method. Recently, by combining with an indirect extrusion process, the extruded Mg–1 wt.% Ca alloy was achieved, reaching a moderate yield strength (YS) of 185 MPa and an ultimate strength (UTS) of 239 MPa [7]. However, a further improvement ⇑ Corresponding author. E-mail address: [email protected] (G. Qin). http://dx.doi.org/10.1016/j.jallcom.2015.01.068 0925-8388/Ó 2015 Elsevier B.V. All rights reserved.
on the strength of binary Mg–Ca alloys is still expected to maintain a long-term structural stability during biodegradation. In fact, the considerable efforts have been devoted to resolve these issues by alloying and/or deformation in this field [9,10]. Precipitation hardening has been proven as an effective route in newly developed wrought Mg alloys, such as the low-cost Mg–Zn-, Mg–Sn- and Mg–Ca-based systems [11–15]. For example, the twin roll cast Mg–6.2Zn–0.5Zr–0.2Ca alloy shows both high YS (>300 MPa) and excellent formability due to uniform precipitation of Mg(Zn, Zr) particles [12]. With presence of high melting point Mg2Sn phases, the both Mg–8Sn–1Zn–1Al wt.% and Mg–10Sn–3Zn–1Al wt.% alloy could be extruded at a wide ram speed of 0.1–6.7 mm/s, and YS of the samples range from 199 MPa to 319 MPa, with the elongation of 6–15% [13,14]. Moreover, grain refinement hardening has also been employed widely. For example, the novel method of aging prior to extrusion was applied on Mg–7.6Al–0.4Zn wt.% alloy, thus grain size was reduced to be 8.4 lm and YS of the sample could reach 210 MPa [16]. The Mg–3Al–1Zn wt.% alloy with YS of 220 MPa was successfully extruded at low temperature of 250 °C by indirect extrusion method, and the grain size of 1.8 lm was obtained [17]. Recently, high-pressure torsion was applied on both the binary Mg–3.4 at.% Zn and ZK60 alloys, ultra-fine grains in size of 0.14–1 lm were produced [18,19]. All these studies evidently indicate that both the grain refinement and precipitation of hardening phases can significantly contribute to the enhancement of mechanical properties of Mg alloys. More interestingly, Yu et al. [20] showed that an increased density of dynamic precipitates can be generated at a relatively low
H. Pan et al. / Journal of Alloys and Compounds 630 (2015) 272–276
extrusion temperature for the Mg–8Al–0.5Zn wt.% alloys, effectively suppressing the growth of dynamically recrystallized grains of a-Mg matrix, thus a high YS of 403 MPa was obtained. It is also well known that a homogenization treatment procedure can increase the density of nano-scale heterogeneous particles, resulting from the precipitation during the following extrusion [15,21]. In this context, we herein applied an indirect extrusion to the sufficiently homogenized binary Mg–1Ca ingot at a relatively low temperature. By integrating with conventional casting, homogenization and indirect extrusion together, we reasonably expect to achieve high-performance mechanical properties of Mg–Ca alloys. 2. Experimental The Mg–1Ca alloy ingot was prepared by melting commercial pure Mg, Ca in an electrical resistance furnace, and 0.1 wt.% Mn was added to remove the Fe impurity in the Mg raw material during melting the Mg–Ca alloy [22], because trace Fe had serious deterioration effects on corrosion resistance. Finally, the melt was poured into the pre-heated steel mold with the protection of mixture gas of CO2 and SF6 (100:1). The actual contents of Ca and Mn elements in the ingot were chemically analyzed to be 1.05 wt.% and 0.097 wt.%. Homogenization treatment was conducted at 500 °C for 12 h with protection of graphite powders, and then the sample was quenched into cold water. The ingot was indirectly-extruded at 300 °C and 350 °C, respectively (named as 1Ca-300 and 1Ca-350) with an extrusion ratio of 20 and a ram speed of 2 mm/s. The mechanical properties of bars were tested on the Shimadzu AG-X Plus 250 kN at an initial tensile strain rate of 1 103 s1. The X-ray diffraction (XRD, Philips PW3040/60 X’Pert PRO with Cu Ka radiation) was used to identify constituent phases of the as-cast and as-extruded samples. The developed textures of extruded samples were measured via XRD. The microstructures of specimens were examined by optical microscope (OM, Olympus GX), scanning electron microscope (SEM, JEOL JEM-2100F, accelerating voltage of 20 kV, working distance of 10 mm) equipped with energy dispersive spectrometer (EDS) and transmission electron microscope (TEM, JEOL JSM-7001F). Thin foils for TEM were prepared by mechanical polishing (40 lm) and then ion beam thinning (GATAN, working voltage of 4.5 V and thinning angle of 7° between the incident beam and the sample surface, at vacuum and cooled atmosphere by liquid nitrogen). TEM observation was conducted at an accelerating voltage of 200 kV.
3. Results and discussion Fig. 1 shows SEM images of the as-cast and extruded bars. The as-cast sample (Fig. 1a) exhibits dendritic a-Mg microstructures with inter-dendritic eutectic-phases. After the homogenization and extrusion processing (Fig. 1b), most of the bulk phases in the as-cast sample were dissolved into matrix, while the other phases are seriously fractured and distributed homogeneously. XRD patterns (Fig. 1c) and EDS results (not shown) furthermore confirm that the bulk and dotty contrasts correspond to the Mg2Ca phases. Fig. 2 shows a bimodal microstructure of the as-extruded samples, consisting of equiaxed DRXed grains and elongated unDRXed grains (the area indicated by the arrows in Fig. 2a and b). The volume fraction of unDRXed grains in the sample extruded at 300 °C is 15% (Fig. 2a), and the average unDRXed and DRXed grain sizes are 10 lm and 1.0 lm (Fig. 2a and c), respectively. With the extrusion temperature increasing to 350 °C (Fig. 2b and d), the volume fraction of unDRXed grains, average unDRXed and DRXed grain
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size evolve to be 5%, 5 lm and 1.2 lm, respectively. A large number of Mg- and Ca-enriched nano-particles, confirmed by EDS, are also observed both among the matrix and along the grain boundaries, as marked by the arrows in Fig. 2c and d. Fig. 2c1 and c2 shows the enlarged morphology of a typical nano-particle and the corresponding Fast Fourier Transformation (FFT) pattern, which confirm that the nano-particle is the Mg2Ca phase and a new orientation relationship with matrix can be determined as [1 1 2 0]p || [1 2 1 0]a and (0 0 0 1)p || (1 0 1 1)a. It is believed that these nano-Mg2Ca phases could pin the DRXed grain boundaries and suppress the grain growth during extrusion, leading to the severe refinement of DRXed grains (Fig. 2c and d) [23]. The (0 0 0 2) X-ray pole figures of the Mg–Ca samples extruded at both 300 °C and 350 °C are analyzed to assess the grain orientation evolution, as shown in Fig. 2e. The sample extruded at 300 °C exhibits an unusual fiber texture with h0 0 0 1i axis of the Mg matrix tilting about 30° away from the radial direction, and a maximum texture intensity of 16.9. With increasing extrusion temperature to 350 °C, more random texture with a decreased intensity of 10.3 is detected. Previous studies demonstrated that the DRXed grains usually represent a relatively weak texture. For example, the grains nucleated by the particle stimulated nucleation (PSN) mechanism are proven to be random in orientation, which leads to the weak texture [15,24]. So the low texture intensity in the sample extruded at 350 °C might be attributed to the increased volume fraction of the DRXed grains (Fig. 2b). Fig. 3a and b shows the bright field TEM images of the Mg–Ca samples extruded at both 300 °C and 350 °C with electron beam parallel to [1 1 2 0]a, and the corresponding selected area electron diffraction (SAED) patterns are displayed in Fig. 3A and B, respectively. A considerable number of plate-like precipitates, with 1 nm in thickness and 40 nm in length, lie on the basal plane of the a-Mg matrix for the sample extruded at 300 °C (Fig. 3a). TEM image in a higher magnification taken from the area A marked by a circle in Fig. 3a is shown in Fig. 3c. Together with EDS analyzes, the rectangular regions (position 1) and the plate-like regions (position 2) are verified to be Ca-enriched (Fig. 3c1). The rectangular phase with 35 nm in length and 18 nm in width, was also observed in the over-aged Mg–0.3Ca–0.3Zn (wt.%) alloys, which was identified as Mg2Ca phase in Ref. [25]. Moreover, it is usually recognized that Mg2Ca phase is the only precipitation phase in the Mg-rich region of Mg–Ca binary alloys, as confirmed by the electron diffraction in this work. Based on these evidences, the rectangular regions are reasonably thought to be Mg2Ca phases in the Mg–Ca alloys. On the other hand, the plate-like phase contains less content of Ca element (Fig. 3c2). Due to the limited size in thickness, it is hard to determine its accurate composition by TEM–EDS. However, it is clearly shown that Ca is enriched in this plate-like phase as compared to the a-Mg matrix. The corresponding FFT patterns (Fig. 3 inset) fail to reveal any apparent other reflections, besides the streaks parallel to (0 0 0 1)a plane, which are consistent with the
Fig. 1. SEM images of the (a) as-cast sample and (b) sample extruded at 300 °C observed over cross-section of rods, (c) XRD patterns for the as-cast and the extruded samples at 300 °C.
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Fig. 2. (a) and (b) Optical micrographs and (c) and (d) TEM images of the as-extruded samples. (a) and (c) are the sample extruded at 300 °C, and (b) and (d) are the sample extruded at 350 °C. The enlarged morphology of typical nano-particles in (c) and the corresponding FFT patterns are included in c1 and c2. Textures of the extruded samples are also shown in (e). Extrusion direction and radial direction are marked as ED and RD, respectively.
Fig. 3. Bright field TEM images of samples extruded at the (a) 300 °C and (b) 350 °C with electron beam parallel to [1 1 2 0]a. Diffraction patterns of region A and B are displayed. (c) The higher magnification TEM image taken from the area marked as A in (a), containing the rectangular nano-phases and G.P. zones. c1and c2 are EDX results for the point 1 and 2 (in (c)), respectively. The corresponding FFT patterns of the area 2 marked as square are inserted as inset in (c). The radial direction of the extruded rod is marked as RD.
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Fig. 4. Engineering tensile stress–strain curves of the Mg–1Ca samples extruded at 300 °C and 350 °C, respectively.
patterns of G.P. zones observed in Mg–Ca–Zn alloys [25,26] and other Mg–RE alloys [27]. Based on these observations, the platelike phases are considered to be coherent with a-Mg matrix and are thus regarded as G.P. zones. In the sample extruded at 350 °C, similar microstructures of numerous G.P. zones can be observed (Fig. 3b). However, by comparing two samples, it can be found that size of G.P. zones (about 0.5 nm in width and 18 nm in length) in the sample extruded at 350 °C is much smaller than the one at 300 °C and bulk Mg2Ca phases are rarely detected. In fact, G.P. zones formation requires both mass transfer and presence of the sufficient density of vacancies. A high density of defects, such as vacancies and dislocations, could be generated in Mg alloys processed by plastic deformation [28]. Even though excess vacancies would decrease rapidly to the half of its original state within a few hours due to the migrations to the dislocations and grain boundaries, it finally presents a slow decay and approaches an equivalent level [29]. It can be expected that the present extrusion processing at high temperature can also provide a large number of defects. Moreover, high concentrations of vacancies could have been formed during the high-temperature homogenization process (500 °C) which could be remained in the a-Mg matrix by quenching. The stress field around these defects would become the driving force for mass transfer [30], resulting in the formation of the solute clusters. The high density of vacancies can also approach the solute to form the vacancy–Ca clusters. Moreover, the dislocations pipe and vacancies flux can provide fast diffusion paths for the solute clusters [31], which eventually evolve to be the G.P. zones. However, it is worthy to note that the G.P. zones formation strongly depends on the concentration of dislocations/vacancies in metallic alloys [32]. At the extrusion temperature of 350 °C for the present Mg–Ca system, the internal stress could be released and the dynamic recovery is promoted [33]. Thus the concentration of dislocations/vacancies can be decreased significantly, and the dynamic formation of G.P. zones is restricted, as shown in Fig. 3b, revealing a reduced size of the G.P. zones. Nevertheless, the above observation shows the dynamical precipitation behavior of the extruded Mg–Ca alloy is extremely different from the conventional quench-aging of binary Mg–Ca alloy, where only the basal type Mg2Ca phases are precipitated [34]. As to the
thermodynamic reason for the G.P. zones formation in the Mg–Ca binary system, it needs a further study in the near future. Fig. 4 shows the engineering tensile stress–strain curves of the as-extruded Mg–1Ca samples. The sample extruded at 300 °C exhibits a YS of 310 MPa, an UTS of 330 MPa and an elongation to failure (EL) of 4%. For the sample extruded at 350 °C, the YS and the UTS decrease slightly to be 280 MPa and 310 MPa, respectively, while the EL increases up to 9.5%. The YS of sample extruded at 300 °C is higher than those of previously reported Mg–Ca binary alloys, such as Mg–1Ca (185 MPa), Mg–2Ca (205 MPa) and Mg–3Ca (249 MPa) [7]. Besides, the YS of the present sample is higher than the commercial AZ31 extruded at 270 °C (178 MPa) [35], and is even higher than those of the newly developed high-strength Mg–6Zn–0.4Ag–0.2Ca–0.6Zr (ZQXK6000) alloy (288 MPa) [23], Mg–4.7Zn–0.5Ca (ZX50) alloy (291 MPa) [36] and Mg–9.8Sn–1.2Zn–1.0Al alloy (TZA911) (308 MPa) [35]. The detailed comparisons of yield strengths among these alloys are listed in Table 1. In general, the enhanced YS of extruded sample is evidently related to the grain size of the a-Mg matrix, texture and dynamic precipitation. Grain refinement strengthening mechanism is usu0:5 ally described by the Hall–Petch (HP) equation: rys ¼ r0 þ ky d , where d is the average grain diameter, r0 is the friction stress and ky is the HP slope [37]. Previous studies showed that the parameters for HP relationship in Mg alloys depend on the texture and grain size, and the hardening effect is more obvious than that of Al alloys due to the higher ky value [38]. It is shown that the mean grain size of the as-extruded samples are approximately 1.4–2.3 lm by averaging the unDRXed and DRXed grain regions, and consequently such grains refinement would play important roles in achieving the high YS for the present 1Ca-300 and 1Ca350 alloys. Following the texture that h0 0 0 1i of Mg matrix tilted about 30° away from the radial direction, both the basal and prismatic slip can be activated during tensile deformation since the Schmid factor for the basal slip increases [39]. Considering the low critical shear stress of the basal slip [40], the YS of the sample could be reduced compared to the case with a strong fiber texture. However, the precipitation strengthening due to the copious G.P. zones and nano-Mg2Ca phases can compensate the strength loss from texture. Notably, the orientation of G.P. zones with respect to the Mg matrix is critical for the strengthening effect. Nie [41] reported that plate-like precipitates formed on prismatic planes of Mg matrix are most effective in dispersive strengthening by developing the Orowan equation that is appropriate for Mg alloys with rods/plates precipitate. The comparison result also shows that plate phases formed on the basal plane, e.g. the G.P. zones in present samples, could considerably contribute to the strength of Mg alloys by providing barriers for gliding dislocations, if their interspacing is small enough. Accordingly, the G.P. zones uniformly distributed in as-extruded sampled with interspaces of about 45 nm would readily enhance the YS substantially [42], where the additional stresses have to be applied on dislocations to overcome the obstacles from the second phases. The much higher tensile strength achieved in this study compared with previous work of the binary Mg–Ca wrought alloys can be attributed to the processing method of a lower extrusion temperature and a more sufficient solution treatment [6–8]. On the other hand, the higher ductility of 9.5% in the sample extruded at 350 °C is considered resulting from the higher volume fraction of DRXed grains and
Table 1 Comparison of the YS among the present Mg–1Ca samples extruded at 300 °C, 350 °C and the samples reported in literatures. Samples
1Ca-300
1Ca-350
1Ca [7]
2Ca [7]
3Ca [7]
AZ31 [35]
ZQXK6000 [23]
ZX50 [36]
TZA911 [35]
YS (MPa)
310
280
185
205
249
178
288
291
308
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the weaker texture intensity, where more dislocations on the basal plane could be induced due to the easier activation of basal hai slip mechanism, thus more plastic deformation could be accumulated [43,44]. In the as-extruded Mg–Ca alloy, the fine grain microstructure stems from the dynamic recrystallization [23,35]. The precipitations play important roles in refining a-Mg grain size of the extruded binary Mg–Ca alloy. As seen in Figs. 2 and 3, a large number of nano-scale Mg2Ca particles and the G.P. zones disperse along grain boundaries and among matrix. Since both the high density of dislocations and vacancies are introduced during extrusion, the defects can provide the driving force and the nucleation sites for forming the precipitates, which in turn drag the grain boundary movement and suppress the grain growth of a-Mg matrix [35,36], and thus contribute to an enhanced strength in Mg– 1 wt.% Ca alloy. 4. Conclusions In summary, Mg–1 wt.% Ca alloys were indirectly extruded at 300 °C and 350 °C, respectively, resulting in the formation of ultra-fine DRXed grain size of 1.0–1.2 lm, and a high density of G.P. zones. The extruded bar exhibited a high YS of 310 MPa and an UTS of 330 MPa. The microstructure analyzes evidenced that the high strength was mainly resulted from the refined DRXed grains, as well as the combined effect of the nano-scale heterogeneous Mg2Ca particles and the G.P. zones.
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Journal of Alloys and Compounds journal homepage: www.elsevier.com/locate/jalcom
Letter
Achieving high strength in indirectly-extruded binary Mg–Ca alloy containing Guinier–Preston zones Hucheng Pan, Gaowu Qin ⇑, Yuping Ren, Liqing Wang, Shineng Sun, Xiangying Meng Key Laboratory for Anisotropy and Texture of Materials (Ministry of Education), Northeastern University, Shenyang 110819, China
a r t i c l e
i n f o
Article history: Received 28 October 2014 Received in revised form 10 January 2015 Accepted 13 January 2015 Available online 19 January 2015 Keywords: Magnesium alloys Mg–Ca Indirect extrusion G.P. zone Mechanical properties
a b s t r a c t We developed a high strength Mg–1 wt.% Ca binary alloy by an integrated method comprising of conventional casting, homogenization and indirect extrusion. The as-extruded Mg–Ca alloy exhibits superior mechanical properties compared with the currently reported results, possessing a tensile yield strength of 310 MPa and an ultimate tensile strength of 330 MPa, which is attributed to the combined effects of fine dynamically recrystallized grains and nanoscale Mg2Ca heterogeneous phases. It is worthy to emphasize that a high density of Guinier–Preston (G.P.) zones appear and are expected to play important roles in precipitation hardening and pinning the grain growth of a-Mg matrix. Ó 2015 Elsevier B.V. All rights reserved.
1. Introduction Magnesium alloys have received increasing interests as structural materials due to unique advantages such as low density and specific strength, and thus they have been widely utilized in vehicles to achieve a lightweight design [1–3]. Besides that, Mg alloys, with proper mechanical and biocompatible properties, have great potentials as load-bearing degradable biomaterials, for example, artificial human bones [4,5]. Among various candidates, it has been confirmed that the corrosion resistance of binary Mg–Ca alloys (Ca < 2 wt.%) is the mostly comparable with pure Mg metal because of the anodic role of Mg2Ca phase in the Mg–Ca binary alloys [6], which thus results in a lower degradation rate than the other Mg-based biomaterials [7]. Moreover, the released Ca component is beneficial for the bone regeneration. With these advantages, Mg–Ca binary alloys are expected to be one of the most promising candidates for high-performance biomaterials [7]. Recent studies have been focused on improving mechanical properties of Mg–Ca binary alloys [6–8]. For example, Seong et al. [6] improved the strength of binary Mg–Ca alloy through a microstructure refinement via a high-ratio differential speed rolling method. Recently, by combining with an indirect extrusion process, the extruded Mg–1 wt.% Ca alloy was achieved, reaching a moderate yield strength (YS) of 185 MPa and an ultimate strength (UTS) of 239 MPa [7]. However, a further improvement ⇑ Corresponding author. E-mail address: [email protected] (G. Qin). http://dx.doi.org/10.1016/j.jallcom.2015.01.068 0925-8388/Ó 2015 Elsevier B.V. All rights reserved.
on the strength of binary Mg–Ca alloys is still expected to maintain a long-term structural stability during biodegradation. In fact, the considerable efforts have been devoted to resolve these issues by alloying and/or deformation in this field [9,10]. Precipitation hardening has been proven as an effective route in newly developed wrought Mg alloys, such as the low-cost Mg–Zn-, Mg–Sn- and Mg–Ca-based systems [11–15]. For example, the twin roll cast Mg–6.2Zn–0.5Zr–0.2Ca alloy shows both high YS (>300 MPa) and excellent formability due to uniform precipitation of Mg(Zn, Zr) particles [12]. With presence of high melting point Mg2Sn phases, the both Mg–8Sn–1Zn–1Al wt.% and Mg–10Sn–3Zn–1Al wt.% alloy could be extruded at a wide ram speed of 0.1–6.7 mm/s, and YS of the samples range from 199 MPa to 319 MPa, with the elongation of 6–15% [13,14]. Moreover, grain refinement hardening has also been employed widely. For example, the novel method of aging prior to extrusion was applied on Mg–7.6Al–0.4Zn wt.% alloy, thus grain size was reduced to be 8.4 lm and YS of the sample could reach 210 MPa [16]. The Mg–3Al–1Zn wt.% alloy with YS of 220 MPa was successfully extruded at low temperature of 250 °C by indirect extrusion method, and the grain size of 1.8 lm was obtained [17]. Recently, high-pressure torsion was applied on both the binary Mg–3.4 at.% Zn and ZK60 alloys, ultra-fine grains in size of 0.14–1 lm were produced [18,19]. All these studies evidently indicate that both the grain refinement and precipitation of hardening phases can significantly contribute to the enhancement of mechanical properties of Mg alloys. More interestingly, Yu et al. [20] showed that an increased density of dynamic precipitates can be generated at a relatively low
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extrusion temperature for the Mg–8Al–0.5Zn wt.% alloys, effectively suppressing the growth of dynamically recrystallized grains of a-Mg matrix, thus a high YS of 403 MPa was obtained. It is also well known that a homogenization treatment procedure can increase the density of nano-scale heterogeneous particles, resulting from the precipitation during the following extrusion [15,21]. In this context, we herein applied an indirect extrusion to the sufficiently homogenized binary Mg–1Ca ingot at a relatively low temperature. By integrating with conventional casting, homogenization and indirect extrusion together, we reasonably expect to achieve high-performance mechanical properties of Mg–Ca alloys. 2. Experimental The Mg–1Ca alloy ingot was prepared by melting commercial pure Mg, Ca in an electrical resistance furnace, and 0.1 wt.% Mn was added to remove the Fe impurity in the Mg raw material during melting the Mg–Ca alloy [22], because trace Fe had serious deterioration effects on corrosion resistance. Finally, the melt was poured into the pre-heated steel mold with the protection of mixture gas of CO2 and SF6 (100:1). The actual contents of Ca and Mn elements in the ingot were chemically analyzed to be 1.05 wt.% and 0.097 wt.%. Homogenization treatment was conducted at 500 °C for 12 h with protection of graphite powders, and then the sample was quenched into cold water. The ingot was indirectly-extruded at 300 °C and 350 °C, respectively (named as 1Ca-300 and 1Ca-350) with an extrusion ratio of 20 and a ram speed of 2 mm/s. The mechanical properties of bars were tested on the Shimadzu AG-X Plus 250 kN at an initial tensile strain rate of 1 103 s1. The X-ray diffraction (XRD, Philips PW3040/60 X’Pert PRO with Cu Ka radiation) was used to identify constituent phases of the as-cast and as-extruded samples. The developed textures of extruded samples were measured via XRD. The microstructures of specimens were examined by optical microscope (OM, Olympus GX), scanning electron microscope (SEM, JEOL JEM-2100F, accelerating voltage of 20 kV, working distance of 10 mm) equipped with energy dispersive spectrometer (EDS) and transmission electron microscope (TEM, JEOL JSM-7001F). Thin foils for TEM were prepared by mechanical polishing (40 lm) and then ion beam thinning (GATAN, working voltage of 4.5 V and thinning angle of 7° between the incident beam and the sample surface, at vacuum and cooled atmosphere by liquid nitrogen). TEM observation was conducted at an accelerating voltage of 200 kV.
3. Results and discussion Fig. 1 shows SEM images of the as-cast and extruded bars. The as-cast sample (Fig. 1a) exhibits dendritic a-Mg microstructures with inter-dendritic eutectic-phases. After the homogenization and extrusion processing (Fig. 1b), most of the bulk phases in the as-cast sample were dissolved into matrix, while the other phases are seriously fractured and distributed homogeneously. XRD patterns (Fig. 1c) and EDS results (not shown) furthermore confirm that the bulk and dotty contrasts correspond to the Mg2Ca phases. Fig. 2 shows a bimodal microstructure of the as-extruded samples, consisting of equiaxed DRXed grains and elongated unDRXed grains (the area indicated by the arrows in Fig. 2a and b). The volume fraction of unDRXed grains in the sample extruded at 300 °C is 15% (Fig. 2a), and the average unDRXed and DRXed grain sizes are 10 lm and 1.0 lm (Fig. 2a and c), respectively. With the extrusion temperature increasing to 350 °C (Fig. 2b and d), the volume fraction of unDRXed grains, average unDRXed and DRXed grain
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size evolve to be 5%, 5 lm and 1.2 lm, respectively. A large number of Mg- and Ca-enriched nano-particles, confirmed by EDS, are also observed both among the matrix and along the grain boundaries, as marked by the arrows in Fig. 2c and d. Fig. 2c1 and c2 shows the enlarged morphology of a typical nano-particle and the corresponding Fast Fourier Transformation (FFT) pattern, which confirm that the nano-particle is the Mg2Ca phase and a new orientation relationship with matrix can be determined as [1 1 2 0]p || [1 2 1 0]a and (0 0 0 1)p || (1 0 1 1)a. It is believed that these nano-Mg2Ca phases could pin the DRXed grain boundaries and suppress the grain growth during extrusion, leading to the severe refinement of DRXed grains (Fig. 2c and d) [23]. The (0 0 0 2) X-ray pole figures of the Mg–Ca samples extruded at both 300 °C and 350 °C are analyzed to assess the grain orientation evolution, as shown in Fig. 2e. The sample extruded at 300 °C exhibits an unusual fiber texture with h0 0 0 1i axis of the Mg matrix tilting about 30° away from the radial direction, and a maximum texture intensity of 16.9. With increasing extrusion temperature to 350 °C, more random texture with a decreased intensity of 10.3 is detected. Previous studies demonstrated that the DRXed grains usually represent a relatively weak texture. For example, the grains nucleated by the particle stimulated nucleation (PSN) mechanism are proven to be random in orientation, which leads to the weak texture [15,24]. So the low texture intensity in the sample extruded at 350 °C might be attributed to the increased volume fraction of the DRXed grains (Fig. 2b). Fig. 3a and b shows the bright field TEM images of the Mg–Ca samples extruded at both 300 °C and 350 °C with electron beam parallel to [1 1 2 0]a, and the corresponding selected area electron diffraction (SAED) patterns are displayed in Fig. 3A and B, respectively. A considerable number of plate-like precipitates, with 1 nm in thickness and 40 nm in length, lie on the basal plane of the a-Mg matrix for the sample extruded at 300 °C (Fig. 3a). TEM image in a higher magnification taken from the area A marked by a circle in Fig. 3a is shown in Fig. 3c. Together with EDS analyzes, the rectangular regions (position 1) and the plate-like regions (position 2) are verified to be Ca-enriched (Fig. 3c1). The rectangular phase with 35 nm in length and 18 nm in width, was also observed in the over-aged Mg–0.3Ca–0.3Zn (wt.%) alloys, which was identified as Mg2Ca phase in Ref. [25]. Moreover, it is usually recognized that Mg2Ca phase is the only precipitation phase in the Mg-rich region of Mg–Ca binary alloys, as confirmed by the electron diffraction in this work. Based on these evidences, the rectangular regions are reasonably thought to be Mg2Ca phases in the Mg–Ca alloys. On the other hand, the plate-like phase contains less content of Ca element (Fig. 3c2). Due to the limited size in thickness, it is hard to determine its accurate composition by TEM–EDS. However, it is clearly shown that Ca is enriched in this plate-like phase as compared to the a-Mg matrix. The corresponding FFT patterns (Fig. 3 inset) fail to reveal any apparent other reflections, besides the streaks parallel to (0 0 0 1)a plane, which are consistent with the
Fig. 1. SEM images of the (a) as-cast sample and (b) sample extruded at 300 °C observed over cross-section of rods, (c) XRD patterns for the as-cast and the extruded samples at 300 °C.
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Fig. 2. (a) and (b) Optical micrographs and (c) and (d) TEM images of the as-extruded samples. (a) and (c) are the sample extruded at 300 °C, and (b) and (d) are the sample extruded at 350 °C. The enlarged morphology of typical nano-particles in (c) and the corresponding FFT patterns are included in c1 and c2. Textures of the extruded samples are also shown in (e). Extrusion direction and radial direction are marked as ED and RD, respectively.
Fig. 3. Bright field TEM images of samples extruded at the (a) 300 °C and (b) 350 °C with electron beam parallel to [1 1 2 0]a. Diffraction patterns of region A and B are displayed. (c) The higher magnification TEM image taken from the area marked as A in (a), containing the rectangular nano-phases and G.P. zones. c1and c2 are EDX results for the point 1 and 2 (in (c)), respectively. The corresponding FFT patterns of the area 2 marked as square are inserted as inset in (c). The radial direction of the extruded rod is marked as RD.
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Fig. 4. Engineering tensile stress–strain curves of the Mg–1Ca samples extruded at 300 °C and 350 °C, respectively.
patterns of G.P. zones observed in Mg–Ca–Zn alloys [25,26] and other Mg–RE alloys [27]. Based on these observations, the platelike phases are considered to be coherent with a-Mg matrix and are thus regarded as G.P. zones. In the sample extruded at 350 °C, similar microstructures of numerous G.P. zones can be observed (Fig. 3b). However, by comparing two samples, it can be found that size of G.P. zones (about 0.5 nm in width and 18 nm in length) in the sample extruded at 350 °C is much smaller than the one at 300 °C and bulk Mg2Ca phases are rarely detected. In fact, G.P. zones formation requires both mass transfer and presence of the sufficient density of vacancies. A high density of defects, such as vacancies and dislocations, could be generated in Mg alloys processed by plastic deformation [28]. Even though excess vacancies would decrease rapidly to the half of its original state within a few hours due to the migrations to the dislocations and grain boundaries, it finally presents a slow decay and approaches an equivalent level [29]. It can be expected that the present extrusion processing at high temperature can also provide a large number of defects. Moreover, high concentrations of vacancies could have been formed during the high-temperature homogenization process (500 °C) which could be remained in the a-Mg matrix by quenching. The stress field around these defects would become the driving force for mass transfer [30], resulting in the formation of the solute clusters. The high density of vacancies can also approach the solute to form the vacancy–Ca clusters. Moreover, the dislocations pipe and vacancies flux can provide fast diffusion paths for the solute clusters [31], which eventually evolve to be the G.P. zones. However, it is worthy to note that the G.P. zones formation strongly depends on the concentration of dislocations/vacancies in metallic alloys [32]. At the extrusion temperature of 350 °C for the present Mg–Ca system, the internal stress could be released and the dynamic recovery is promoted [33]. Thus the concentration of dislocations/vacancies can be decreased significantly, and the dynamic formation of G.P. zones is restricted, as shown in Fig. 3b, revealing a reduced size of the G.P. zones. Nevertheless, the above observation shows the dynamical precipitation behavior of the extruded Mg–Ca alloy is extremely different from the conventional quench-aging of binary Mg–Ca alloy, where only the basal type Mg2Ca phases are precipitated [34]. As to the
thermodynamic reason for the G.P. zones formation in the Mg–Ca binary system, it needs a further study in the near future. Fig. 4 shows the engineering tensile stress–strain curves of the as-extruded Mg–1Ca samples. The sample extruded at 300 °C exhibits a YS of 310 MPa, an UTS of 330 MPa and an elongation to failure (EL) of 4%. For the sample extruded at 350 °C, the YS and the UTS decrease slightly to be 280 MPa and 310 MPa, respectively, while the EL increases up to 9.5%. The YS of sample extruded at 300 °C is higher than those of previously reported Mg–Ca binary alloys, such as Mg–1Ca (185 MPa), Mg–2Ca (205 MPa) and Mg–3Ca (249 MPa) [7]. Besides, the YS of the present sample is higher than the commercial AZ31 extruded at 270 °C (178 MPa) [35], and is even higher than those of the newly developed high-strength Mg–6Zn–0.4Ag–0.2Ca–0.6Zr (ZQXK6000) alloy (288 MPa) [23], Mg–4.7Zn–0.5Ca (ZX50) alloy (291 MPa) [36] and Mg–9.8Sn–1.2Zn–1.0Al alloy (TZA911) (308 MPa) [35]. The detailed comparisons of yield strengths among these alloys are listed in Table 1. In general, the enhanced YS of extruded sample is evidently related to the grain size of the a-Mg matrix, texture and dynamic precipitation. Grain refinement strengthening mechanism is usu0:5 ally described by the Hall–Petch (HP) equation: rys ¼ r0 þ ky d , where d is the average grain diameter, r0 is the friction stress and ky is the HP slope [37]. Previous studies showed that the parameters for HP relationship in Mg alloys depend on the texture and grain size, and the hardening effect is more obvious than that of Al alloys due to the higher ky value [38]. It is shown that the mean grain size of the as-extruded samples are approximately 1.4–2.3 lm by averaging the unDRXed and DRXed grain regions, and consequently such grains refinement would play important roles in achieving the high YS for the present 1Ca-300 and 1Ca350 alloys. Following the texture that h0 0 0 1i of Mg matrix tilted about 30° away from the radial direction, both the basal and prismatic slip can be activated during tensile deformation since the Schmid factor for the basal slip increases [39]. Considering the low critical shear stress of the basal slip [40], the YS of the sample could be reduced compared to the case with a strong fiber texture. However, the precipitation strengthening due to the copious G.P. zones and nano-Mg2Ca phases can compensate the strength loss from texture. Notably, the orientation of G.P. zones with respect to the Mg matrix is critical for the strengthening effect. Nie [41] reported that plate-like precipitates formed on prismatic planes of Mg matrix are most effective in dispersive strengthening by developing the Orowan equation that is appropriate for Mg alloys with rods/plates precipitate. The comparison result also shows that plate phases formed on the basal plane, e.g. the G.P. zones in present samples, could considerably contribute to the strength of Mg alloys by providing barriers for gliding dislocations, if their interspacing is small enough. Accordingly, the G.P. zones uniformly distributed in as-extruded sampled with interspaces of about 45 nm would readily enhance the YS substantially [42], where the additional stresses have to be applied on dislocations to overcome the obstacles from the second phases. The much higher tensile strength achieved in this study compared with previous work of the binary Mg–Ca wrought alloys can be attributed to the processing method of a lower extrusion temperature and a more sufficient solution treatment [6–8]. On the other hand, the higher ductility of 9.5% in the sample extruded at 350 °C is considered resulting from the higher volume fraction of DRXed grains and
Table 1 Comparison of the YS among the present Mg–1Ca samples extruded at 300 °C, 350 °C and the samples reported in literatures. Samples
1Ca-300
1Ca-350
1Ca [7]
2Ca [7]
3Ca [7]
AZ31 [35]
ZQXK6000 [23]
ZX50 [36]
TZA911 [35]
YS (MPa)
310
280
185
205
249
178
288
291
308
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the weaker texture intensity, where more dislocations on the basal plane could be induced due to the easier activation of basal hai slip mechanism, thus more plastic deformation could be accumulated [43,44]. In the as-extruded Mg–Ca alloy, the fine grain microstructure stems from the dynamic recrystallization [23,35]. The precipitations play important roles in refining a-Mg grain size of the extruded binary Mg–Ca alloy. As seen in Figs. 2 and 3, a large number of nano-scale Mg2Ca particles and the G.P. zones disperse along grain boundaries and among matrix. Since both the high density of dislocations and vacancies are introduced during extrusion, the defects can provide the driving force and the nucleation sites for forming the precipitates, which in turn drag the grain boundary movement and suppress the grain growth of a-Mg matrix [35,36], and thus contribute to an enhanced strength in Mg– 1 wt.% Ca alloy. 4. Conclusions In summary, Mg–1 wt.% Ca alloys were indirectly extruded at 300 °C and 350 °C, respectively, resulting in the formation of ultra-fine DRXed grain size of 1.0–1.2 lm, and a high density of G.P. zones. The extruded bar exhibited a high YS of 310 MPa and an UTS of 330 MPa. The microstructure analyzes evidenced that the high strength was mainly resulted from the refined DRXed grains, as well as the combined effect of the nano-scale heterogeneous Mg2Ca particles and the G.P. zones.
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