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Deformation behavior at room temperature ranges of fine-grained Mg–Mn system alloys
Materials Science & Engineering A 766 (2019) 138384
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Deformation behavior at room temperature ranges of fine-grained Mg–Mn system alloys
T
Hidetoshi Somekawa∗, Dudekula Althaf Basha, Alok Singh Research Center for Structural Materials, National Institute for Materials Science, 1-2-1 Sengen, Tsukuba, Ibaraki, 305-0047, Japan
ARTICLE INFO
ABSTRACT
Keywords: Magnesium Mechanical property Grain boundary sliding Deformation mechanism Grain boundary
Tensile properties and deformation behavior were investigated on Mg dilute binary alloys (Mg-0.1Bi and Mg-0.1Zr alloys) and Mg–Mn based ternary alloys (Mg-0.3Mn-0.1Bi and Mg-0.3Mn-0.1Zr alloys). These alloys were processed by hot extrusion to obtain fine-grained structures with an average grain size of ~1 μm. Microstructural observations showed that Mn element was segregated at grain boundaries in the ternary alloys. Precipitates consisting of the αMn and/or Mg3Bi2 phase were dispersed in the matrix in the alloys containing Mn and/or Bi elements. All the extruded alloys showed good tensile property of ductility. The elongation-to-failure in tension of Mg-0.3Mn-0.1Zr alloy was more than 100% even at a quasi-static strain rate of 1 × 10−3/s. The Mg-0.1Bi alloy exhibited the largest value of 250% among all in the present examined conditions. These values are superior to those of conventional wrought processed Mg alloys. These superior properties are found to be due to the contribution of grain boundary sliding (GBS) to deformation. Bi and Zr are alternative elements to Mn for use in binary alloys. On the other hand, in low strain rate regimes, a combination of Mn and such micro-alloying elements led to an increase in flow stress, but reduced the elongation-to-failure. This resulted from a high density of precipitates in the matrix, which became cavity nucleation sites. Moreover, non-equilibrium grain boundary structures, which induced atomistic strains, were also unlikely to promote GBS. The normalized stress vs. strain relations including the previous results in the temperature ranges of 298–423 K is found to be fitted by the well-known constitutive equation.
1. Introduction Alloying has been recognized as the most effective method to improve the mechanical properties of metallic materials. This approach has also been applied to magnesium (Mg) and its alloys, which have the lowest density among the conventional metallic materials, to overcome the serious issues of poor ductility and formability near room temperature. The most desirable effect of alloying in Mg is to reduce the difference in critical resolved shear stress (CRSS) between the basal and nonbasal planes. Using single crystal Mg alloys, it has been shown that the addition of aluminum (Al), zinc (Zn) or lithium (Li) elements brings about a decrease in the differences in CRSS [1–4]. Recent studies via first-principles calculations have provided that some particular kinds of elements, such as rare-earth (RE) elements and alkali group elements, have the same effect [5–8]. Many observations of deformed microstructures have also shown activation of non-basal dislocation slips in Mg alloys containing RE [9–19] and alkali group alloying elements [8,20–22]. Mg-RE (RE = Y, Gd, Ce etc) alloys exhibit much greater deformability and larger ductility as compared to the conventional Mg alloys [23–26]. Even for the non-RE alloying elements, good properties due
∗
to non-basal dislocation activities are observed in, e.g., Mg–Al–Sn and Mg–Al–Zn–Mn–Ca alloys [22,27]. There are several plastic deformation mechanisms (not only dislocation slip but also diffusion, grain boundary sliding (GBS) and twinning) of Mg and its alloys; however, most previous studies have focused on the control of dislocation slips. Alloying with manganese (Mn) has been reported to improve the room temperature ductility [28–31]. Mn is likely to segregate at grain boundaries during the wrought process. This feature of grain boundary segregation plays a major role in enhancing room temperature GBS. Fine-grained Mg–Mn alloys with a high volume fraction of grain boundaries exhibit an elongation-to-failure in tension of more than 100% even in quasi-static strain rate regimes [28]. These Mg–Mn alloys also have similar stretch formability to that of the conventional Al alloys [29]. These results suggest that GBS is possible and a promising deformation mechanism to resolve the critical problems of poor ductility and formability. A high contribution of GBS on deformation is recognized as an alternative strategy for controlling dislocation slips, such as non-basal slip activities. In order to enhance GBS, grain boundary control, i.e., grain boundary characteristic and structure, is well-known to be an effective method [32–36]. The results obtained from
Corresponding author. E-mail address: [email protected] (H. Somekawa).
https://doi.org/10.1016/j.msea.2019.138384 Received 2 June 2019; Received in revised form 31 August 2019; Accepted 4 September 2019 Available online 04 September 2019 0921-5093/ © 2019 Elsevier B.V. All rights reserved.
Materials Science & Engineering A 766 (2019) 138384
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Tensile tests were carried out to investigate the mechanical behavior on all the extruded alloys at room temperature (298 K). The initial strain rates were 1 × 10−2, 10−3, 10−4 and 10−5/s. A strain rate change tensile test was also performed at intermediate temperatures using the Mg-0.3Mn-0.1Bi and Mg-0.3Mn-0.1Zr alloys. The strain rate in tension was changed from 1 × 10−5 to 1 × 10−4/s with a set strain of 0.1. The testing temperatures were selected as 323, 373 and 423 K. Since this test was not performed at room temperature, the holding time for specimen was more than 15 min after reaching the set temperature. The specimens used in all the tensile tests had a gauge length of 10 mm and a diameter of 2.5 mm. Tensile specimens were made by machining from the extruded bars in the direction parallel to the ED. The surface after the tensile test at room temperature at the strain rate of 1 × 10−5/ s was observed by scanning electron microscopy (SEM).
experiments and numerical simulations have pointed out that grain boundaries with higher grain boundary energy accelerate GBS [32,33]. High-angle grain boundaries show larger grain boundary misfits via GBS than that at the special grain boundaries, such as low-angle grain boundaries and coincident site lattice boundaries [34,35]. From the atomic structural viewpoint, equilibrium and non-equilibrium grain boundary structures also affect the temperature and/or the strain rate at which GBS occurs [36]. On application to Mg binary alloys, there are two unique alloying elements: zirconium (Zr) and bismuth (Bi). Firstprinciples calculations have indicated that addition of Zr has an ability to reduce the grain boundary energy [37]. In the case of alloying with Bi, microstructure analysis has revealed that equilibrium grain boundary structures form during the wrought process [38,39]. These two alloying elements have a great potential to enhance GBS. Recently, Xu et al. have investigated the effect of Mn or Zr addition on agehardening response using multi-component system Mg alloy (Mg-Gd-YZn system alloy) [40]; however, to the best of our knowledges, there are no reports about the impact of these alloying elements on the microstructures and mechanical properties of Mg–Mn system ternary alloys. In this study, we have investigated and considered the effect of these micro-alloying elements on the tensile properties and deformation behavior of binary and ternary alloys.
3. Results and discussion The microstructures of extruded alloys observed by EBSD method are shown in Fig. 1. The inverse pole figures show that all the extruded alloys consist of recrystallized grains; most of the grain boundaries are highangle grain boundaries (misorientation angle of more than 15°). The average grain size, d, measured from the EBSD analysis are 0.8–1.2 μm, as provided in Table 1. Although the extrusion temperatures are slightly different, these extruded alloys have similar average grain sizes, because of performing under similar maximum extrusion loads. The binary alloys are able to acquire fine-grained structures at lower extrusion temperatures as compared to those of the ternary alloys. This table also includes the average misorientation angles, θ, of 40°–50° for all the extruded alloys. They are found to be independent of the alloying element. It is also found that these alloys have a basal texture, owing to the extrusion process. The XRD plots are shown in Fig. 2. Apart from the α-Mg, there are no other peaks in the Mg-0.1Zr alloy. In contrast, the alloys containing Bi elements (Mg-0.1Bi and Mg-0.3Mn-0.1Bi alloys) also show the binary Mg3Bi2 phase. In the extruded Mg–Mn–Bi ternary alloy, the existence of α-Mn peak is also confirmed. Although the maximum solubility of Bi and Mn in Mg are 1.13 at.% and 1.0 at%, respectively [41], which are significantly larger than the present chemical composition (0.1 and 0.3 at.%), the solubility of both elements dramatically decrease at the extrusion temperature (388–423 K, as listed in Table 1). Hence, the Mg3Bi2 phase and α-Mn precipitate during the wrought process. The same microstructural features are observed in the previously reported extruded Mg-Bi [38,39] and Mg–Mn binary alloys [28,42]. Typical microstructures of extruded ternary alloys observed by TEM are shown in Fig. 3. These are high-angle annular dark-field (HAADF) images, in which brighter contrast indicates concentration of heavier elements. Some of bright regions, which display the aggregation of elements with large atomic number (the atomic numbers of Mg, Mn, Zr and Bi are 12, 25, 40 and 83, respectively [43]), can be confirmed in both images. They confirm the existence of precipitates formed by these heavy elements, that is agreement with the XRD results (Fig. 2). Fig. 3 also includes the results of energy dispersive X-ray spectroscopy (EDX) analysis in the square area enclosed by a white dashed line. These EDX maps reveal a higher Mn concentration at grain boundaries in both the extruded ternary alloys. Thus, Mn is found to segregate at grain boundaries. Such grain boundary segregation behavior is reported in several wrought processed Mg alloys [9,19,33,44–50] and extruded Mg–Mn binary alloys [28,30]. While the results of TEM observations of the binary alloys are not shown here, segregation of alloying elements at grain boundaries are not observed. This non-segregation grain boundary behavior is the same as that of previous reports [38,39]. The nominal stress vs. nominal strain curves of all the extruded alloys in tensile tests at various strain rates are shown in Fig. 4. All the plots have the same x- and y-axis scale for ease of comparison. Strain rate affects the tensile behavior in all the extruded alloys. The flow stress decreases, but the elongation-to-failure in tension increases at lower strain rate, irrespective of the alloying elements. Even at room
2. Experimental procedure Mg-0.1 at.%X binary alloys and Mg-0.3 at.%Mn-0.1 at.%X ternary alloys (X = Bi or Zr) were used in this study, hereafter denoted as the Mg-0.1Bi, Mg-0.1Zr, Mg-0.3Mn-0.1Bi and Mg-0.3Mn-0.1Zr alloys. Except for Mg-0.1Bi alloy, these alloys were cast by combining a master alloy (Mg-2.6 at.%Mn or Mg-1.5 at.%Zr alloy) with commercial purity Mg. These cast alloys were solution-treated at the temperature of 773 K for 2 h for binary alloys and 24 h for ternary alloys. The solution-treated alloys were subsequently extruded at an extrusion ratio of 25:1 into 8 mm diameter rods. The billet was kept in furnace at the set extrusion temperature for more than 0.5 h, and then was extruded at a speed of 0.2 mm/s. The extrusion temperatures of these alloys are listed in Table 1. The reason for different extrusion temperatures was the limitation of maximum load in the extrusion machine. The phases of all the extruded alloys were determined by X-ray diffraction (XRD) with Cu-Kα radiation. The microstructures of extruded alloys were observed by an electron backscatter diffraction (EBSD, in a field emission gun scanning electron microscope (FE-SEM) equipped with an EDAX-TSL EBSD system) and by transmission electron microscopy (TEM) and highresolution electron microscopy (HREM). The scanning step size for EBSD measurement was 75 nm, and the EBSD results were analyzed using the EDAX software. The area measured and observed by XRD, EBSD, TEM and HREM was at the center of the cross-section containing the extrusion-direction and the transverse-direction (ED-TD). After embedding each alloy in resin, the samples for EBSD observation were prepared by grinding under diamond slurries and colloidal silica slurry, and then were etched using acetic solution for several seconds. The specimens for XRD measurement were prepared by hand polishing to mirror surface. TEM and HREM samples were prepared by ion-milling. Table 1 The experimental condition and results obtained from microstructural analysis and room temperature tensile testing. Text, K
Mg-0.1Bi Mg-0.1Zr Mg-0.3Mn-0.1Bi Mg-0.3Mn-0.1Zr
388 393 423 403
d, μm
0.8 1.0 1.1 1.2
θ, °
44.3 40.9 50.9 47.8
m-value 10−4–10−5/s
10−2–10−3/s
0.25 0.22 0.21 0.18
0.10 0.10 0.12 0.12
where Text is the extrusion temperature, d is the average grain size and θ is the average misorientation angle. 2
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Fig. 1. Initial microstructures observed by EBSD method in extrusion-transverse direction plane; (a) Mg-0.1Bi, (b) Mg-0.1Zr, (c) Mg-0.3Mn-0.1Bi and (d) Mg-0.3Mn0.1Zr alloys. The scale bars in these images are 20 μm.
0.3 at.%Mn binary alloy (hereafter denoted as the Mg-0.3Mn alloy) processed by extrusion [28,52], which have average grain sizes of 1.2 and 1.8 μm, respectively, measured from the EBSD analysis. The flow stress in Fig. 5(b) is cited from stress at a strain of 0.1, in accordance with the Japanese Industrial Standard Method in Ref. [53]. Fig. 5(a) clearly shows the effect of micro-alloying on elongation-to-failure in tension. The binary alloys are superior in ductility compared to pure Mg and Mg-0.3Mn alloy. Focusing on the binary alloys, to attain and to retain good property of ductility, Zr and Bi elements are found to be alternatives to Mn. As for the ternary alloys, the elongation-to-failure in tension is similar to that of the Mg-0.3Mn alloy in the strain rate regimes of 10−2–10−5/s. The elongation-to-failure of the ternary alloys is also larger than that of the present binary alloys at the strain rates of 10−2–10−3/s; nevertheless, the superiority is unlikely to be obtained in the low strain rate range (1 × 10−5/s). The combined addition of Mn and these micro-alloying elements brings about an increase in the flow stress, and is effective for improving ductility in the quasi-static strain rate regimes. Fig. 5(b) shows the dependence of the flow stress on the strain rate. The slope in this figure corresponds to the strain rate sensitivity, i.e., the m-value, which has a close relation to the plastic deformation mechanism [54,55]. The m-value for each alloy is calculated by dividing into two different strain rate regimes, as summarized in Table 1. The m-value decreases with increase in strain rate, and the m-values of the binary alloys are slightly higher than those of the ternary alloys. Nevertheless, all the extruded alloys have much larger m-values than the wrought processed conventional Mg alloys. The m-values of the extruded AZ31 alloy with fine-grained structures are reported to be 0.01 and 0.03 at the quasi-static and low strain rates, respectively [51], where deformation mechanism is dominantly dislocation slip under tensile stress state. This reveals that the plastic deformation mechanism of the present extruded alloys is different from that of wrought processed common Mg alloys. It is known that the contribution of GBS to deformation becomes large with increase in the m-value [54–56]. Typical surface features after tensile tests at room temperature are shown in Fig. 6 for the ternary alloys. The tensile direction is the horizontal in both images. These images apparently show traces of GBS, marked by white arrows. The same feature of GBS is also observed on the surface after room temperature tensile tests in the fine-grained Mg-0.3Bi alloy [39] and the fine-grained Mg-0.3Mn binary alloy [30], which both exhibits a high ductility. The present huge elongation-to-failure in tension is assumed to be related to GBS. GBS is well-recognized as the rate-controlling deformation mechanism for superplastic flow [54–57]. This deformation mechanism is affected by external factors (such as testing temperature and strain rate), and internal factors (microstructure). As external factors, a high temperature and a low strain rate cause GBS to make a large
Fig. 2. XRD peaks of the extruded alloys. Blue and red lines correspond to the α-Mn and the Mg3Bi2 phase, respectively. (For interpretation of the references to colour in this figure legend, the reader is referred to the Web version of this article.)
temperature, all the extruded alloys show an elongation-to-failure of more than 50% at a quasi-static strain rate of 1 × 10−3/s. The Mg0.3Mn-0.1Zr alloy exhibits the best ductility of 105% among the present alloys at this strain rate. In addition, the elongation-to-failure of the binary alloys is over 200% at the strain rate of 1 × 10−5/s. The elongation-to-failure of extruded fine-grained conventional Mg alloys, such as Mg–3Al–1Zn (AZ31) alloy, is at the most 30–50% in the strain rate regimes of 10−2–10−5/s [51]. All the present alloys are found to have good ductility properties. This figure also provides that even though all the extruded alloys have similar average grain sizes (as listed in Table 1), the ternary alloys have higher flow stress than that of the binary alloys. The combination of Mn and micro-alloying elements plays a role in increasing the flow stress, mainly due to the presence of precipitates. The variation in elongation-to-failure in tension and flow stress as a function of initial strain rate is shown in Fig. 5. This figure includes the results of tensile tests obtaining from the fine-grained pure Mg and Mg3
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Fig. 3. TEM (HAADF mode) images of the extruded Mg–Mn ternary alloys; (a) Mg-0.3Mn-0.1Bi and (b) Mg-0.3Mn-0.1Zr alloys. EDS mapping, which is measured in area enclosed with white dashed line using nano-ordered scanning size, is inset in each figure.
Fig. 4. Nominal stress vs. strain curves obtained from tensile tests at room-temperature; (a) Mg-0.1Bi, (b) Mg-0.1Zr, (c) Mg-0.3Mn-0.1Bi and (d) Mg-0.3Mn-0.1Zr alloys.
contribution to deformation during superplastic flow, due to enhancement of diffusion. The m-value is high and is theoretically obtained to be 0.5 under such conditions. In order to examine this point, the results of tensile strain rate change tests at intermediated temperatures are shown in Fig. 7 for the ternary alloys. Stress vs. strain behavior is shown up to a strain of 1.5 for both figures. It is difficult to identify the actual elongation-to-failure in tension, because strain rate change tests were performed. However, the fracture strain is more than 2.0 at the temperatures of 373 K and 423 K. The strain rate change is fixed to be from 1 × 10−5 to 1 × 10−4/s, to eliminate the consideration for effect of the strain rate on the m-value. In Fig. 7, tensile behavior is influenced by the testing temperature; the flow stress decreases with increase in testing temperature for both alloys. The m-values obtained in the strain rates of 10−5–10−4/s at each temperature are summarized in Table 2. The m-values exist in between 0.17 and 0.20 at the intermediated
temperature ranges of 323 K and 423 K. It is interesting to note that the m-value generally increases with higher temperature, which suggests an increase in activation of GBS. However, these values are close to those obtained at room temperature. One of the reasons is assumed to be that the testing temperature is still low to affect major deformation via diffusivity. The other reason is that unique grain boundary structure in ternary alloys is influential for prevention of diffusivity, as described in later section. In any case, the rate-controlling deformation is likely to be the same for testing at temperatures between 298 K and 423 K. That is to say, in the present alloys, GBS occurs and its contribution to deformation is similar in these temperature ranges. As for the internal factors, the grain size and misorientation angle are critical for deformation mechanism and m-value. Grain refinement and/or inducing a high fraction of high-angle grain boundaries lead to not only a large contribution of GBS to deformation but also an increase in m-value. 4
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Fig. 5. The variation of (a) elongation-to-failure in tension and (b) flow stress as a function of strain rate of the extruded Mg and its alloys. The results of fine-grained Mg and Mg-0.3Mn alloy produced by extrusion are cited from previous studies [28,52].
In the Mg–Mn binary alloys, high-angle grain boundaries are confirmed to readily become segregation sites for Mn, which play a part in enhancement of GBS [28,51]. In addition, a Mg–Mn alloy having a dense fraction of high-angle grain boundaries exhibits a higher m-value [51]. However, in this study, these internal factors can be ruled out as important factors, because the grain size and misorientation angle of all the extruded alloys are similar, as shown in Fig. 1 and Table 1. From the atomistic viewpoints, grain boundary structure is also realized as affecting the deformation behavior [36]. The microstructures in the vicinity of grain boundaries observed by HREM are shown in Fig. 8 for the ternary alloys. Grain boundary structures of both alloys show existence of strains, such as dislocations, as well as atomic-level ordered facets and steps. These microstructural features display that these alloys have non-equilibrium grain boundaries. Our previous studies have interestingly reported that the extruded Mg–Bi binary alloys have equilibrium grain boundaries, regardless of the chemical composition in Bi [38,39]. This is inconsistent with the present result; the combinational addition of Mn and Bi is found to form a different grain boundary structure. Numerical analysis has revealed that Mn element is one of the favorable elements to segregate at grain boundaries among other elements, owing to large grain boundary segregation energy [38]. Grain boundary segregation behavior, as shown in Fig. 3, is assumed to induce atomistic strain, i.e., the origin or site for non-equilibrium grain boundary structure formation during recrystallization in the wrought process. Therefore, it is again noted that such a grain boundary structure has the possibility to affect the deformation behavior. Comparing the present results for binary and ternary alloys, the ternary alloys containing Mn show lower m-value than the binary alloys, which display their smaller contribution of GBS on deformation. This is closely related to the fact that the non-equilibrium grain boundary structure has a characteristic of preventing dislocation climb and slip; thus, these grain boundaries are likely to be impediment for GBS. Similar to the results of m-value, the elongation-to-failure of the ternary alloys is lower than that of the binary alloys.
XRD measurement and TEM observation, as shown in Figs. 2 and 3, show that the ternary alloys have a dense dispersion of precipitates. When GBS occurs, cavitation becomes readily the origin of fracture, where cavity nucleation site is well-observed at grain boundary triple junction and at the interface between the matrix and particle [28,32,54–57]. The same behavior as that for room temperature cavity formation is confirmed in fine-grained Mg-0.3Mn alloys [28]. Not only a non-equilibrium grain boundary structure but also a dense precipitation leads to a decrease in elongation-to-failure in tension. Compared with conventional fine-grained Mg alloys, e.g., AZ31 alloy, it again says that the present ternary alloys apparently show a larger elongation-to-failure and a higher m-value. The Mg–Mn based alloys consist of non-equilibrium grain boundary structure; on the other hand, one of the reasons for obtaining good properties are grain boundary segregation of Mn, which is the most effective solute element to enhance GBS among the other solute elements [30]. Grain boundary segregation behavior is an important means of microstructural control for the Mg–Mn system alloys, as shown in Fig. 3. For the binary alloys, although such grain boundary structures (non-equilibrium grain boundaries and alloying element segregated grain boundaries) are unlikely to form [38,39], they have superior ductility via GBS to that of the present ternary alloys, in particular at lower strain rate regimes. The equilibrium grain boundaries are assumed to have a possibility to promote grain boundary diffusion as compared to that of unique grain boundaries. However, it is necessary to examine this point in detail in the future. Finally, we compare the stress vs. strain behavior between the present and previous results showing huge elongation-to-failure. The constitutive equation for metallic materials is generally given as follows [54,57]: ·
=A
Gb kT
n
G
b d
p
D
(1)
Fig. 6. The results of surface observation by SEM; (a) Mg-0.3Mn-0.1Bi and (b) Mg-0.3Mn-0.1Zr alloys. White arrows indicate grain boundary sliding and horizontal direction is tensile direction, which is parallel to extrusion direction. 5
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Fig. 7. The results obtained from strain rate change tests at intermediated temperature ranges; (a) Mg-0.3Mn-0.1Bi and (b) Mg-0.3Mn-0.1Zr alloys. Tensile strain rate is changed at strain of 0.1 in this testing. Table 2 The m-values obtained from strain rate change testing at intermediate temperature ranges for the extruded Mg–Mn system ternary alloys. m-value
Mg-0.3Mn-0.1Bi Mg-0.3Mn-0.1Zr
323 K
373 K
423 K
0.20 0.17
0.19 0.18
0.18 0.18
·
where is the strain rate, σ is the flow stress, T is the temperature, G is the shear modulus, b is the Burgers vector, d is the grain size, p is the grain size exponent, D is the diffusion rate and A, k and R are the well· known constants. The variation in (kT/Gb) (d/b)2 as a function of (σ/ G) is shown in Fig. 9. Although the grain size exponent, p, and the parameter of D are still being debated, they are cited to be p = 2 and the grain boundary diffusion coefficients, respectively, on the basis of previous reports [52,58]. This figure also includes the previous data obtained from several testing methods, such as compressive and hardness tests, using fine-grained Mg and its alloys [29,30,39,52,58–60]. When GBS is the rate-controlling mechanism, the slope in this figure, corresponding to the n-value, would be 2 [54–57]. However, the gradient of the present fitting line is larger than 2, which suggests that GBS “partially” contributes to deformation. It is noted that all the present experimental data normalized by several parameters are in good agreement with the previous results. This indicates that, in the finegrained Mg and its alloys, deformation behavior to attain good property for ductility is unlikely to be influenced by the testing method and experimental conditions, such as the temperature and strain rate; occuring the same deformation mechanism.
·
Fig. 9. The variation in /D(kT/Gb) (b/d)2 as a function of (σ/G) in finegrained magnesium and its alloys [29,30,39,52,58–60]. The stress vs. strain behavior is evaluated by tensile, compressive and hardness testing methods.
4. Conclusion The effect of the combined addition of Mn and micro-alloying elements, such as Bi and Zr, on tensile properties and deformation behavior was investigated on extruded Mg dilute binary alloys and Mg–Mn
Fig. 8. Microstructures of vicinity in grain boundary observed by HREM; (a) Mg-0.3Mn-0.1Bi and (b) Mg-0.3Mn-0.1Zr alloys. White arrow in each image indicates grain boundary. 6
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based ternary alloys. The following conclusions were made.
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(1) All the extruded alloys had the fine-grained structures, i.e., an average grain size of approximately 1 μm, and similar average misorientation angles. However, the extrusion temperature of ternary alloys was higher than that of binary alloys to obtain the same microstructures. The microstructural observations also revealed that the alloys containing Mn and/or Bi elements showed the existence of precipitates in the matrix. Segregation of Mn at grain boundaries was confirmed in the Mg–Mn ternary alloys, regardless of the types of micro-alloying elements. (2) The results obtained from room temperature tensile tests displayed that all the extruded alloys had good properties of ductility. The elongation-to-failure in tension of more than 50% was obtained even at quasi-static strain rate regimes. In the Mg binary alloys, the elements of Bi and Zr are found to be alternatives to Mn to maintain good ductility. The elongation-to-failure increased with lower strain rate, due to the higher contribution of grain boundary sliding to deformation. The micro-alloying elements did not affect the deformation mechanism in the testing temperatures of 298–423 K. (3) The combination of Mn and micro-alloying elements brought about increase in the flow stress, but were unable to attain large elongation-to-failure in tension, in particular in low strain rate regimes. This was closely related to the dense dispersion of precipitates and the formation of non-equilibrium grain boundary structures. Acknowledgement The authors are grateful to Ms. R. Komatsu and Ms. Y. Kobayashi (National Institute for Materials Science) for their technical help. This work was partially supported by the JSPS Grant-in-Aid (C) for Scientific Research in No. 16K06783 and 19K05068. References [1] A. Akhtar, E. Teghtsoonian, Solid solution strengthening of magnesium single crystals-I Alloying behavior in basal slip, Acta Metall. 17 (1969) 1339–1350. [2] A. Akhtar, E. Teghtsoonian, Solid solution strengthening of magnesium single crystals-II the effect of solute on the ease of prismatic slip, Acta Metall. 17 (1969) 1331–1356. [3] H. Yoshinaga, R. Horiuchi, Work hardening characteristics of the basal slip of magnesium single crystals, Mater. Trans., JIM 3 (1962) 220–226. [4] H. Yoshinaga, R. Horiuchi, Deformation mechanisms in magnesium single crystals compressed in the direction parallel to hexagonal axis, Mater. Trans., JIM 4 (1963) 1–8. [5] T. Tsuru, Y. Udagawa, M. Yamaguchi, M. Itakura, Y. Kaji, Solution softening in magnesium alloys; the effect of solid solutions on the dislocation core structure and nonbasal slips, J. Phy. Cond. Mater. 25 (5p) (2013) 022202. [6] J.A. Yasi, L.G. Hctor, D.R. Trinkle, Prediction of thermal cross-slip stress in magnesium alloy from a geometric interaction model, Acta Mater. 60 (2012) 2350–2358. [7] S. Sandlobes, Z. Pei, M. Friak, L.F. Zhu, F. Wang, S. Zaefferer, D. Raabe, J. Neugebauer, Ductility improvement of Mg alloys by solid solution; Ab initio modeling, synthesis and mechanical properties, Acta Mater. 70 (2014) 92–104. [8] H. Somekawa, M. Yamaguchi, Y. Osawa, A. Singh, M. Itakura, T. Tsuru, T. Mukai, Materials design for magnesium alloys with high deformability, Philos. Mag. 95 (2015) 869–885. [9] A. Kula, X. Jia, R.K. Mishra, M. Niewczas, Flow stress and work hardening of Mg-Y alloys, Int. J. Plast. 92 (2017) 96–121. [10] K. Zhang, H. Wen, M.A. Kumar, F. Chen, L. Zhang, I.J. Beyerlein, J.M. Schoenung, S. Mahajan, E.J. Lavernia, Yield symmetry and reduced strength differential in Mg2.5Y alloy, Acta Mater. 120 (2016) 75–85. [11] A. Singh, H. Somekawa, T. Mukai, Dislocation structures in a near-isotropic Mg-Y extruded alloy, Mater. Sci. Eng. A698 (2017) 238–248. [12] S. Sandlobes, S. Zaefferer, I. Schestakow, S. Yi, R. Gonzalez-Martinez, On the role of non-basal deformation mechanism for the ductility of Mg and Mg-Y alloys, Acta Mater. 59 (2011) 429–439. [13] S. Sandlobes, M. Friak, J. Neugebauer, D. Raabe, Basal and non-basal dislocation slip in Mg-Y, Mater. Sci. Eng. A576 (2013) 61–68. [14] Z. Wu, R. Ahmad, B. Yin, S. Sandlobes, W.A. Curtin, Mechanistic origin and prediction of enhanced ductility in magnesium alloys, Science 359 (2018) 447–451. [15] J.J. Bhattacharyya, F. Wang, P.D. Wu, W.R. Whittington, H. El Kadiri, S.R. Agnew, Demonstration of alloying, thermal activation, and latent hardening effects on quasi-static and dynamic polycrystal plasticity of Mg alloy, WE43-T5, plate, Int. J. Plast. 81 (2016) 123–151.
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H. Somekawa, et al. Mg-rare earth alloys, Metal. Mater. Trans. 43A (2012) 1347–1362. [50] M. Bugnet, A. Kula, M. Niewczas, G.A. Botton, Segregation and clustering of solutes at grain boundaries in Mg-rare earth solid solutions, Acta Mater. 79 (2014) 66–73. [51] H. Somekawa, D.A. Basha, A. Singh, Change in dominant deformation mechanism of Mg alloy via grain boundary control, Mater. Sci. Eng. A746 (2019) 162–166. [52] H. Somekawa, T. Mukai, Hall-Petch breakdown in fine-grained pure magnesium at low strain rates, Metall. Mater. Trans. A 46A (2015) 894–902. [53] Method for Evaluation of Tensile Properties of Metallic Superplastic Materials, Japanese Industrial Standards, JIS H-7501, (2016). [54] O.D. Sherby, J. Wadsworth, Superplasticity recent advances and future directions, Prog. Mater. Sci. 33 (1989) 169–221. [55] R.Z. Valiev, T.G. Langdon, Principles of equal-channel angular pressing as a processing tool for grain refinement, Prog. Mater. Sci. 51 (2006) 881–981.
[56] A.P. Zhilyaev, T.G. Langdon, Using high-pressure torsion for metal processing; Fundamentals and applications, Prog. Mater. Sci. 53 (2008) 893–979. [57] R.S. Mishra, T.R. Bicler, A.K. Mukerjee, Superplasticity in powder metallurgy aluminum alloys and composites, Acta Mater. 43 (1995) 877–891. [58] R.B. Figueiredo, S. Sabbaghianrad, A. Giwa, J.R. Greer, T.G. Langdon, Evidence for exceptional low temperature ductility in polycrystalline magnesium processed by severe plastic deformation, Acta Mater. 122 (2017) 322–331. [59] R.B. Figueiredo, F.S.J. Pogglali, C.L.P. Silva, P.R. Cetlin, T.G. Langdon, The influence of grain size and strain rate on the mechanical behavior of pure magnesium, J. Mater. Sci. 51 (2016) 3013–3024. [60] C.H. Cacers, G.E. Mann, J.R. Griffithes, Grain size hardening in Mg and Mg-Zn solid solutions, Metal. Mater. Trans. 42A (2011) 1950–1959.
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Deformation behavior at room temperature ranges of fine-grained Mg–Mn system alloys
T
Hidetoshi Somekawa∗, Dudekula Althaf Basha, Alok Singh Research Center for Structural Materials, National Institute for Materials Science, 1-2-1 Sengen, Tsukuba, Ibaraki, 305-0047, Japan
ARTICLE INFO
ABSTRACT
Keywords: Magnesium Mechanical property Grain boundary sliding Deformation mechanism Grain boundary
Tensile properties and deformation behavior were investigated on Mg dilute binary alloys (Mg-0.1Bi and Mg-0.1Zr alloys) and Mg–Mn based ternary alloys (Mg-0.3Mn-0.1Bi and Mg-0.3Mn-0.1Zr alloys). These alloys were processed by hot extrusion to obtain fine-grained structures with an average grain size of ~1 μm. Microstructural observations showed that Mn element was segregated at grain boundaries in the ternary alloys. Precipitates consisting of the αMn and/or Mg3Bi2 phase were dispersed in the matrix in the alloys containing Mn and/or Bi elements. All the extruded alloys showed good tensile property of ductility. The elongation-to-failure in tension of Mg-0.3Mn-0.1Zr alloy was more than 100% even at a quasi-static strain rate of 1 × 10−3/s. The Mg-0.1Bi alloy exhibited the largest value of 250% among all in the present examined conditions. These values are superior to those of conventional wrought processed Mg alloys. These superior properties are found to be due to the contribution of grain boundary sliding (GBS) to deformation. Bi and Zr are alternative elements to Mn for use in binary alloys. On the other hand, in low strain rate regimes, a combination of Mn and such micro-alloying elements led to an increase in flow stress, but reduced the elongation-to-failure. This resulted from a high density of precipitates in the matrix, which became cavity nucleation sites. Moreover, non-equilibrium grain boundary structures, which induced atomistic strains, were also unlikely to promote GBS. The normalized stress vs. strain relations including the previous results in the temperature ranges of 298–423 K is found to be fitted by the well-known constitutive equation.
1. Introduction Alloying has been recognized as the most effective method to improve the mechanical properties of metallic materials. This approach has also been applied to magnesium (Mg) and its alloys, which have the lowest density among the conventional metallic materials, to overcome the serious issues of poor ductility and formability near room temperature. The most desirable effect of alloying in Mg is to reduce the difference in critical resolved shear stress (CRSS) between the basal and nonbasal planes. Using single crystal Mg alloys, it has been shown that the addition of aluminum (Al), zinc (Zn) or lithium (Li) elements brings about a decrease in the differences in CRSS [1–4]. Recent studies via first-principles calculations have provided that some particular kinds of elements, such as rare-earth (RE) elements and alkali group elements, have the same effect [5–8]. Many observations of deformed microstructures have also shown activation of non-basal dislocation slips in Mg alloys containing RE [9–19] and alkali group alloying elements [8,20–22]. Mg-RE (RE = Y, Gd, Ce etc) alloys exhibit much greater deformability and larger ductility as compared to the conventional Mg alloys [23–26]. Even for the non-RE alloying elements, good properties due
∗
to non-basal dislocation activities are observed in, e.g., Mg–Al–Sn and Mg–Al–Zn–Mn–Ca alloys [22,27]. There are several plastic deformation mechanisms (not only dislocation slip but also diffusion, grain boundary sliding (GBS) and twinning) of Mg and its alloys; however, most previous studies have focused on the control of dislocation slips. Alloying with manganese (Mn) has been reported to improve the room temperature ductility [28–31]. Mn is likely to segregate at grain boundaries during the wrought process. This feature of grain boundary segregation plays a major role in enhancing room temperature GBS. Fine-grained Mg–Mn alloys with a high volume fraction of grain boundaries exhibit an elongation-to-failure in tension of more than 100% even in quasi-static strain rate regimes [28]. These Mg–Mn alloys also have similar stretch formability to that of the conventional Al alloys [29]. These results suggest that GBS is possible and a promising deformation mechanism to resolve the critical problems of poor ductility and formability. A high contribution of GBS on deformation is recognized as an alternative strategy for controlling dislocation slips, such as non-basal slip activities. In order to enhance GBS, grain boundary control, i.e., grain boundary characteristic and structure, is well-known to be an effective method [32–36]. The results obtained from
Corresponding author. E-mail address: [email protected] (H. Somekawa).
https://doi.org/10.1016/j.msea.2019.138384 Received 2 June 2019; Received in revised form 31 August 2019; Accepted 4 September 2019 Available online 04 September 2019 0921-5093/ © 2019 Elsevier B.V. All rights reserved.
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Tensile tests were carried out to investigate the mechanical behavior on all the extruded alloys at room temperature (298 K). The initial strain rates were 1 × 10−2, 10−3, 10−4 and 10−5/s. A strain rate change tensile test was also performed at intermediate temperatures using the Mg-0.3Mn-0.1Bi and Mg-0.3Mn-0.1Zr alloys. The strain rate in tension was changed from 1 × 10−5 to 1 × 10−4/s with a set strain of 0.1. The testing temperatures were selected as 323, 373 and 423 K. Since this test was not performed at room temperature, the holding time for specimen was more than 15 min after reaching the set temperature. The specimens used in all the tensile tests had a gauge length of 10 mm and a diameter of 2.5 mm. Tensile specimens were made by machining from the extruded bars in the direction parallel to the ED. The surface after the tensile test at room temperature at the strain rate of 1 × 10−5/ s was observed by scanning electron microscopy (SEM).
experiments and numerical simulations have pointed out that grain boundaries with higher grain boundary energy accelerate GBS [32,33]. High-angle grain boundaries show larger grain boundary misfits via GBS than that at the special grain boundaries, such as low-angle grain boundaries and coincident site lattice boundaries [34,35]. From the atomic structural viewpoint, equilibrium and non-equilibrium grain boundary structures also affect the temperature and/or the strain rate at which GBS occurs [36]. On application to Mg binary alloys, there are two unique alloying elements: zirconium (Zr) and bismuth (Bi). Firstprinciples calculations have indicated that addition of Zr has an ability to reduce the grain boundary energy [37]. In the case of alloying with Bi, microstructure analysis has revealed that equilibrium grain boundary structures form during the wrought process [38,39]. These two alloying elements have a great potential to enhance GBS. Recently, Xu et al. have investigated the effect of Mn or Zr addition on agehardening response using multi-component system Mg alloy (Mg-Gd-YZn system alloy) [40]; however, to the best of our knowledges, there are no reports about the impact of these alloying elements on the microstructures and mechanical properties of Mg–Mn system ternary alloys. In this study, we have investigated and considered the effect of these micro-alloying elements on the tensile properties and deformation behavior of binary and ternary alloys.
3. Results and discussion The microstructures of extruded alloys observed by EBSD method are shown in Fig. 1. The inverse pole figures show that all the extruded alloys consist of recrystallized grains; most of the grain boundaries are highangle grain boundaries (misorientation angle of more than 15°). The average grain size, d, measured from the EBSD analysis are 0.8–1.2 μm, as provided in Table 1. Although the extrusion temperatures are slightly different, these extruded alloys have similar average grain sizes, because of performing under similar maximum extrusion loads. The binary alloys are able to acquire fine-grained structures at lower extrusion temperatures as compared to those of the ternary alloys. This table also includes the average misorientation angles, θ, of 40°–50° for all the extruded alloys. They are found to be independent of the alloying element. It is also found that these alloys have a basal texture, owing to the extrusion process. The XRD plots are shown in Fig. 2. Apart from the α-Mg, there are no other peaks in the Mg-0.1Zr alloy. In contrast, the alloys containing Bi elements (Mg-0.1Bi and Mg-0.3Mn-0.1Bi alloys) also show the binary Mg3Bi2 phase. In the extruded Mg–Mn–Bi ternary alloy, the existence of α-Mn peak is also confirmed. Although the maximum solubility of Bi and Mn in Mg are 1.13 at.% and 1.0 at%, respectively [41], which are significantly larger than the present chemical composition (0.1 and 0.3 at.%), the solubility of both elements dramatically decrease at the extrusion temperature (388–423 K, as listed in Table 1). Hence, the Mg3Bi2 phase and α-Mn precipitate during the wrought process. The same microstructural features are observed in the previously reported extruded Mg-Bi [38,39] and Mg–Mn binary alloys [28,42]. Typical microstructures of extruded ternary alloys observed by TEM are shown in Fig. 3. These are high-angle annular dark-field (HAADF) images, in which brighter contrast indicates concentration of heavier elements. Some of bright regions, which display the aggregation of elements with large atomic number (the atomic numbers of Mg, Mn, Zr and Bi are 12, 25, 40 and 83, respectively [43]), can be confirmed in both images. They confirm the existence of precipitates formed by these heavy elements, that is agreement with the XRD results (Fig. 2). Fig. 3 also includes the results of energy dispersive X-ray spectroscopy (EDX) analysis in the square area enclosed by a white dashed line. These EDX maps reveal a higher Mn concentration at grain boundaries in both the extruded ternary alloys. Thus, Mn is found to segregate at grain boundaries. Such grain boundary segregation behavior is reported in several wrought processed Mg alloys [9,19,33,44–50] and extruded Mg–Mn binary alloys [28,30]. While the results of TEM observations of the binary alloys are not shown here, segregation of alloying elements at grain boundaries are not observed. This non-segregation grain boundary behavior is the same as that of previous reports [38,39]. The nominal stress vs. nominal strain curves of all the extruded alloys in tensile tests at various strain rates are shown in Fig. 4. All the plots have the same x- and y-axis scale for ease of comparison. Strain rate affects the tensile behavior in all the extruded alloys. The flow stress decreases, but the elongation-to-failure in tension increases at lower strain rate, irrespective of the alloying elements. Even at room
2. Experimental procedure Mg-0.1 at.%X binary alloys and Mg-0.3 at.%Mn-0.1 at.%X ternary alloys (X = Bi or Zr) were used in this study, hereafter denoted as the Mg-0.1Bi, Mg-0.1Zr, Mg-0.3Mn-0.1Bi and Mg-0.3Mn-0.1Zr alloys. Except for Mg-0.1Bi alloy, these alloys were cast by combining a master alloy (Mg-2.6 at.%Mn or Mg-1.5 at.%Zr alloy) with commercial purity Mg. These cast alloys were solution-treated at the temperature of 773 K for 2 h for binary alloys and 24 h for ternary alloys. The solution-treated alloys were subsequently extruded at an extrusion ratio of 25:1 into 8 mm diameter rods. The billet was kept in furnace at the set extrusion temperature for more than 0.5 h, and then was extruded at a speed of 0.2 mm/s. The extrusion temperatures of these alloys are listed in Table 1. The reason for different extrusion temperatures was the limitation of maximum load in the extrusion machine. The phases of all the extruded alloys were determined by X-ray diffraction (XRD) with Cu-Kα radiation. The microstructures of extruded alloys were observed by an electron backscatter diffraction (EBSD, in a field emission gun scanning electron microscope (FE-SEM) equipped with an EDAX-TSL EBSD system) and by transmission electron microscopy (TEM) and highresolution electron microscopy (HREM). The scanning step size for EBSD measurement was 75 nm, and the EBSD results were analyzed using the EDAX software. The area measured and observed by XRD, EBSD, TEM and HREM was at the center of the cross-section containing the extrusion-direction and the transverse-direction (ED-TD). After embedding each alloy in resin, the samples for EBSD observation were prepared by grinding under diamond slurries and colloidal silica slurry, and then were etched using acetic solution for several seconds. The specimens for XRD measurement were prepared by hand polishing to mirror surface. TEM and HREM samples were prepared by ion-milling. Table 1 The experimental condition and results obtained from microstructural analysis and room temperature tensile testing. Text, K
Mg-0.1Bi Mg-0.1Zr Mg-0.3Mn-0.1Bi Mg-0.3Mn-0.1Zr
388 393 423 403
d, μm
0.8 1.0 1.1 1.2
θ, °
44.3 40.9 50.9 47.8
m-value 10−4–10−5/s
10−2–10−3/s
0.25 0.22 0.21 0.18
0.10 0.10 0.12 0.12
where Text is the extrusion temperature, d is the average grain size and θ is the average misorientation angle. 2
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Fig. 1. Initial microstructures observed by EBSD method in extrusion-transverse direction plane; (a) Mg-0.1Bi, (b) Mg-0.1Zr, (c) Mg-0.3Mn-0.1Bi and (d) Mg-0.3Mn0.1Zr alloys. The scale bars in these images are 20 μm.
0.3 at.%Mn binary alloy (hereafter denoted as the Mg-0.3Mn alloy) processed by extrusion [28,52], which have average grain sizes of 1.2 and 1.8 μm, respectively, measured from the EBSD analysis. The flow stress in Fig. 5(b) is cited from stress at a strain of 0.1, in accordance with the Japanese Industrial Standard Method in Ref. [53]. Fig. 5(a) clearly shows the effect of micro-alloying on elongation-to-failure in tension. The binary alloys are superior in ductility compared to pure Mg and Mg-0.3Mn alloy. Focusing on the binary alloys, to attain and to retain good property of ductility, Zr and Bi elements are found to be alternatives to Mn. As for the ternary alloys, the elongation-to-failure in tension is similar to that of the Mg-0.3Mn alloy in the strain rate regimes of 10−2–10−5/s. The elongation-to-failure of the ternary alloys is also larger than that of the present binary alloys at the strain rates of 10−2–10−3/s; nevertheless, the superiority is unlikely to be obtained in the low strain rate range (1 × 10−5/s). The combined addition of Mn and these micro-alloying elements brings about an increase in the flow stress, and is effective for improving ductility in the quasi-static strain rate regimes. Fig. 5(b) shows the dependence of the flow stress on the strain rate. The slope in this figure corresponds to the strain rate sensitivity, i.e., the m-value, which has a close relation to the plastic deformation mechanism [54,55]. The m-value for each alloy is calculated by dividing into two different strain rate regimes, as summarized in Table 1. The m-value decreases with increase in strain rate, and the m-values of the binary alloys are slightly higher than those of the ternary alloys. Nevertheless, all the extruded alloys have much larger m-values than the wrought processed conventional Mg alloys. The m-values of the extruded AZ31 alloy with fine-grained structures are reported to be 0.01 and 0.03 at the quasi-static and low strain rates, respectively [51], where deformation mechanism is dominantly dislocation slip under tensile stress state. This reveals that the plastic deformation mechanism of the present extruded alloys is different from that of wrought processed common Mg alloys. It is known that the contribution of GBS to deformation becomes large with increase in the m-value [54–56]. Typical surface features after tensile tests at room temperature are shown in Fig. 6 for the ternary alloys. The tensile direction is the horizontal in both images. These images apparently show traces of GBS, marked by white arrows. The same feature of GBS is also observed on the surface after room temperature tensile tests in the fine-grained Mg-0.3Bi alloy [39] and the fine-grained Mg-0.3Mn binary alloy [30], which both exhibits a high ductility. The present huge elongation-to-failure in tension is assumed to be related to GBS. GBS is well-recognized as the rate-controlling deformation mechanism for superplastic flow [54–57]. This deformation mechanism is affected by external factors (such as testing temperature and strain rate), and internal factors (microstructure). As external factors, a high temperature and a low strain rate cause GBS to make a large
Fig. 2. XRD peaks of the extruded alloys. Blue and red lines correspond to the α-Mn and the Mg3Bi2 phase, respectively. (For interpretation of the references to colour in this figure legend, the reader is referred to the Web version of this article.)
temperature, all the extruded alloys show an elongation-to-failure of more than 50% at a quasi-static strain rate of 1 × 10−3/s. The Mg0.3Mn-0.1Zr alloy exhibits the best ductility of 105% among the present alloys at this strain rate. In addition, the elongation-to-failure of the binary alloys is over 200% at the strain rate of 1 × 10−5/s. The elongation-to-failure of extruded fine-grained conventional Mg alloys, such as Mg–3Al–1Zn (AZ31) alloy, is at the most 30–50% in the strain rate regimes of 10−2–10−5/s [51]. All the present alloys are found to have good ductility properties. This figure also provides that even though all the extruded alloys have similar average grain sizes (as listed in Table 1), the ternary alloys have higher flow stress than that of the binary alloys. The combination of Mn and micro-alloying elements plays a role in increasing the flow stress, mainly due to the presence of precipitates. The variation in elongation-to-failure in tension and flow stress as a function of initial strain rate is shown in Fig. 5. This figure includes the results of tensile tests obtaining from the fine-grained pure Mg and Mg3
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Fig. 3. TEM (HAADF mode) images of the extruded Mg–Mn ternary alloys; (a) Mg-0.3Mn-0.1Bi and (b) Mg-0.3Mn-0.1Zr alloys. EDS mapping, which is measured in area enclosed with white dashed line using nano-ordered scanning size, is inset in each figure.
Fig. 4. Nominal stress vs. strain curves obtained from tensile tests at room-temperature; (a) Mg-0.1Bi, (b) Mg-0.1Zr, (c) Mg-0.3Mn-0.1Bi and (d) Mg-0.3Mn-0.1Zr alloys.
contribution to deformation during superplastic flow, due to enhancement of diffusion. The m-value is high and is theoretically obtained to be 0.5 under such conditions. In order to examine this point, the results of tensile strain rate change tests at intermediated temperatures are shown in Fig. 7 for the ternary alloys. Stress vs. strain behavior is shown up to a strain of 1.5 for both figures. It is difficult to identify the actual elongation-to-failure in tension, because strain rate change tests were performed. However, the fracture strain is more than 2.0 at the temperatures of 373 K and 423 K. The strain rate change is fixed to be from 1 × 10−5 to 1 × 10−4/s, to eliminate the consideration for effect of the strain rate on the m-value. In Fig. 7, tensile behavior is influenced by the testing temperature; the flow stress decreases with increase in testing temperature for both alloys. The m-values obtained in the strain rates of 10−5–10−4/s at each temperature are summarized in Table 2. The m-values exist in between 0.17 and 0.20 at the intermediated
temperature ranges of 323 K and 423 K. It is interesting to note that the m-value generally increases with higher temperature, which suggests an increase in activation of GBS. However, these values are close to those obtained at room temperature. One of the reasons is assumed to be that the testing temperature is still low to affect major deformation via diffusivity. The other reason is that unique grain boundary structure in ternary alloys is influential for prevention of diffusivity, as described in later section. In any case, the rate-controlling deformation is likely to be the same for testing at temperatures between 298 K and 423 K. That is to say, in the present alloys, GBS occurs and its contribution to deformation is similar in these temperature ranges. As for the internal factors, the grain size and misorientation angle are critical for deformation mechanism and m-value. Grain refinement and/or inducing a high fraction of high-angle grain boundaries lead to not only a large contribution of GBS to deformation but also an increase in m-value. 4
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Fig. 5. The variation of (a) elongation-to-failure in tension and (b) flow stress as a function of strain rate of the extruded Mg and its alloys. The results of fine-grained Mg and Mg-0.3Mn alloy produced by extrusion are cited from previous studies [28,52].
In the Mg–Mn binary alloys, high-angle grain boundaries are confirmed to readily become segregation sites for Mn, which play a part in enhancement of GBS [28,51]. In addition, a Mg–Mn alloy having a dense fraction of high-angle grain boundaries exhibits a higher m-value [51]. However, in this study, these internal factors can be ruled out as important factors, because the grain size and misorientation angle of all the extruded alloys are similar, as shown in Fig. 1 and Table 1. From the atomistic viewpoints, grain boundary structure is also realized as affecting the deformation behavior [36]. The microstructures in the vicinity of grain boundaries observed by HREM are shown in Fig. 8 for the ternary alloys. Grain boundary structures of both alloys show existence of strains, such as dislocations, as well as atomic-level ordered facets and steps. These microstructural features display that these alloys have non-equilibrium grain boundaries. Our previous studies have interestingly reported that the extruded Mg–Bi binary alloys have equilibrium grain boundaries, regardless of the chemical composition in Bi [38,39]. This is inconsistent with the present result; the combinational addition of Mn and Bi is found to form a different grain boundary structure. Numerical analysis has revealed that Mn element is one of the favorable elements to segregate at grain boundaries among other elements, owing to large grain boundary segregation energy [38]. Grain boundary segregation behavior, as shown in Fig. 3, is assumed to induce atomistic strain, i.e., the origin or site for non-equilibrium grain boundary structure formation during recrystallization in the wrought process. Therefore, it is again noted that such a grain boundary structure has the possibility to affect the deformation behavior. Comparing the present results for binary and ternary alloys, the ternary alloys containing Mn show lower m-value than the binary alloys, which display their smaller contribution of GBS on deformation. This is closely related to the fact that the non-equilibrium grain boundary structure has a characteristic of preventing dislocation climb and slip; thus, these grain boundaries are likely to be impediment for GBS. Similar to the results of m-value, the elongation-to-failure of the ternary alloys is lower than that of the binary alloys.
XRD measurement and TEM observation, as shown in Figs. 2 and 3, show that the ternary alloys have a dense dispersion of precipitates. When GBS occurs, cavitation becomes readily the origin of fracture, where cavity nucleation site is well-observed at grain boundary triple junction and at the interface between the matrix and particle [28,32,54–57]. The same behavior as that for room temperature cavity formation is confirmed in fine-grained Mg-0.3Mn alloys [28]. Not only a non-equilibrium grain boundary structure but also a dense precipitation leads to a decrease in elongation-to-failure in tension. Compared with conventional fine-grained Mg alloys, e.g., AZ31 alloy, it again says that the present ternary alloys apparently show a larger elongation-to-failure and a higher m-value. The Mg–Mn based alloys consist of non-equilibrium grain boundary structure; on the other hand, one of the reasons for obtaining good properties are grain boundary segregation of Mn, which is the most effective solute element to enhance GBS among the other solute elements [30]. Grain boundary segregation behavior is an important means of microstructural control for the Mg–Mn system alloys, as shown in Fig. 3. For the binary alloys, although such grain boundary structures (non-equilibrium grain boundaries and alloying element segregated grain boundaries) are unlikely to form [38,39], they have superior ductility via GBS to that of the present ternary alloys, in particular at lower strain rate regimes. The equilibrium grain boundaries are assumed to have a possibility to promote grain boundary diffusion as compared to that of unique grain boundaries. However, it is necessary to examine this point in detail in the future. Finally, we compare the stress vs. strain behavior between the present and previous results showing huge elongation-to-failure. The constitutive equation for metallic materials is generally given as follows [54,57]: ·
=A
Gb kT
n
G
b d
p
D
(1)
Fig. 6. The results of surface observation by SEM; (a) Mg-0.3Mn-0.1Bi and (b) Mg-0.3Mn-0.1Zr alloys. White arrows indicate grain boundary sliding and horizontal direction is tensile direction, which is parallel to extrusion direction. 5
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Fig. 7. The results obtained from strain rate change tests at intermediated temperature ranges; (a) Mg-0.3Mn-0.1Bi and (b) Mg-0.3Mn-0.1Zr alloys. Tensile strain rate is changed at strain of 0.1 in this testing. Table 2 The m-values obtained from strain rate change testing at intermediate temperature ranges for the extruded Mg–Mn system ternary alloys. m-value
Mg-0.3Mn-0.1Bi Mg-0.3Mn-0.1Zr
323 K
373 K
423 K
0.20 0.17
0.19 0.18
0.18 0.18
·
where is the strain rate, σ is the flow stress, T is the temperature, G is the shear modulus, b is the Burgers vector, d is the grain size, p is the grain size exponent, D is the diffusion rate and A, k and R are the well· known constants. The variation in (kT/Gb) (d/b)2 as a function of (σ/ G) is shown in Fig. 9. Although the grain size exponent, p, and the parameter of D are still being debated, they are cited to be p = 2 and the grain boundary diffusion coefficients, respectively, on the basis of previous reports [52,58]. This figure also includes the previous data obtained from several testing methods, such as compressive and hardness tests, using fine-grained Mg and its alloys [29,30,39,52,58–60]. When GBS is the rate-controlling mechanism, the slope in this figure, corresponding to the n-value, would be 2 [54–57]. However, the gradient of the present fitting line is larger than 2, which suggests that GBS “partially” contributes to deformation. It is noted that all the present experimental data normalized by several parameters are in good agreement with the previous results. This indicates that, in the finegrained Mg and its alloys, deformation behavior to attain good property for ductility is unlikely to be influenced by the testing method and experimental conditions, such as the temperature and strain rate; occuring the same deformation mechanism.
·
Fig. 9. The variation in /D(kT/Gb) (b/d)2 as a function of (σ/G) in finegrained magnesium and its alloys [29,30,39,52,58–60]. The stress vs. strain behavior is evaluated by tensile, compressive and hardness testing methods.
4. Conclusion The effect of the combined addition of Mn and micro-alloying elements, such as Bi and Zr, on tensile properties and deformation behavior was investigated on extruded Mg dilute binary alloys and Mg–Mn
Fig. 8. Microstructures of vicinity in grain boundary observed by HREM; (a) Mg-0.3Mn-0.1Bi and (b) Mg-0.3Mn-0.1Zr alloys. White arrow in each image indicates grain boundary. 6
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based ternary alloys. The following conclusions were made.
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(1) All the extruded alloys had the fine-grained structures, i.e., an average grain size of approximately 1 μm, and similar average misorientation angles. However, the extrusion temperature of ternary alloys was higher than that of binary alloys to obtain the same microstructures. The microstructural observations also revealed that the alloys containing Mn and/or Bi elements showed the existence of precipitates in the matrix. Segregation of Mn at grain boundaries was confirmed in the Mg–Mn ternary alloys, regardless of the types of micro-alloying elements. (2) The results obtained from room temperature tensile tests displayed that all the extruded alloys had good properties of ductility. The elongation-to-failure in tension of more than 50% was obtained even at quasi-static strain rate regimes. In the Mg binary alloys, the elements of Bi and Zr are found to be alternatives to Mn to maintain good ductility. The elongation-to-failure increased with lower strain rate, due to the higher contribution of grain boundary sliding to deformation. The micro-alloying elements did not affect the deformation mechanism in the testing temperatures of 298–423 K. (3) The combination of Mn and micro-alloying elements brought about increase in the flow stress, but were unable to attain large elongation-to-failure in tension, in particular in low strain rate regimes. This was closely related to the dense dispersion of precipitates and the formation of non-equilibrium grain boundary structures. Acknowledgement The authors are grateful to Ms. R. Komatsu and Ms. Y. Kobayashi (National Institute for Materials Science) for their technical help. This work was partially supported by the JSPS Grant-in-Aid (C) for Scientific Research in No. 16K06783 and 19K05068. References [1] A. Akhtar, E. Teghtsoonian, Solid solution strengthening of magnesium single crystals-I Alloying behavior in basal slip, Acta Metall. 17 (1969) 1339–1350. [2] A. Akhtar, E. Teghtsoonian, Solid solution strengthening of magnesium single crystals-II the effect of solute on the ease of prismatic slip, Acta Metall. 17 (1969) 1331–1356. [3] H. Yoshinaga, R. Horiuchi, Work hardening characteristics of the basal slip of magnesium single crystals, Mater. Trans., JIM 3 (1962) 220–226. [4] H. Yoshinaga, R. Horiuchi, Deformation mechanisms in magnesium single crystals compressed in the direction parallel to hexagonal axis, Mater. Trans., JIM 4 (1963) 1–8. [5] T. Tsuru, Y. Udagawa, M. Yamaguchi, M. Itakura, Y. Kaji, Solution softening in magnesium alloys; the effect of solid solutions on the dislocation core structure and nonbasal slips, J. Phy. Cond. Mater. 25 (5p) (2013) 022202. [6] J.A. Yasi, L.G. Hctor, D.R. Trinkle, Prediction of thermal cross-slip stress in magnesium alloy from a geometric interaction model, Acta Mater. 60 (2012) 2350–2358. [7] S. Sandlobes, Z. Pei, M. Friak, L.F. Zhu, F. Wang, S. Zaefferer, D. Raabe, J. Neugebauer, Ductility improvement of Mg alloys by solid solution; Ab initio modeling, synthesis and mechanical properties, Acta Mater. 70 (2014) 92–104. [8] H. Somekawa, M. Yamaguchi, Y. Osawa, A. Singh, M. Itakura, T. Tsuru, T. Mukai, Materials design for magnesium alloys with high deformability, Philos. Mag. 95 (2015) 869–885. [9] A. Kula, X. Jia, R.K. Mishra, M. Niewczas, Flow stress and work hardening of Mg-Y alloys, Int. J. Plast. 92 (2017) 96–121. [10] K. Zhang, H. Wen, M.A. Kumar, F. Chen, L. Zhang, I.J. Beyerlein, J.M. Schoenung, S. Mahajan, E.J. Lavernia, Yield symmetry and reduced strength differential in Mg2.5Y alloy, Acta Mater. 120 (2016) 75–85. [11] A. Singh, H. Somekawa, T. Mukai, Dislocation structures in a near-isotropic Mg-Y extruded alloy, Mater. Sci. Eng. A698 (2017) 238–248. [12] S. Sandlobes, S. Zaefferer, I. Schestakow, S. Yi, R. Gonzalez-Martinez, On the role of non-basal deformation mechanism for the ductility of Mg and Mg-Y alloys, Acta Mater. 59 (2011) 429–439. [13] S. Sandlobes, M. Friak, J. Neugebauer, D. Raabe, Basal and non-basal dislocation slip in Mg-Y, Mater. Sci. Eng. A576 (2013) 61–68. [14] Z. Wu, R. Ahmad, B. Yin, S. Sandlobes, W.A. Curtin, Mechanistic origin and prediction of enhanced ductility in magnesium alloys, Science 359 (2018) 447–451. [15] J.J. Bhattacharyya, F. Wang, P.D. Wu, W.R. Whittington, H. El Kadiri, S.R. Agnew, Demonstration of alloying, thermal activation, and latent hardening effects on quasi-static and dynamic polycrystal plasticity of Mg alloy, WE43-T5, plate, Int. J. Plast. 81 (2016) 123–151.
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