Chapter 2 Literature Review
2.5 Superplasticity of Magnesium alloys
Superplastic forming (SPF), defined as elongations of at least 100% and strain rate sensitivity close to 0.5, is an effective method to fabricate hard–to–form materials into complex shapes [59,60]. For SPF to be used in industry, the development of high strain–rate superplasticity (HSRSP), defined as superplasticity occurring at strain rates at or above 1.0 × 10-2 s-1 [61], is needed, especially for Mg alloys with poor formability. R.B. Figueiredo et al. [62] proposed two strategies for achieving HSRSP in Mg alloys processed by equal–channel angular extrusion (ECAE): (I) by pressing the alloys through a reduced number of passes in order to increase the thermal stability of the microstructure; and (II) by increasing the processing temperature to permit the occurrence of superplastic flow at higher testing temperatures. Another desirable property for developing superplasticity in a material is low temperature superplasticity (LTSP), defined as superplasticity occurring at temperatures at or below 0.55 Tm, where Tm is the alloy melting temperature [61]. The presence of LTSP is an attractive property in Mg alloys because of their susceptibility to surface oxidation when formed at elevated temperatures and their low formability at temperatures close to RT.
Many previous experiments established that superplasticity requires a small polycrystalline grain size (typically less than 10 μm [63]) and these small grains are generally achieved through the application of SPD. As mentioned in chapter 2.2, ECAE is one of the most popular SPD methods and has proved to be effective in refining grains in various Mg alloys, resulting in improved ductility, strength, and
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superplasticity [14,26,64–71]. K. Matsubara et al. [26] reported that the Mg–9% Al alloy processed by a combination of extrusion and ECAE exhibited a maximum elongation of 840% at 200 °C with a strain rate of 3.3 × 10-4 s-1. R.B. Figueiredo et al.
[64] reported that an ECAE processed ZK60 alloy showed a maximum elongation of 3050% at 200 °C with a strain rate of 1.0 × 10-4 s-1. V.N. Chuvil’deev et al. [65] found that the ECAE processed AZ91 alloy possessed 570% in elongation at 300 °C with a strain rate of 3.0 × 10-3 s-1 and the ECAE processed ZK60 alloy exhibited 810% in as shown in Fig. 2.24. M. Kawasaki et al. [70] also provided a very detailed tabulation of all papers reporting superplasticity in metals processed by the ECAE process.
Superplastic deformation is an integrated process that combines grain boundary sliding (GBS), dislocation movement, and diffusion in intracrystalline. The m value represents the proportion of GBS, and it is well known that high strain rate sensitivity (typically m close to 0.5) is a characteristic of superplastic metals and alloys [72,73].
The stress exponent (n value), which is reciprocal of the m value, is calculated to determine deformation mechanism of the materials. It has been reported that the n value of the GBS and the dislocation creep mechanisms are 2 and 3, respectively [73,74]. Moreover, GBS is usually accommodated by slip controlled by diffusion [75].
To further understand the deformation mechanism of mateirals, the activation energy for the deformation is calculated under constant strain rate using the following formula [76]:
19
) / 1 (
) (ln nR T
Q
where Q is the apparent activation energy, n is the stress exponent, R is the gas constant (R = 8.31 J/(K．mol)), σ is the true stress, and T is the absolute temperature.
Therefore, according to the relationship of lnσ versus 1/T at different strain rates during the deformation, the activation energy Q can be calculated, and then, the deformation mechanism of materials can be determined.
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Figure 2.1. Microstructure of the AS21 alloy [1].
Figure 2.2. Microstructure of the AE42 alloy [4].
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Figure 2.3. Creep properties of the AE42, AE41, AS41, and AZ91 alloys [5].
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Figure 2.4. Ternary phase diagram of Mg–Zn–Al alloy.
Figure 2.5. Creep behaviors between the AZ91 and ZA–series magnesium alloys [9].
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Figure 2.6. Microstructure of the AJ43 alloy [10].
Figure 2.7. Creep behaviors of Mg–Al–Sr and Mg–Al–Sr–Ca alloys [10].
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Figure 2.8. Microstructure of the Mg–Zn–Al–Ca–RE alloy [13].
Figure 2.9 Creep behaviors between As41, AE42, and Mg–Zn–Al–Ca–RE alloys [13].
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Figure 2.10. The cross–sectional figure of the ECAE die [15].
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Figure 2.11. Three types of die–angle combination in the ECAE–die design [16].
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Figure 2.12. The relationship between the amount of accumulated strain and die angles of Φ and Ψ [17].
Figure 2.13. Four types of ECAE routes [18].
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Figure 2.14. Appearance of the 4340 steel subjected to the ECAE process at 350 °C with different pressing rates [24].
Figure 2.15. Appearance of the AZ31 alloy subjected to the ECAE process with different temperatures and pressing rates [25].
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Figure 2.16. Microstructure of the as–cast AZ31 alloy after one ECAE pass, (a) dislocation: b[1210], g[1011], and (b) subgrain: b[1210],
] 0 1 10
[
g [49].
Figure 2.17. Microstructure of the as–cast AZ31 alloy after four ECAE passes [49].
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Figure 2.18. Microstructures of the compressed Mg–2.0 Zn–0.3 Zr– 0.9 Y alloy at 250
°C with different strain rates (a) 0.001 s-1, (b) 0.01 s-1, (c) 0.1 s-1, and (d) 1 s-1 [50].
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Figure 2.19. SEM micrograph of the ZA84 alloy subjected to the SHT for (a) 24 h, (b) 48 h, (c) 72 h, and (d) 100 h [53].
Figure 2.20. Hardness vs. aging time of the ZA84 alloy after the T6 treatment [53].
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Figure 2.21 (a) A comparison of aging characteristics of the dynamically aged and the as–solutionised samples at 170 °C [54].
Figure 2.21 (b) Tensile properties of the 6069 and 6061 alloys after dynamic aging and static peak–aging at 170 °C [54].
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Figure 2.22. Comparison of elongation to failure as a function of deformation temperature in the (a) ZK60 and (b) AZ91 alloys [65].
Figure 2.23. The tensile stress–strain curves of the as–rolled LZ82 alloy [68].
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Figure 2.24. Appearance of the tensile specimens after four ECAE passes at 200 °C and pulling to failure at 200 °C; the upper specimen is untested [69].
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Chapter 3 Effects of Equal–Channel Angular Extrusion on the Microstructure and Tensile Properties of the ZA85 Magnesium Alloy
3.1 Introduction
In recent years, the development of magnesium alloys, which generally have excellent properties such as low density, high specific strength, superior damping capacity, high thermal conductivity, and good electromagnetic shielding characteristics, has been attracting much attention [1–4]. These properties make magnesium alloys suitable for a broad range of applications in electronic devices and the aircraft and automobile industries among others. Among various magnesium alloys, Mg–Al–Zn (AZ) alloys are widely used because of their desirable mechanical properties, corrosion resistance, and castability. However, the application of these alloys is limited at temperatures above 120 °C. This is because their heat resistance is inferior to that of aluminum alloys at high temperatures [5]. This phenomenon is attributed to the presence of intermetallic Mg17Al12 (β–phase), which mainly precipitates along grain boundaries and exhibits a low decomposition temperature of approximately 470 °C. Thus, grain boundary sliding occurs even at temperatures below 150 °C [6,7]. It has been reported that the addition of rare earth (RE) elements to magnesium improves its properties at elevated temperatures [8–10]. However, the use of Mg-RE alloys is limited owing to their inferior ductility and the high cost of RE elements. Another way to improve the high–temperature performance of AZ alloys is to suppress the formation of the β–phase [11]. It has been reported that a ternary addition of a large amount of zinc to binary Mg–Al alloys, with a Zn:Al
41
composition of approximately 2:1, can completely suppress the formation of the β–phase [12,13]. The main precipitate of Mg–Zn–Al (ZA) alloys is Mg32(Al,Zn)49
(τ–phase), which has a higher melting point and decomposition temperature than the β–phase [14]; therefore, ZA alloys exhibit better properties at elevated temperatures compared with commercial AZ alloys.
Another disadvantage of commercial magnesium alloys is their poor formability and low ductility at room temperature (RT) as a result of their hexagonal close-packed (HCP) crystal structure, which limits their practical applications. Microstructure refinement is a promising method to increase the ductility and strength of magnesium alloys. Severe plastic deformation (SPD) has been introduced in materials processing to produce ultrafine-grained microstructures [15]. Equal–channel angular extrusion (ECAE) is one of the most popular SPD methods and can produce a homogeneous submicron or nanocrystalline microstructure in bulk materials [16,17]. A block with two intersecting channels that have identical cross sections is used as an ECAE die.
Severe deformation occurs via simple shear in the zone where the two channels intersect. Large amounts of strain can accumulate by repeated pressing because the channel cross sections are identical. ECAE is proven to be effective in refining grains in various magnesium alloys, resulting in improved ductility, strength, and superplasticity [18–22].
ECAE research on Mg alloys has focused mainly on AZ alloys. The effects of ECAE on ZA alloys, which have better high-temperature properties compared with AZ alloys, have not been investigated yet. Therefore, we investigate the microstructure and tensile properties of the as–cast ZA85 alloy after ECAE via route BC [23] at 180, 220, and 250 °C.
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3.2 Experimental Procedures
The experimental alloy was prepared from commercially pure Mg, Al, and Zn (>99.9%). A steel crucible and an electron resistance furnace were used for melting and alloying with SF6 as the protective atmosphere. Steel molds were used for casting the alloy. The as–cast alloy was air cooled from the molten state. The chemical composition of the experimental alloy was determined by energy dispersive spectroscopy (EDS). The results of chemical analyses were averaged over three different regions that were chosen randomly from the ingot. The chemical composition of the alloy is 8.34 wt.% Zn, 4.74 wt.% Al with the balance Mg, as shown in Table 3.1. For ECAE, specimens of dimensions 17 mm × 17 mm × 60 mm were cut from the ingot, and an ECAE die with an angle of 120° was used. Boron nitride was used as the lubricant during ECAE. The ECAE die was preheated to 180, 220, and 250 °C and maintained for 15 min before inserting a lubricated ECAE specimen into the entrance channel. All specimens were held inside the ECAE die for 5 min before pressing. These specimens were processed via route BC in which after each pass, the specimen was rotated through 90° in the same direction at a pressing speed of 2 mm/min. Microstructures of the as–cast and ECAE materials were examined by standard metallographic procedures. The polished surface was etched with 3 mL acetic acid solution, 5 mL deionized water, 35 mL ethanol, and 1 g picric acid. The microstructures were observed by optical microscopy, scanning electron microscopy (SEM), and transmission electron microscopy (TEM). Regarding the ECAE specimens, the surfaces perpendicular to ND (y–plane) were observed [24].
The average grain size, grain size distribution, and area fraction of grain size were obtained with Image Pro software (IpWin32). A Rockwell indenter with a load of 100 kgf was used for a Rockwell hardness B (HRB) test at RT. The HRB values were averaged over 10 tests under each set of conditions. The ECAE specimens were
43
longitudinally cut to obtain tensile specimens with a gauge section of 6 mm × 3 mm × 2 mm. Tensile tests were conducted at RT and 200 °C with an initial strain rate of 1 × 10−3 s−1 using an Instron 8501 universal testing machine. A furnace mounted on the
The coarse precipitates, identified as the τ–phase by X–ray diffraction (the same as in previous literature [12,14,25]), are distributed along the grain boundaries, as shown in Fig. 3.2. The chemical composition of matrix and τ–phase checked by EDS are shown in Table 3.2. Moreover, several defects such as blow holes and shrinkage voids are clearly observed in the as–cast specimen. These defects result from the air trapped in the melting alloy during casting and from the difference between the cooling rates in the inner and outer regions of the ingot.
Fig. 3.3 shows the optical micrographs of the ECAE–processed specimens after different number of passes were conducted at 180 °C. After fewer than four ECAE passes, all microstructures became inhomogeneous with fine recrystallized grains along the grain boundaries and a number of large distorted grains. This resultant microstructure, termed as “bimodal,” was also observed by Chang et al. [26]. These differ from the microstructure of aluminum alloys subjected to the ECAE process.
Nakashima et al. [27] reported that after two ECAE passes at RT, the microstructure of aluminum alloys became homogeneous with ultrafine grains less than 1 μm in size.
44
This phenomenon is attributed to the difference between the grain refinement mechanisms of magnesium and aluminum alloys. For aluminum alloys, Berbon et al.
[28] proposed a grain refinement mechanism during ECAE deformation. In the first pass, several dislocations are introduced within the grain because of the applied strain.
These dislocations then rearrange into low–energy dislocation structures (LEDSs).
The dislocations generated in the following passes then transform the LEDSs into subgrains. With increasing number of ECAE passes, boundary misorientation would increase to form high–angle grain boundaries. However, the grain refinement mechanism of magnesium alloys by ECAE is mainly dynamic recrystallization [29, 30] because of the relatively few slip systems in HCP metals during the ECAE process at testing temperatures. In HCP metals, the slip system at RT is mainly the basal slip system. At high temperatures, non–basal slip systems such as prismatic and pyramidal slip systems can become activated. However, in this study, ECAE processing temperatures are below 250 °C, which is not high enough to activate all non–basal slip systems [31]. Therefore, dynamic recrystallization is responsible for grain refinement. From the TEM micrograph shown in Fig. 3.4, it can be observed that after six ECAE passes at 180 °C, there are several dislocation–free grains attributable to dynamic recrystallization. The density of dislocation increases owing to the large amount of strain accumulated by the repetition of ECAE processes. Then, dynamic recrystallization occurs in the area of high dislocation density, producing numerous fine grains and reducing dislocation density. It should be noted that the microstructures exist in a preferential orientation after one, two, and six passes but not after four passes. In this study, because ECAE is conducted via route BC, in which the specimen is rotated 90° in the same direction after each pass, the grain structure would return to an equiaxed structure after every four passes.
Fig. 3.5 shows the area fraction of grain sizes for the ECAE–processed
45
specimens. As can be observed in Fig. 3.5(a), the area fraction of fine grains (less than 10 μm) increased with the number of ECAE passes at 180 °C. This indicates that the area fraction of dynamically recrystallized grains progressively increased with strain.
In addition, the area fraction of large grains apparently decreased with subsequent passes. This suggests that ECAE can result in a uniform microstructure. Furthermore, defects such as blow holes and shrinkage voids in the as–cast specimen were eliminated after the ECAE process, as shown in Fig. 3.3.
The SEM micrographs (Fig. 3.6) show the effect of ECAE on the precipitates.
The precipitate size reduced with increasing number of ECAE passes. After four passes, the size was significantly reduced to less than 10 μm. After six passes, it was further reduced to an average of 1 μm with uniform distribution. This proves that ECAE not only reduces grain size but also shatters coarse precipitates. In general, the precipitates shattered by shear stress during the ECAE process ought to have an irregular shape and a rough surface. In this study, the ECAE process is conducted at high temperatures with a low pressing speed; hence, the precipitate surface becomes smoother to reduce surface energy.
Figs. 3.7 and 3.8 show the optical micrographs of the ECAE–processed specimens after different number of passes at 220 and 250 °C, respectively. Compared to the results at the lower ECAE temperature of 180 °C, the degree of dynamic recrystallization increased with ECAE temperature so that a more uniform microstructure was obtained at higher temperatures. Figs. 3.7(c) and 3.8(c) show that the microstructures obtained after four ECAE passes at 220 and 250 °C were more uniform than those at 180 °C. These new fine grains grew slightly during the ECAE process at 220 and 250 °C. This indicates that a higher temperature is more beneficial for dynamic recrystallization. Therefore, the degree of dynamic recrystallzation increases with ECAE temperature, leading to much more uniform microstructures at
46
higher ECAE temperatures with the same number of ECAE passes. The area fraction of grains less than 10 μm in size reached 80% after six passes at 180 °C. On the other hand, most of the dynamically recrystallized grains had grown to 10–20 μm during ECAE at 220 °C, as shown in Fig. 3.5(b). The area fraction of grains of sizes 10–20 μm was approximately 1.5% after one pass and increased to approximately 14.5% and 44% after two and four passes, respectively. After ECAE at 250 °C, grain growth was more evident. In particular, most of the dynamically recrystallized grains had grown to sizes 20–30 μm, as shown in Fig. 3.5(c). Furthermore, the area fraction of grains with sizes 20–30 μm increased to approximately 38% after four passes.
Fig. 3.9 shows the average grain size with different number of ECAE passes at different temperatures. It can be observed that the degree of grain refinement increased with ECAE temperature after one pass. The average grain sizes of the specimens were 31, 19, and 16 μm after one pass at 180, 220, and 250 °C, respectively. However, the grain refinement rate decreased with additional passes at 220 °C. Note that the average grain size increased from 14 μm after two passes to 20 μm after four passes at 250 °C. In contrast, at the lower ECAE temperature of 180 °C, the average grain size reduced to 8 μm after four ECAE passes. It was further reduced to 4 μm after six ECAE passes without conspicuous grain growth.
Fig. 3.10 shows that hardness increases with the number of ECAE passes, which can be attributed to two factors. The first one is grain refinement. As can be observed in Figs. 3.9 and 3.10, a smaller grain size is accompanied by greater hardness. The second factor is the shattered precipitates. These uniformly distributed fine precipitates would hinder grain boundary sliding as well as dislocation slips, which leads to better mechanical properties. After six ECAE passes at 180 °C, hardness increased significantly from HRB 19 to HRB 46.
Fig. 3.11 shows the results of the tensile tests on the ZA85 alloy conducted at RT.
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The trend of the tensile properties is similar to the trend of the hardness properties.
For the as-cast specimen, UTS was 175 MPa at RT. After one ECAE pass at 180, 220, and 250 °C, UTS increased to 187, 198, and 220 MPa, respectively. It can be observed that ECAE temperature varies directly with the strength of the alloy. In addition, UTS and YS increased with additional ECAE passes at 180 and 220 °C but decreased with increasing number of ECAE passes from two to four; this is owing to the grain growth effect at 250 °C. At 180, 220, 250 °C, the UTS of the four–pass ECAE–processed specimens were 373, 348, and 242 MPa, respectively. At 180 °C, the UTS of the six–pass ECAE–processed specimens reached 402 MPa. Compared with those of the as–cast ZA85 alloy, the UTS and YS at RT were improved by up to 230% and 215%, respectively. The tensile properties at RT in the present study are also superior to
For the as-cast specimen, UTS was 175 MPa at RT. After one ECAE pass at 180, 220, and 250 °C, UTS increased to 187, 198, and 220 MPa, respectively. It can be observed that ECAE temperature varies directly with the strength of the alloy. In addition, UTS and YS increased with additional ECAE passes at 180 and 220 °C but decreased with increasing number of ECAE passes from two to four; this is owing to the grain growth effect at 250 °C. At 180, 220, 250 °C, the UTS of the four–pass ECAE–processed specimens were 373, 348, and 242 MPa, respectively. At 180 °C, the UTS of the six–pass ECAE–processed specimens reached 402 MPa. Compared with those of the as–cast ZA85 alloy, the UTS and YS at RT were improved by up to 230% and 215%, respectively. The tensile properties at RT in the present study are also superior to