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DOI: 10.1016/j.jmrt.2019.08.048
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Available online 11 September 2019
Influence of extrusion temperature on dynamic deformation behaviors and mechanical properties of Mg-8Al-0.5Zn-0.2Mn-0.3Ca-0.2Y alloy
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Sang-Hoon Kima, Sang Won Leea, Byoung Gi Moonb, Ha Sik Kimb, Young Min Kimb, Sung Hyuk Parka,
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sh.park@knu.ac.kr

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a School of Materials Science and Engineering, Kyungpook National University, Daegu 41566, Republic of Korea
b Implementation Research Division, Korea Institute of Materials Science, Changwon 51508, Republic of Korea
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Received 19 November 2018. Accepted 26 August 2019
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Table 1. EDS analysis results and predicted intermetallic compounds of homogenized AZXW8000 alloy.
Table 2. Microstructural characteristics and tensile properties of extruded AZXW8000 alloys.
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Abstract

This study demonstrates that the dynamic recrystallization (DRX) and dynamic precipitation (DP) behaviors and resultant mechanical properties of the recently developed nonflammable AZXW8000 alloy vary significantly with the extrusion temperature. With an increase in the extrusion temperature from 250 to 350°C, the DRX behavior of the alloy improves through the transition of dominant recrystallization mechanisms from twinning-induced DRX and continuous DRX to discontinuous DRX, which causes an increase in the DRX fraction of the extruded alloy. During extrusion at 250°C, continuous Mg17Al12 precipitates, which are statically formed during preheating before extrusion, fully dissolve into the matrix and numerous fine Mg17Al12 precipitates are formed dynamically throughout the material. These fine precipitates refine the dynamic recrystallized (DRXed) grains through grain-boundary pinning. As a result, the nonflammable AZXW8000 alloy extruded at 250°C exhibits a high yield strength of 307MPa and an ultimate tensile strength of 398MPa. The tensile strength of the extruded alloy decreases with increasing extrusion temperature, which is attributed primarily to significant weakening of the precipitation hardening and strain hardening effects. However, the tensile ductility improves with increasing extrusion temperature owing to a decrease in the amounts of precipitates and unDRXed grains.

Keywords:
Magnesium
Ca and Y addition
Extrusion
Dynamic recrystallization
Dynamic precipitation.
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1Introduction

Given the low densities and high specific strengths of Mg and its alloys, their use has attracted considerable attention as a possible approach to reducing the weight of automotive components [1–3]. However, the high reactivity of Mg and its alloys causes them to oxidize easily and even ignite at elevated temperatures. This issue poses a safety hazard during the manufacture and use of any such components and limits the applicability of these alloys in many industries. Since the oxide layer of Mg has an inherently loose pore structure, its oxidation and ignition cannot be effectively prevented at high temperatures [4,5]. Alloying can improve these low oxidation and ignition resistances by increasing the density of the loose oxide layer [6–9]. Among the various possible ways to add alloying elements, the combined addition of Ca and Y, instead of their individual addition, has been reported to have strong synergistic effects on the improvement in the oxidation and ignition resistances of Mg alloys [10–12]. In addition, our recent research [13] has demonstrated that the combined addition of 0.5wt% Ca and 0.2wt% Y to an AZ31 alloy leads to the formation of thermally stable second-phase particles, which, in turn, promotes dynamic recrystallization (DRX) during extrusion and consequently improves the tensile strength of the extruded alloy without degrading its tensile ductility.

To date, most Mg components have been predominantly manufactured by casting. However, wrought Mg alloys have great potential for application to automotive components because they can provide high structural safety and enable weight reduction on account of their superior mechanical properties to cast Mg alloys [3,14]. In particular, the extrusion process provides the advantage of facilitating the manufacture of components of various shapes in just a single pass, unlike other metal forming processes such as rolling and forging [15]. During the hot extrusion process, not only the shape of materials but also their microstructure varies greatly. The formation of new fine grains and precipitates can occur during extrusion, for which the DRX and dynamic precipitation (DP) phenomena, respectively, are specifically responsible. DRX occurs through dissipation of the internal strain energy accumulated in the material during extrusion, which results in the formation of small dynamically recrystallized (DRXed) grains. The DP phenomenon occurs when thermal and strain energies are provided simultaneously, and it is particularly promoted during extrusion processes because higher strains are applied in extrusion than in other processes. Therefore, numerous fine precipitates can be formed during hot extrusion in highly alloyed Mg alloys. These DRX and DP behaviors of Mg alloys are strongly dependent on the deformation temperature. Since the microstructural characteristics of extruded Mg alloys are governed by these temperature-dependent DRX and DP behaviors, the influence of extrusion temperature on the microstructural evolution during extrusion and on the mechanical properties resulting from extrusion needs to be investigated. Although extensive studies have been conducted on the effect of extrusion temperature on the microstructure and mechanical properties of various extruded Mg alloys [16–20], no in-depth study has yet been conducted on the variations in the DRX and DP behaviors and the resultant microstructure and mechanical properties with the extrusion temperature in recently developed nonflammable Mg alloys containing Ca and Y. These elements are capable of overcoming the major drawbacks (i.e., high ignitability and low mechanical properties) of Mg alloys. Therefore, in the present study, small amounts of Ca and Y were added to a commercial AZ80 alloy, and then, the microstructural characteristics and tensile properties of the alloy extruded at relatively low and high temperatures (250 and 350°C, respectively) were systematically analyzed with a particular focus on the changes in the DRX and DP behaviors with the extrusion temperature.

2Materials and methods

A Mg-8Al-0.5Zn-0.2Mn-0.3Ca-0.2Y (wt%) (AZXW8000) alloy was used in the present study. The amounts of added Ca and Y were optimized in previous studies in order to improve the ignition resistance [10–12]. To prepare cast billets for direct extrusion, a commercial AZ80 alloy ingot, high-purity (99.9%) Ca, and a Mg-30Y (wt%) master alloy were melted in an electric resistance furnace at 750°C in an inert atmosphere consisting of a CO2–SF6 gas mixture to prevent any oxidation. The molten metals were held at 720°C for 10min to enable their stabilization and then poured into a steel mold preheated to 200°C. The cast billets were homogenized at 445°C for 24h. Cylindrical samples (Ø80mm×200mm) for extrusion were machined from the homogenized billets and preheated at the extrusion temperature (either 250 or 350°C) along with dies having an angle of 90° for 1h prior to extrusion. Direct extrusion was performed at a temperature of 250 or 350°C, a ram speed of 1mm/s, and an extrusion ratio of 8. The extruded rectangular bars were 39.3mm in width and 16mm in thickness. After extrusion, each extrusion butt remaining in the container was immediately water-quenched for microstructural analysis. The AZXW8000 alloys extruded at 250 and 350°C are hereafter referred to as AZXW8000-250 and AZXW8000-350 alloys, respectively.

The microstructural characteristics of the homogenized billet and extruded bars were analyzed by optical microscopy (OM), field-emission scanning electron microscopy (FE-SEM), energy-dispersive X-ray spectroscopy (EDS), X-ray diffraction (XRD), and electron backscatter diffraction (EBSD). The average grain size of the homogenized billets was measured in a 400mm2 area in the OM images by the linear intercept method. The area fraction of DRXed grains and the size of unDRXed grains in the extruded bars were measured in a relatively large area of 122mm2 in the OM images. The size distribution, total number, and area fraction of the undissolved particles in the homogenized billet and the area fraction of the precipitates in each extruded bar were measured in an area of 2.9mm2 in the SEM images by means of an image analyzer (IMT iSolution DT). The solute concentration in the matrix was determined by averaging the EDS measurement results taken at 50 different positions. XRD measurements were performed using a Rigaku D/Max-2500 VL X-ray diffractometer to identify second-phase particles in the homogenized billet and extruded bars. EBSD data were analyzed using the TexSEM Laboratories orientation imaging microscopy (TSL OIM) 7.0 software. The size of DRXed grains, texture, and Schmid factor (SF) for basal slip of the extruded bars were analyzed using EBSD data having confidence index values greater than 0.1. For the tensile test, dogbone-shaped (gage section: Ø6mm×25mm) specimens were machined from each of the extruded bars. The loading axis of the specimens corresponded to the extrusion direction (ED). The tensile test was performed at room temperature by the use of an Instron 4206 universal testing machine at a strain rate of 1.0×10−3s−1.

3Results3.1Microstructural characteristics of homogenized alloy prior to extrusion

Fig. 1(a) and (b) show the optical and SEM microstructures of the homogenized AZXW8000 billet. The average grain size is 515μm, which is larger than that of homogenized Mg-Al-Zn alloy billets without added Ca and Y (150–422μm), which were prepared under the same casting conditions [13,21]. This indicates that the addition of small amounts of Ca and Y results in grain coarsening of a cast AZ80 alloy, which is consistent with previously reported results [13,22]. After homogenization, Mg17Al12 phase particles, which are generally formed in Mg-Al-based alloys along the grain boundaries and cell dendrites during solidification, are fully dissolved in the α-Mg matrix. However, the particles of thermally stable phases remain undissolved, Fig. 1(b). The measured EDS and XRD results reveal that all the undissolved particles contain either the Ca element or the Y element, Fig. 2(a) and Table 1. The polygonal-type Y-containing phases are identified as Al8Mn4Y and Al2Y from the measured EDS and XRD results and calculated equilibrium phase diagram. The presence of these phases is in good agreement with observations of Y-containing Mg-Al-Zn alloys such as AZ91-0.15Ca-0.1Y [23], AZ31-2Y [24], and AZ91D-2Y [25] (wt%). In addition, granular-type Al2Ca phase particles are present in the homogenized billet and the total area fraction of these Ca- or Y-containing undissolved particles is 1.3%, Fig. 1(b). In order to analyze the size distribution of these undissolved particles, the sizes of the particles and the number of particles larger than 1μm were measured, Fig. 1(c). Most particles have sizes of 1–6μm, measured as the equivalent circle diameter, and the number of these particles in an area of 1mm2 is as high as 1573. This means that a large number of particles are present in the alloy even after homogenization heat treatment. The amounts of solute atoms dissolved in the α-Mg matrix after homogenization, which were obtained through EDS measurements, are shown in Fig. 1(d). The amount of Al dissolved in the matrix (6.67wt%) is smaller than that added to the alloy (8.0wt%) and almost no Ca and Y are detected in the matrix. These results indicate that the added Ca and Y elements are almost entirely consumed to form Ca- and Y-containing phases, respectively, and reduce the supersaturation of solute atoms for precipitation [26]. In addition, since all of these phases (i.e., Al2Ca, Al8Mn4Y, and Al2Y, see Fig. 2(a) and Table 1) contain the Al element, their formation causes a decrease in the Al content of the matrix.

Fig. 1.

(a) Optical and (b) SEM micrographs of homogenized AZXW8000 alloy. (c) Particle size distribution and number of particles per unit area of homogenized alloy. (d) Average Al, Zn, Ca, and Y contents dissolved in a-Mg matrix of homogenized alloy, as measured by EDS. davg and fparticle denote the average grain size and the area fraction of undissolved particles, respectively.

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Fig. 2.

XRD results of AZXW8000 alloys (a) homogenized and (b, c) extruded at (b) 250°C and (c) 350°C.

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Table 1.

EDS analysis results and predicted intermetallic compounds of homogenized AZXW8000 alloy.

PointElements (at%)Intermetallic compound
Mg  Al  Zn  Mn  Ca 
17.93  60.40  0.17  –  21.49  –  Al2Ca 
24.43  49.54  –  20.24  0.19  5.60  Al8Mn4
6.72  62.73  0.31  4.51  1.33  24.40  Al2

Prior to the extrusion experiments, the extrusion billets, as well as the container and dies, are preheated up to the extrusion temperature (250 or 350°C). Therefore, static precipitation can occur in the homogenized billet during preheating owing to the supersaturation of the Al atoms in the α-Mg matrix. As shown in Fig. 3(a), both discontinuous and continuous Mg17Al12 precipitates are formed in the billet after preheating at 250°C for 1h. Coarse discontinuous precipitates with a lamellar structure consisting of the Mg17Al12 and α-Mg phases are observed along the grain boundaries, and a number of fine lath-type continuous precipitates are observed inside the grains, Fig. 3(a). In the billet preheated at 350°C for 1h, however, only small amounts of granular-type continuous precipitates and no discontinuous precipitates are observed, Fig. 3(b). Since an increase in the heat treatment temperature causes an increase in the Al solubility in the α-Mg matrix, a much larger amount of Mg17Al12 precipitates is formed in the billet preheated at the lower temperature of 250°C.

Fig. 3.

SEM micrographs of homogenized AZXW8000 alloys preheated for 1h at (a) 250°C and (b) 350°C.

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3.2Microstructural characteristics of extruded alloys

Fig. 4 shows the optical and SEM micrographs of the extruded AZXW8000 alloys. The AZXW8000-250 alloy has a bimodal structure consisting of fine DRXed grains and coarse unDRXed grains, whereas the AZXW8000-350 alloy has an almost fully DRXed grain structure. When the extrusion temperature increases from 250 to 350°C, the DRX fraction increases from 80.5 to 98.1% and the average size of the unDRXed grains decreases slightly from 108 to 97μm. In addition, the unDRXed grains are more severely elongated along the ED in the AZXW8000-250 alloy, which has a lower extrusion temperature. Consequently, the average aspect ratio (length-to-width ratio) of the unDRXed grains in the AZXW8000-250 alloy (1.95) is larger than that of the unDRXed grains in the AZXW8000-350 alloy (1.36), Fig. 4(c) and (d). The homogenized AZXW8000 alloy contains numerous undissolved particles with sizes in the range of 1–6μm, Fig. 1(b). These undissolved particles are fragmented and rearranged along the ED during extrusion, Fig. 4(c) and (d).

Fig. 4.

(a, b) Optical and (c, d) SEM micrographs of AZXW8000 alloys extruded at (a, c) 250°C and (b, d) 350°C. fDRX, dunDRX, and ARunDRX denote the area fraction of DRXed grains, average size of unDRXed grains, and average aspect ratio (length-to-width ratio) of unDRXed grains, respectively.

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Fig. 5 shows the distribution of precipitates in the DRXed and unDRXed regions of the extruded alloys. Although 0.3wt% Ca and 0.2wt% Y are added to the AZ80 alloy, these added elements are almost entirely consumed to form the Ca- or Y-containing undissolved particles, respectively, Fig. 1(d). Hence, all the dynamic precipitates in the extruded alloys are the Mg17Al12 phase formed by the diffusion and reorganization of Al atoms supersaturated in the α-Mg matrix during the extrusion process. This is evidenced from the XRD results of the extruded alloys, Fig. 2(b) and (c). It reveals that XRD peaks of the Mg17Al12 phase are formed after extrusion. In both the AZXW8000-250 and AZXW8000-350 alloys, dynamic Mg17Al12 precipitates with sizes of 0.1–1.1μm are distributed along the grain boundaries and grain interior, Fig. 5. However, the area fractions of the precipitates in both the DRXed and unDRXed regions of the AZXW8000-250 alloy (17.7 and 19.9%, respectively) are considerably higher than those in the AZXW8000-350 alloy (2.3 and 2.1%, respectively). The size of precipitates is also larger in the AZXW8000-250 alloy than in the AZXW8000-350 alloy. This is because the lower Al solubility at the lower extrusion temperature increases the driving force for precipitation of the Mg17Al12 phase, which results in an increase in the amount of Mg17Al12 precipitates formed during extrusion. In addition, extrusion at lower temperatures leads to the formation of more dislocations, which act as nucleation sites for precipitation, than does extrusion at higher temperatures. This consequently promotes the DP behavior during extrusion. Accordingly, more abundant and larger precipitates are formed in the AZXW8000-250 alloy than in the AZXW8000-350 alloy. The difference in the DP behavior between the AZXW8000-250 and AZXW8000-350 alloys and the characteristics of the precipitates are discussed in more detail in Section 4.1.

Fig. 5.

SEM micrographs showing Mg17Al12 precipitates in AZXW8000 alloys extruded at (a, b) 250°C and (c, d) 350°C: (a, c) DRXed region and (b, d) unDRXed region. fppt denotes the area fraction of precipitates.

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The inverse pole figure maps and ED inverse pole figures for the overall, DRXed, and unDRXed regions of the extruded alloys are shown in Fig. 6. With an increase in the extrusion temperature from 250 to 350°C, the average size of the DRXed grains increases from 4.9 to 17.5μm, Fig. 6(b) and (f). Since grain growth aimed at reduction of the grain boundary energy occurs more easily at higher temperatures, the average size of the DRXed grains increases with increasing extrusion temperature. In addition, abundant fine precipitates formed during extrusion in the AZXW8000-250 alloy effectively inhibit the growth of DRXed grains during and after extrusion through grain-boundary pinning, Fig. 5(a).

Fig. 6.

Inverse pole figure maps and ED inverse pole figures of AZXW8000 alloys extruded at (a–d) 250°C and (e–i) 350°C: (a, e) overall region, (b, f) DRXed region, (c, g) unDRXed region, and (d, i) rectangular area in (a) and (e). dDRX and θavg denote the average size of DRXed grains and the average value of misorientation angles between the unDRXed grain and the DRXed grains, respectively.

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Both extruded alloys exhibit a fiber texture that is typical of extruded Mg alloys, wherein the basal planes of most grains are aligned parallel to the ED [27]. The texture intensities of the unDRXed regions of the alloys (5.6 and 9.0 for AZXW8000-250 and AZXW8000-350, respectively) of the alloys are stronger than those of their DRXed regions (1.8 and 3.3 for AZXW8000-250 and AZXW8000-350, respectively), Fig. 6. Furthermore, the texture of the DRXed grains is distributed more widely than is that of the unDRXed grains, which have a strong <10-10> texture. It is known that the formation of DRXed grains during hot deformation causes a rapid reduction in the orientation coherency with parent grains, and the texture of these DRXed grains tends to be relatively weaker than that of the continuously deformed unDRXed grains [28]. This stronger texture of unDRXed grains than of DRXed grains has been observed in various extruded Mg alloys [29,30]. Although the area fraction of the unDRXed grains of the AZXW8000-250 alloy (19.5%) is considerably higher than that of the AZXW8000-350 alloy (1.9%), Fig. 4(a) and (b). The overall texture intensities of these two alloys are almost the same (3.4 and 3.5 for the AZXW8000-250 and AZXW8000-350 alloys, respectively), Fig. 6(a) and (e). The reasons for this similarity are discussed in Section 4.1.

Fig. 7 shows the tensile stress–strain curves of the extruded alloys, and Table 2 lists their corresponding tensile properties. As the extrusion temperature increases from 250 to 350°C, the tensile yield strength (TYS) and ultimate tensile strength (UTS) of the extruded AZXW8000 alloy decrease by 113MPa (from 307 to 194MPa) and by 71MPa (from 398 to 327MPa), respectively. However, the tensile elongation (EL) of the alloy increases significantly, by 140% (from 5.8 to 13.8%). In addition, with the increase in the extrusion temperature from 250 to 350°C, the strain hardening exponent (n value) of the alloy increases from 0.16 to 0.21. The dominant factors causing these variations in the strength, ductility, and strain hardening capability of the extruded alloy with the extrusion temperature are discussed in terms of the strengthening mechanisms, fracture mode, and related microstructural characteristics, respectively, in Section 4.2.

Fig. 7.

Tensile stress–strain curves of extruded AZXW8000 alloys.

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Table 2.

Microstructural characteristics and tensile properties of extruded AZXW8000 alloys.

AlloyMicrostructural characteristicsaTensile propertiesb
fDRX(%)  dDRX(μm)  dunDRX(μm)  fppt-DRX(%)  fppt-unDRX(%)  SFbasal  KAMunDRX  TYS(MPa)  UTS(MPa)  EL(%)  n 
AZXW8000-250  80.5  4.9  108  17.7  19.9  0.19  2.55  307 (± 9)  398 (± 16)  5.8 (±1.2)  0.16 
AZXW8000-350  98.1  17.5  97  2.3  2.1  0.19  1.32  194 (± 7)  327 (± 1)  13.8 (± 1.9)  0.21 
a

fDRX, dDRX, dunDRX, fppt-DRX, fppt-DRX, SFbasal, and KAMunDRX denote the area fraction of DRXed grains, average size of DRXed grains, average size of unDRXed grains, area fraction of precipitates in the DRXed region, area fraction of precipitates in the unDRXed region, average SF value for basal slip under tension along the ED, and average KAM value in the unDRXed region, respectively.

b

TYS, UTS, EL, and n denote the tensile yield strength, ultimate tensile strength, total elongation, and strain hardening exponent, respectively.

4Discussion4.1Dynamic recrystallization and dynamic precipitation behaviors during extrusion at different temperatures

As mentioned in Section 3.2, an increase in the extrusion temperature leads to an increase in the DRX fraction of the extruded alloy. In order to analyze the variations in the DRX behaviors with the extrusion temperature, the microstructures were observed along the metal flow line of the extrusion butts, which refer to the parts of the billet remaining in the container after extrusion, as shown in Fig. 8. Regions A, B, and C in Fig. 8 correspond to the billet region that is somewhat far from the extrusion die, the billet region close to the extrusion die, and the extruded region that has passed through the extrusion die, respectively. The amount of strain imposed in these regions decreases in the following order: region C>region B>region A. In the AZXW8000-250 alloy, at the early stage of extrusion, many twin bands are formed in the coarse initial grains, and DRX occurs along the twin bands and grain boundaries, Fig. 8(c). As the extrusion progresses, the DRXed region further expands toward the interior of the grains, which leads to a gradual reduction in the size and area fraction of the unDRXed grains, Fig. 8(a) and (b). On the other hand, in the AZXW8000-350 alloy, very few twin bands are observed in the initially deformed billet region, and the DRXed grains are formed mainly along the grain boundaries, Fig. 8(f). As the extrusion progresses, the DRXed region rapidly expands toward the interior of the grains, and the extruded material finally has an almost fully DRXed grain structure, Fig. 8(d) and (e).

Fig. 8.

Optical micrographs of extrusion butts of AZXW8000 alloys extruded at (a–c) 250°C and (d–f) 350°C: (a, d) region C, (b, e) region B, and (c, f) region A marked in the low-magnification optical micrograph included on the left side in the figure. Schematic illustrations describing dynamic recrystallization behaviors of AZXW8000 alloys extruded at (g) 250°C and (h) 350°C.

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The DRX mechanism can be categorized as continuous DRX (CDRX), twinning-induced DRX (TDRX), and discontinuous DRX (DDRX). The activation of these DRX mechanisms is known to depend strongly on the deformation temperature. In Mg alloys, CDRX and TDRX are more dominant at low temperatures, and DDRX is more dominant at high temperatures [31,32]. At lower temperatures, where nonuniform deformation occurs, dislocations easily accumulate at the original grain boundaries and high local stress is generated there, which results in cross-slip to non-basal planes and the occurrence of climb [31]. As a result, low-angle grain boundaries are formed in the vicinity of the original grain boundaries through the rearrangement of the dislocations by the promoted cross-slip and climb behaviors, and the formed low-angle grain boundaries are transformed into high-angle grain boundaries upon further deformation. This consequently results in the formation of new grains with distinct boundaries [31,33]. This mechanism is called CDRX, which can occur not only at the original grain boundaries but also at the twin boundaries formed during deformation. In addition, the mutual intersection of twin bands formed during deformation and the formation of transverse low-angle grain boundaries within these twin bands lead to the subdivision of the formed twin bands, which are transformed into high-angle grain boundaries by the CDRX mechanism. This process of formation of new grains in deformation twins is known as TDRX [32]. In Mg alloys, twinning can easily occur during deformation owing to the lack of active slip systems, and the activation of twinning depends on the deformation temperature. At high temperatures above 300°C, non-basal slip systems can be activated on account of the significant reduction in their critical resolved shear stress (CRSS). Therefore, the occurrence of twinning during deformation is suppressed owing to the increased number of active slip systems [34–36]. Therefore, in the AZXW8000-350 alloy extruded at the high temperature of 350°C, no deformation twins are formed during extrusion, Fig. 8(d)–(f). This means that TDRX does not occur at this extrusion temperature. On the other hand, in the AZXW8000-250 alloy extruded at the relatively lower temperature of 250°C, a larger number of narrow twins are formed in the grains during extrusion, which effectively serve as nucleation sites for DRX, Fig. 8(c).

Meanwhile, when Mg alloys are deformed at temperatures higher than ∼300°C, DDRX, which involves grain boundary bulging and subgrain formation by dislocation climb, is predominantly activated instead of CDRX. Thus, new DRXed grains are initially formed along the original grain boundaries [31,32]. In addition, at high extrusion temperatures, the DDRX behavior is accelerated by the increase in grain boundary mobility induced by high thermal energy, and this accelerated DDRX behavior subsequently promotes the formation of new DRXed grains [16,17]. Therefore, at the early stage of extrusion of the AZXW8000-350 alloy, DRX vigorously occurs along the grain boundaries owing to the accelerated DDRX behavior at this relatively high deformation temperature of 350°C, Fig. 8(f). This consequently results in the formation of a typical necklace structure that is involved in DDRX. Because of this enhanced DDRX behavior, the AZXW8000-350 alloy has an almost fully DRXed grain structure, unlike the AZXW8000-250 alloy. In addition, it is known that particles larger than 1μm in size can act as nucleation sites for DRX because of the large strain energy accumulated around them during hot deformation [15,37]. This is the so-called particle-stimulated nucleation (PSN) mechanism, which has been widely observed in various extruded Mg alloys containing second-phase particles larger than 1μm [13,38,39]. As mentioned earlier, a number of undissolved Ca- or Y-containing particles with sizes of 1–6μm are present in the homogenized billet, Fig. 1(b) and (c). DRXed grains are observed in the vicinity of these undissolved particles in both the extruded alloys, which means that these particles result in the formation of DRXed grains during extrusion through the PSN mechanism, Fig. 4(c) and (d). Although the DRX by PSN does not directly lead to a dramatic increase in the DRX fraction, the particles can act as nucleation sites for CDRX or DDRX inside the coarse initial grains. As a result, DRX can occur not only along the grain boundaries but also inside the initial grains, which can lead to a meaningful increase in the DRX fraction of the extruded alloy. Moreover, the resultant size and amount of these DRXed grains formed by PSN vary with the extrusion temperature. The DRXed grains formed by PSN are smaller and more abundant in the AZXW8000-250 alloy than in the AZXW8000-350 alloy, Fig. 4(c) and (d). This is attributed to the variation in the density of dislocations accumulated around the particles with the extrusion temperature. At low extrusion temperatures, dislocations easily accumulate around particles and induce PSN, whereas at high extrusion temperatures, accumulation of dislocations around particles is less pronounced owing to the annihilation of dislocations through enhanced dislocation cross-slip and dislocation climb. Meanwhile, the sizes of the static precipitates formed through preheating and the dynamic precipitates formed during extrusion are not large enough to cause PSN (0.1–0.45μm and 0.1–1.1μm for the static and dynamic precipitates, respectively) [40]. When nanosized particles form a cluster larger than 1μm in size, they can induce PSN during hot extrusion [37,41]. In the present study, however, the nanosized Mg17Al12 precipitates are homogeneously distributed throughout the material without the formation of clusters, Figs. 3 and 5. Therefore, they do not cause PSN during hot extrusion. Consequently, DRX occurs through the TDRX, CDRX, and PSN mechanisms in the AZXW8000-250 alloy, Fig. 8(g). In contrast, it occurs through the DDRX and PSN mechanisms in the AZXW8000-350 alloy, Fig. 8(h). However, the PSN-induced DRX is less pronounced than the DRX caused by the other mechanisms, owing to the low area fraction of the undissolved particles (1.3%).

In terms of the influence of the various DRX mechanisms on the grain texture, the textures of the DRXed grains formed by CDRX [42], TDRX [43], and PSN [37,44] are more random than are those of the DRXed grains formed by DDRX. Therefore, the texture intensity of the DRXed grains of the AZXW8000-250 alloy (1.8), in which TDRX and CDRX are the predominant DRX mechanisms, is relatively lower than that of the AZXW8000-350 alloy (3.3), in which DDRX is the main DRX mechanism, Fig. 6(b) and (f). In order to analyze the difference in the textures of the DRXed grains formed by different DRX mechanisms, misorientation angles between an unDRXed grain and 25 adjacent DRXed grains were measured in both extruded alloys. Here, the misorientation angle refers to the angular difference between the c-axes of the unDRXed and DRXed grains. Fig. 6(d) and (i) shows the distribution of crystallographic orientations of the DRXed and unDRXed grains in the ED inverse pole figure and the average value of the measured misorientation angles, respectively. It is difficult to individually identify DRXed grains formed through each DRX mechanism owing to the subsequent DRX that occurs during extrusion. Therefore, it is assumed that fine grains formed adjacent to a coarse unDRXed grain are DRXed grains formed through CDRX and DDRX in the AZXW8000-250 and AZXW8000-350 alloys, respectively. In both the alloys, the DRXed grains have considerably different orientations from the unDRXed grain. However, as can be seen in the inverse pole figures in Fig. 6(d) and (i), the crystallographic orientations of the DRXed grains are distributed farther away from those of the unDRXed grain in the AZXW8000-250 alloy than in the AZXW8000-350 alloy. As a result, the average misorientation angle in the AZXW8000-250 alloy (38.6°) is higher than that in the AZXW8000-350 alloy (32.4°). This result indicates that the activation of CDRX, which is the main DRX mechanism in the AZXW8000-250 alloy, leads to the formation of DRXed grains with a weaker texture than does that of DDRX, which is the main DRX mechanism in the AZXW8000-350 alloy. As a result, the texture intensity of the DRXed region of the AZXW8000-250 alloy (1.8) is lower than that of the AZXW8000-350 alloy (3.3), Fig. 6(b) and (f). In addition, the AZXW8000-350 alloy has an almost fully DRXed grain structure with a DRX fraction of 98.1% owing to the promoted DDRX behavior, Fig. 4(b). After the completion of DRX during extrusion, lattice rotation of the DRXed grains occurs upon further deformation, which strengthens their basal texture. Accordingly, since the DRXed grains of the AZXW8000-350 alloy are subjected to considerable plastic deformation, they have a stronger basal texture than do the DRXed grains of the AZXW8000-250 alloy, which has a partially DRXed grain structure. Therefore, although the fraction of unDRXed grains in the AZXW8000-350 alloy, which have a stronger texture than the DRXed grains, is smaller than that in the AZXW8000-250 alloy, the overall texture intensity of the former alloy (3.5) is similar to that of the latter (3.4).

As mentioned in Section 3.1, static precipitates are formed along the grain boundaries and in the grains during preheating prior to extrusion, Fig. 3. However, their shape and distribution change greatly after extrusion, Fig. 5. In the extruded alloys, discontinuous precipitates with a lamellar structure are no longer observed at the grain boundaries, and new granular-type precipitates are homogeneously distributed throughout the material, instead of the lath-type continuous precipitates present in the homogenized alloy, Fig. 5. This indicates that not only DRX but also DP occurs in the AZXW8000 alloy during extrusion. To analyze this DP behavior, SEM micrographs of the immediately water-quenched extrusion butts of the two alloys are observed, as shown in Fig. 9. In the AZXW8000-250 alloy, at the early stage of extrusion, the DRXed grains are formed along the initial grain boundaries, which is where statically formed discontinuous precipitates are mainly distributed, and the discontinuous precipitates, which have a lamellar structure, are more severely fragmented in the DRXed region than in the unDRXed region, Fig. 9(a). This severe fragmentation of the precipitates in the DRXed region can be attributed to subgrain rotation by CDRX. This phenomenon is in good agreement with the result observed for an AZ91 alloy subjected to hot compressive deformation [45]. Furthermore, the amount of lath-type continuous precipitates, which are present within the grains of the homogenized billet, decreases and granular-type precipitates are formed in the DRXed and unDRXed regions and the twin bands, Fig. 9(b). This disappearance of the statically formed continuous precipitates and the appearance of the new granular precipitates indicate the simultaneous occurrence of dissolution of static Mg17Al12 precipitates and formation of dynamic Mg17Al12 precipitates during extrusion.

Fig. 9.

SEM micrographs of extrusion butts of AZXW8000 alloys extruded at (a–d) 250°C and (e, f) 350°C: (a, b) region A and (c–f) region C marked in the optical micrograph included on the left side in Fig. 8.

(0.88MB).

Dissolution of the initial precipitates formed during preheating can occur because the temperature in the deformation zone rises considerably owing to the deformation heat and friction heat generated during direct extrusion. Indeed, even under indirect extrusion, wherein there is no friction between the billet and the container wall (unlike direct extrusion), the actual temperature in the deformation zone increases from 250 to above 300°C during extrusion [46]. In addition, the continuous precipitates statically formed during preheating for a relatively short time (1h) may be coherent with the matrix. It is known that fine coherent particles can dissolve when dislocations bypass them during deformation [47]. Therefore, the very fine sized continuous precipitates in the homogenized billet can be easily dissolved during extrusion. At the early stage of extrusion, the continuous precipitates are partially dissolved in the unDRXed region whereas they are completely dissolved in the DRXed region, Fig. 9(b). This can be attributed to the promotion of atomic diffusion via the DRXed grain boundaries. Although the dissolution rates of the static precipitates differ between the DRXed and unDRXed regions, the lath-type continuous precipitates are not observed in either the unDRXed region or the DRXed region after the billet passes through the extrusion die, owing to the sufficient duration and strain for dissolution, Fig. 9(c) and (d).

The dynamic precipitates are formed in both the DRXed and the unDRXed regions and the twin band, Fig. 9(b). However, they differ considerably in terms of their size and amount. The accumulation of a large number of dislocations during extrusion accelerates the formation of dynamic precipitates because the dislocations provide diffusion paths for Al solute atoms and nucleation sites for precipitation. Therefore, the dynamic precipitates formed in the twin band are finer and denser than are those formed in the DRXed and unDRXed regions, because the higher dislocation density in the twin band promotes the diffusion of solute atoms and provides more nucleation sites for precipitation [48]. Furthermore, the density of the precipitates in the DRXed region is lower than that in the unDRXed region and twin band, and the precipitates distributed along the grain boundaries in the DRXed region are coarser than those distributed in the other regions, Fig. 9(b). This presence of less abundant and larger precipitates in the DRXed region is related to the Ostwald ripening phenomenon [49,50]. DRX leads to the formation of new high-angle grain boundaries that facilitate the diffusion of solute atoms. Through grain boundary diffusion, finer precipitates become small and disappear whereas larger precipitates gradually enlarge. As a result, fine precipitates remain inside the grains and coarser precipitates are formed along the grain boundaries in the DRXed region, as shown in Fig. 9(d). The amount of dynamic precipitates in the AZXW8000-350 alloy is significantly smaller than that in the AZXW8000-250 alloy because the higher extrusion temperature increases the solubility of Al and leads to the introduction of fewer dislocations during extrusion, Fig. 9(e) and (f).

These microstructural characteristics resulting from the precipitate dissolution and DP during extrusion may not be retained in the extruded alloys. The above-mentioned DP behaviors are analyzed by examining the microstructures of the extrusion butts, which are water-cooled immediately after extrusion. However, this artificial fast cooling is in contrast to the natural air-cooling of the extruded materials after they pass through the extrusion die. Accordingly, additional static precipitation can occur because of the residual heat inside the extruded material. Indeed, lath-type continuous precipitates are observed in the unDRXed region of the extruded AZXW8000-250 alloy, Fig. 5(b). These precipitates are not present in the extrusion butt, Fig. 9(c). Accordingly, they are concluded to be formed by static precipitation during air cooling after extrusion. In the extruded AZXW8000-350 alloy, irregular-type precipitates, which are not present in the extrusion butt as shown in Fig. 9(f), are observed along the DRXed grain boundaries, Fig. 5(c). These static precipitates, formed partially along the grain boundaries, inhibit the growth of DRXed grains during air cooling after extrusion. As a result, the DRXed grain boundaries in the extrusion butt are relatively straight, Fig. 9(f), whereas those in the extruded alloy are considerably serrated because of the local pinning effect induced by the static precipitates, Fig. 5(c).

4.2Variation in tensile properties with increasing extrusion temperature

The variations in the DRX and DP behaviors caused by the increase in the extrusion temperature result in microstructural differences of the extruded alloy, which, in turn, affects their tensile properties. The TYS and UTS of the extruded alloys decrease with increasing extrusion temperature, which can be explained in terms of various strengthening mechanisms as follows.

  • (1)

    Grain-boundary hardening – With an increase in the extrusion temperature, the average size of the DRXed grains increases from 4.9 to 17.5μm, Fig. 6. However, the area fraction of the coarse unDRXed grains decreases from 19.5 to 1.9% owing to the promoted DRX behavior at the higher temperature, Fig. 4. The average grain size of the overall region consisting of the fine DRXed grains and coarse unDRXed grains, dtot, can be obtained from the individual area fractions and grain sizes of the DRXed and unDRXed regions by utilizing the rule of mixtures as follows.

where fDRX and funDRX are the area fractions of the DRXed and unDRXed regions, respectively, and dDRX and dunDRX are the average grain sizes of the DRXed and unDRXed regions, respectively. From Eq. (1), the average grain sizes of the AZXW8000-250 and AZXW8000-350 alloys are calculated to be 25.0 and 19.0μm, respectively. Although the AZXW8000-250 alloy has significantly finer DRXed grains owing to the grain-boundary pinning effect induced by the abundant precipitates, the calculated average grain size of this alloy is larger than that of the AZXW8000-350 alloy owing to the larger amount of coarse unDRXed grains in the former alloy. Consequently, the increase in the extrusion temperature leads to a slight enhancement of the grain-boundary hardening effect in the extruded AZXW8000 alloy.
  • (2)

    Precipitation hardening and solid-solution hardening – The amount of precipitates in the extruded AZXW8000 alloy decreases significantly with increasing extrusion temperature, Fig. 5. At the higher extrusion temperature, which is unfavorable for precipitation owing to an increase in Al solubility, the inter-precipitate spacing becomes considerably larger. According to the Orowan equation [51], the larger the inter-precipitate spacing, the lower is the resistance to dislocation movement. Therefore, the precipitation hardening effect in the AZXW8000 alloy weakens greatly with increasing extrusion temperature. Meanwhile, as the amount of formed precipitates increases, the amount of Al solute atoms dissolved in the matrix decreases. As a result, the solid-solution hardening effect in the AZXW8000-350 alloy is stronger than that in the AZXW8000-250 alloy. However, in general, the precipitation hardening effect induced by a three-dimensional volume lattice defect is much stronger than the solid-solution hardening effect induced by a zero-dimensional point lattice defect. Previous studies [52,53] have shown that the hardness and strength of Mg-Al-based alloys are substantially higher in the aged state than in the supersaturated state. Therefore, although the solid-solution hardening effect in the extruded AZXW8000 alloy during tensile deformation weakens with increasing extrusion temperature, the precipitation hardening effect in the alloy improves significantly.

  • (3)

    Texture hardening – The maps and distributions of the SF for basal slip under tension along the ED for each extruded alloy are shown in Fig. 10, which reveals that the SF values of the coarse unDRXed grains (0.11 and 0.15 for the AZXW8000-250 and AZXW8000-350 alloys, respectively) are lower than those of the fine DRXed grains (0.26 and 0.19 for the AZXW8000-250 and AZXW8000-350 alloys, respectively). As the basal planes of the unDRXed grains are arranged nearly parallel to the ED, their crystallographic orientation is unfavorable for basal slip under plastic deformation along the ED. In contrast, the basal planes of the DRXed grains are more inclined from the ED than are those of the unDRXed grains, Fig. 6. This results in higher SF values of the DRXed grains. Therefore, the overall SF value of the extruded alloy varies with the DRX fraction. That is, as the area fraction of the unDRXed grains increases, the average SF value of the extruded alloy decreases. However, despite the difference in the DRX fractions of the AZXW8000-250 and AZXW8000-350 alloys, they have the exact same average SF value (0.19), Fig. 10. As described in Section 4.1, the DRXed grains of the AZXW8000-350 alloy exhibit a more intense basal fiber texture than those of the AZXW8000-250 alloy owing to the dominant DDRX mechanism and lattice rotation upon further deformation after complete recrystallization. Accordingly, the SF value of the DRXed grains of the AZXW8000-350 alloy (0.19) is lower than that of the AZXW8000-250 alloy (0.26), Fig. 10. Consequently, the average SF value of the AZXW8000-350 alloy, which is composed of the DRXed grains having a relatively low SF value (0.19), is equal to that of the AZXW8000-250 alloy, which is composed of DRXed grains having a high SF value (0.26) and unDRXed grains having a low SF value (0.11). Therefore, the increased extrusion temperature does not have any influence on the texture hardening effect in the extruded AZXW8000 alloy.

    Fig. 10.

    Schmid factor (SF) maps for basal slip under tension along ED and SF distributions in overall, DRXed, and unDRXed regions of AZXW8000 alloys extruded at (a) 250°C and (b) 350°C.

    (1.35MB).
  • (4)

    Strain hardening – DRX occurs through the elimination of dislocations introduced into the material during extrusion, and the consequently formed new grains have low internal strain. On the other hand, the unDRXed grains, which are continuously deformed during extrusion, have considerably high internal strains owing to the accumulation of dislocations, and this can lead to the strain hardening of the extruded alloy during tensile deformation. In order to analyze the internal residual strain of the unDRXed region of each extruded alloy, the kernel average misorientation (KAM) maps were drawn and the KAM distributions were measured, as shown in Fig. 11. The KAM value represents the local misorientation at each measurement point [54]. The average KAM values of the unDRXed regions of the AZXW8000-250 and AZXW8000-350 alloys are 2.55 and 1.32, respectively, Fig. 11(c). This indicates that the internal residual strain of the unDRXed region decreases with increasing extrusion temperature. As the deformation temperature increases, dislocation cross-slip occurs more easily within the grains owing to the reduction in the CRSS for non-basal slip systems, and dislocation climb is also activated to a greater extent with an increase in the thermal energy [31,55,56]. These dislocation cross-slip and dislocation climb phenomena are known to contribute to the softening of the material [55]. Therefore, the reduction in the internal residual strain with the increase in the extrusion temperature is attributed to the annihilation of dislocations caused by the promoted dislocation cross-slip and climb. Since the two alloys have different amounts of unDRXed grains, the overall internal strain induced by the unDRXed grains can be considered through the product of the KAM value and the area fraction of the unDRXed grains. Both the KAM value and the area fraction of the unDRXed grains of the extruded alloy decrease with increasing extrusion temperature. Therefore, the product of the KAM value and the area fraction of the unDRXed grains decreases significantly from 0.497 for the AZXW8000-250 alloy to 0.025 for the AZXW8000-350 alloy. This indicates that the strain hardening effect in the extruded AZXW8000 alloy during tensile deformation substantially weakens with an increase in the extrusion temperature.

    Fig. 11.

    (a, b) Kernel average misorientation (KAM) maps and (c) KAM distributions in unDRXed region of extruded AZXW8000 alloys: (a) 250°C and (b) 350°C.

    (0.37MB).

Consequently, in the AZXW8000 alloy, as the extrusion temperature increases, the grain-boundary hardening and solid-solution hardening effects of the extruded material strengthen slightly but the texture hardening effect remains unchanged. However, with an increase in the extrusion temperature, the precipitation hardening and strain hardening effects weaken considerably owing to a significant reduction in the amount of precipitates and internal residual strain, respectively. These are the main causes of the drastic decrease in the tensile strength of the extruded AZXW8000 alloy with increasing extrusion temperature.

Although the increase in the extrusion temperature leads to a decrease in the tensile strengths of the extruded AZXW8000 alloy, the tensile ductility of the alloy increases significantly, from 5.8 to 13.8%, with increasing extrusion temperature. In extruded Mg alloys with a bimodal structure consisting of fine DRXed grains and coarse unDRXed grains, the variation in tensile ductility is strongly related to the DRX fraction [16]. In extruded Mg alloys with a typical fiber basal texture, {10-11} contraction and {10-11}-{10-12} double twins are formed during tensile deformation along the ED and these twins are known to act as initiation sites for microcracking [57–59]. In addition, these undesirable twins are formed more easily and frequently in larger grains because the CRSS for twinning decreases with increasing grain size [60]. Accordingly, the twins can be easily formed in coarse unDRXed grains of extruded Mg alloys during tension along the ED. However, in this study, very few deformation twins are formed during tension in the AZXW8000-350 alloy, and no microcracks are observed to have formed along the twins in either the AZXW8000-350 alloy, which has an almost fully DRXed grain structure, or the AZXW8000-250 alloy, which has a considerable amount of coarse unDRXed grains, Fig. 12(a) and (b). This suppressed twinning behavior in the AZXW8000-250 alloy can be attributed to the precipitates formed during extrusion. The extruded AZXW8000-250 alloy has a number of fine Mg17Al12 precipitates, which strengthen the extruded material but make it brittle. Hence, in the AZXW8000-250 alloy, premature fracture occurs before twinning in the coarse unDRXed grains during tensile deformation, which eventually leads to its poor tensile ductility (5.8%). The AZXW8000-350 alloy has a much smaller amount of precipitates than the AZXW8000-250 alloy, Fig. 5. Accordingly, although the precipitation hardening effect in the AZXW8000-350 alloy is weaker than that in the AZXW8000-250 alloy, the degradation of ductility caused by precipitates is not significant in the AZXW8000-350 alloy. In addition, since the AZXW8000-350 alloy has an almost fully DRXed grain structure, the formation of contraction and double twins in the coarse unDRXed grains and the resultant cracking do not occur during tension. For these reasons, the AZXW8000-350 alloy exhibits a much higher EL (13.8%) than the AZXW8000-250 alloy (5.8%). Consequently, the reductions in the amounts of precipitates and unDRXed grains with increasing extrusion temperature improve the ductility of the extruded AZXW8000 alloy.

Fig. 12.

Optical micrographs of fractured tensile specimens of AZXW8000 alloys extruded at (a) 250°C and (b) 350°C.

(0.45MB).

Fig. 13 shows the true tensile stress–strain curves of the extruded alloys, which reveal that the strain hardening capability of the AZXW8000-350 alloy is higher than that of the AZXW8000-250 alloy. The difference between the TYS and the UTS (Δσ) of the AZXW8000-350 alloy (182MPa) is larger than that of the AZXW8000-250 alloy (123MPa). The strain hardening exponent (n value) of the former (0.21) is also higher than that of the latter (0.16). This higher strain hardening capability of the AZXW8000-350 alloy can be attributed to the combined effects of (i) the smaller grain size, (ii) lower area fraction of unDRXed grains, and (ii) larger amount of solute atoms of this alloy [61,62]. Firstly, with regard to grain size, the DRXed grain size of the AZXW8000-250 alloy (4.9μm) is considerably smaller than that of the AZXW8000-350 alloy (17.5μm). However, the average grain size of the AZXW8000-350 alloy (19.0μm) is smaller than that of the AZXW8000-250 alloy (25.0μm) owing to the lower area fraction of coarse unDRXed grains in the former. This indicates that the AZXW8000-350 alloy has more numerous grain boundaries that act as obstacles to the movement of dislocations, which partially contributes to enhancement of the strain hardening effect during the tensile deformation of this alloy. Secondly, the area fraction of the unDRXed grains of the AZXW8000-250 alloy (19.5%) is considerably higher than that of the AZXW8000-350 alloy (1.9%). Since unDRXed grains of extruded Mg alloys have a high dislocation density owing to the continuous deformation that occurs without DRX during extrusion, the strain hardening capability of these grains during subsequent plastic deformation is considerably lower than that of DRXed grains, which have a low dislocation density. Accordingly, the AZXW8000-350 alloy, having an almost fully DRXed grain structure, has higher strain hardening capability during tensile deformation than the AZXW8000-250 alloy, which contains a considerable amount of unDRXed grains (area fraction: 19.5%), Fig. 4(a) and (b). Lastly, an increase in the amount of solute atoms dissolved in a material can suppress dislocation recovery behavior during plastic deformation, which consequently enhances the strain hardening capability of the material [62]. Because of the increased Al solubility at higher extrusion temperatures, the AZXW8000-350 alloy has a larger amount of solute atoms than the AZXW8000-250 alloy. This large amount of solute atoms can also contribute to the increase in the n value of the AZXW8000-350 alloy.

Fig. 13.

True stress–strain curves of extruded AZXW8000 alloys. The n value and Δσ denote the strain-hardening exponent and the difference between the yield strength and the ultimate tensile strength, respectively.

(0.11MB).
5Conclusions

This study demonstrates that the dynamic recrystallization (DRX) and dynamic precipitation (DP) behaviors of the nonflammable AZXW8000 alloy during extrusion vary significantly with the extrusion temperature. The combined addition of small amounts of Ca and Y to the AZ80 alloy causes the formation of thermally stable Al-Ca, Al-Mn-Y, and Al-Y phase particles, which promote DRX during extrusion through the particle-stimulated nucleation (PSN) effect. With an increase in the extrusion temperature from 250 to 350°C, the dominant DRX mechanisms change from twinning-induced DRX (TDRX) and continuous DRX (CDRX) to discontinuous DRX (DDRX), which results in a considerable increase in the area fraction of DRXed grains. During hot extrusion, the dissolution of the statically formed continuous precipitates and the formation of dynamic precipitates occur simultaneously. The alloy extruded at 250°C has a larger amount of fine Mg17Al12 precipitates, whereas the alloy extruded at 350°C contains a relatively smaller amount of precipitates owing to the increased Al solubility and the enhanced dynamic recovery behavior at the higher temperature. The tensile strengths, i.e., tensile yield strength (TYS) and ultimate tensile strength (UTS), of the extruded alloy decrease with an increase in the extrusion temperature, which is attributed mainly to the significant weakening of the precipitation hardening and strain hardening effects. However, the tensile elongation increases with increasing extrusion temperature owing to the decreased amounts of precipitates and unDRXed grains. In addition, with an increase in the extrusion temperature, the strain hardening capability of the extruded alloy increases, which can be attributed to the increases in the grain boundary fraction, DRX fraction, and solute atom amount. When newly developed Mg-Al-based nonflammable alloys containing small amounts of Ca and Y are used as extruded components in the near future, this acquired understanding of the variations in their dynamic deformation behaviors with extrusion temperature will be greatly helpful in controlling the microstructure of the material and in designing the extrusion process suitably to achieve the required mechanical properties.

Conflicts of interest

The authors declare no conflicts of interest.

Acknowledgements

This study was supported by the R&D Center for Valuable Recycling (Global-Top R&BD Program) of the Ministry of Environment of Korea (Project No. 2016002220003).

References
[1]
B.L. Mordike, Ebert T. Magnesium.
Properties — applications — potential.
Mater Sci Eng A, 302 (2001), pp. 37-45
[2]
S. Ramakrishnan, P. Koltun.
Global warming impact of the magnesium produced in China using the Pidgeon process.
Resour Conserv Recycl, 42 (2004), pp. 49-64
[3]
H. Friedrich, S. Schumann.
Research for a “new age of magnesium” in the automotive industry.
J Mater Process Technol, 117 (2001), pp. 276-281
[4]
Y.M. Kim, C.D. Yim, H.S. Kim, B.S. You.
Key factor influencing the ignition resistance of magnesium alloys at elevated temperatures.
Scr Mater, 65 (2011), pp. 958-961
[5]
N. Birks, G.H. Meier, F.S. Pettit.
Introduction to the high temperature oxidation of metals.
2nd ed., Cambridge University Press, (2006),
[6]
B.S. You, W.W. Park, I.S. Chung.
The effect of calcium addition to magnesium on the microstructure and compositional changes of oxide film formed at high temperature.
Mater Trans, 42 (2001), pp. 1139-1141
[7]
X.Q. Zeng, Q.D. Wang, Y.Z. Lü, W.J. Ding, C. Lu, Y.P. Zhu, et al.
Study on ignition proof magnesium alloy with beryllium and rare earth additions.
Scr Mater, 43 (2000), pp. 403-409
[8]
P.Y. Lin, H. Zhou, N. Sun, W.P. Li, C.T. Wang, M.X. Wang, et al.
Influence of cerium addition on the resistance to oxidation of AM50 alloy prepared by rapid solidification.
Corros Sci, 52 (2010), pp. 416-421
[9]
N.V.R. Kumar, J.J. Blandin, M. Suéry, E. Grosjean.
Effect of alloying elements on the ignition resistance of magnesium alloys.
Scr Mater, 49 (2003), pp. 225-230
[10]
Y.M. Kim, H.S. Kim, B.S. You, C.D. Yim.
Non-flammable magnesium alloy with excellent mechanical properties, and preparation method thereof.
(2013),
[11]
Y.M. Kim, B.S. You, M.S. Shim, N.J. Kim.
Mechanical properties and high- temperature oxidation behavior of Mg-Al-Zn-Ca-Y magnesium alloys.
Magnesium technology, pp. 217-219 http://dx.doi.org/10.1007/978-3-319-48203-3_41
[12]
Y. Go, S.M. Jo, S.H. Park, H.S. Kim, B.S. You, Y.M. Kim.
Microstructure and mechanical properties of non-flammable Mg-8Al-0.3Zn-0.1Mn-0.3Ca-0.2Y alloy subjected to low-temperature, low-speed extrusion.
J Alloys Compd, 739 (2018), pp. 69-76
[13]
S.H. Kim, S.W. Bae, S.W. Lee, B.G. Moon, H.S. Kim, Y.M. Kim, et al.
Microstructural evolution and improvement in mechanical properties of extruded AZ31 alloy by combined addition of Ca and Y.
Mater Sci Eng: A., 725 (2018), pp. 309-318
[14]
M.K. Kulekci.
Magnesium and its alloys applications in automotive industry.
Int J Adv Manuf Technol., 39 (2008), pp. 851-865
[15]
M. Bauser, G. Sauer, K. Siegert.
Extrusion.
2nd ed., ASM International, (2006),
[16]
S.H. Park, B.S. You, R.K. Mishra, A.K. Sachdev.
Effects of extrusion parameters on the microstructure and mechanical properties of Mg–Zn–(Mn)–Ce/Gd alloys.
Mater Sci Eng: A, 598 (2014), pp. 396-406
[17]
M. Hirano, M. Yamasaki, K. Hagihara, K. Higashida, Y. Kawamura.
Effect of extrusion parameters on mechanical properties of Mg97Zn1Y2 alloys at room and elevated temperatures.
Mater Trans, 51 (2010), pp. 1640-1647
[18]
M. Shahzad, L. Wagner.
Influence of extrusion parameters on microstructure and texture developments, and their effects on mechanical properties of the magnesium alloy AZ80.
Mater Sci Eng: A, 506 (2009), pp. 141-147
[19]
T. Murai, S. Matsuoka, S. Miyamoto, Y. Oki.
Effects of extrusion conditions on microstructure and mechanical properties of AZ31B magnesium alloy extrusions.
J Mater Process Technol, 41 (2003), pp. 207-212
[20]
A. Singh, M. Watanabe, A. Kato, A.P. Tsai.
Microstructure and strength of quasicrystal containing extruded Mg–Zn–Y alloys for elevated temperature application.
Mater Sci Eng: A, 385 (2004), pp. 382-396
[21]
S.H. Kim, B.S. You, S.H. Park.
Effect of billet diameter on hot extrusion behavior of Mg–Al–Zn alloys and its influence on microstructure and mechanical properties.
J Alloys Compd, 690 (2017), pp. 417-423
[22]
M. Hakamada, A. Watazu, N. Saito, H. Iwasaki.
Tensile properties of forged Mg-Al-Zn-Ca alloy.
Mater Trans, 49 (2008), pp. 554-558
[23]
S.K. Woo, C.D. Yim, Y.M. Kim, B.S. You.
Effect of Ca and Y on corrosion behavior of extruded AZ series Mg alloys.
Magnesium technology, pp. 323-326 http://dx.doi.org/10.1007/978-3-319-48185-2_60
[24]
W. Guobing, P. Xiaodong, L. Junchen, X. Weidong, W. Qunyi.
Structure heredity effect of Mg-10Y master alloy in AZ31 magnesium alloy.
Rare Metal Mat Eng, 42 (2013), pp. 2009-2013
[25]
Z. Zhao, Q. Chen, Y. Wang, D. Shu.
Microstructures and mechanical properties of AZ91D alloys with Y addition.
Mater Sci Eng: A, 515 (2009), pp. 152-161
[26]
B. Langelier, S. Esmaeili.
Effects of Ce additions on the age hardening response of Mg-Zn alloys.
J Miner Mater Charact Eng, 101 (2015), pp. 1-8
[27]
L. Lu, C. Liu, J. Zhao, W. Zeng, Z. Wang.
Modification of grain refinement and texture in AZ31 Mg alloy by a new plastic deformation method.
J Alloys Compd, 628 (2015), pp. 130-134
[28]
D. Ponge, G. Gottstein.
Necklace formation during dynamic recrystallization: mechanisms and impact on flow behavior.
Acta Mater, 46 (1998), pp. 69-80
[29]
Y.Z. Du, X.G. Qiao, M.Y. Zheng, K. Wu, S.W. Xu.
The microstructure, texture and mechanical properties of extruded Mg–5.3Zn–0.2Ca–0.5Ce (wt%) alloy.
Mater Sci Eng: A, 620 (2015), pp. 164-171
[30]
M.G. Jiang, C. Xu, T. Nakata, H. Yan, R.S. Chen, S. Kamado.
Development of dilute Mg–Zn–Ca–Mn alloy with high performance via extrusion.
J Alloys Compd, 668 (2016), pp. 3-21
[31]
A. Galiyev, R. Kaibyshev, G. Gottstein.
Correlation of plastic deformation and dynamic recrystallization in magnesium alloy ZK60.
Acta Mater, 49 (2001), pp. 1199-1207
[32]
O. Sitdikov, R. Kaibyshev.
Dynamic recrystallization in pure magnesium.
Mater Trans, 42 (2001), pp. 1928-1937
[33]
J.C. Tan, M.J. Tan.
Dynamic continuous recrystallization characteristics in two stage deformation of Mg–3Al–1Zn alloy sheet.
Mater Sci Eng: A, 339 (2003), pp. 124-132
[34]
A. Jain, S.R. Agnew.
Modeling the temperature dependent effect of twinning on the behavior of magnesium alloy AZ31B sheet.
Mater Sci Eng: A, 462 (2007), pp. 29-36
[35]
S.E. Ion, F.J. Humphreys, S.H. White.
Dynamic recrystallisation and the development of microstructure during the high temperature deformation of magnesium.
Acta Metall, 30 (1982), pp. 1909-1919
[36]
L. Jiang, J.J. Jonas, A.A. Luo, A.K. Sachdev, S. Godet.
Twinning-induced softening in polycrystalline AM30 Mg alloy at moderate temperatures.
Scr Mater, 54 (2006), pp. 771-775
[37]
J.D. Robson, D.T. Henry, B. Davis.
Particle effects on recrystallization in magnesium–manganese alloys: particle-stimulated nucleation.
Acta Mater, 57 (2009), pp. 2739-2747
[38]
S.H. Park, H. Yu, J.H. Bae, C.D. Yim, B.S. You.
Microstructural evolution of indirect-extruded ZK60 alloy by adding Ce.
J Alloys Compd, 545 (2012), pp. 139-143
[39]
Z.T. Li, X.D. Zhang, M.Y. Zheng, X.G. Qiao, K. Wu, C. Xu, et al.
Effect of Ca/Al ratio on microstructure and mechanical properties of Mg-Al-Ca-Mn alloys.
Mater Sci En: A, 682 (2017), pp. 423-432
[40]
J.D. Robson, D.T. Henry, B. Davis.
Particle effects on recrystallization in magnesium–manganese alloys: particles pinning.
Mater Sci Eng: A, 528 (2011), pp. 4239-4247
[41]
H. Yu, Y.M. Kim, B.S. You, H.S. Yu, S.H. Park.
Effects of cerium addition on the microstructure, mechanical properties and hot workability of ZK60 alloy.
Mater Sci Eng: A, 559 (2013), pp. 798-807
[42]
H. Borkar, R. Gauvin, M. Pekguleryuz.
Effect of extrusion temperature on texture evolution and recrystallization in extruded Mg–1% Mn and Mg–1% Mn–1.6%Sr alloys.
J Alloys Compd, 555 (2013), pp. 219-224
[43]
K.D. Molodov, T. Al-Samman, D.A. Molodov, G. Gottstein.
Mechanisms of exceptional ductility of magnesium single crystal during deformation at room temperature: multiple twinning and dynamic recrystallization.
Acta Mater, 76 (2014), pp. 314-330
[44]
M. Hradilová, F. Montheillet, A. Fraczkiewicz, C. Desrayaud, P. Lejček.
Effect of Ca-addition on dynamic recrystallization of Mg–Zn alloy during hot deformation.
Mater Sci Eng: A, 580 (2013), pp. 217-226
[45]
S.W. Xu, N. Matsumoto, S. Kamado, T. Honma, Y. Kojima.
Dynamic microstructural changes in Mg–9Al–1Zn alloy during hot compression.
Scr Mater, 61 (2009), pp. 249-252
[46]
J.G. Jung, S.H. Park, B.S. You.
Effect of aging prior to extrusion on the microstructure and mechanical properties of Mg–7Sn–1Al–1Zn alloy.
J Alloys Compd, 627 (2015), pp. 324-332
[47]
F.J. Humphreys, M. Hatherly.
Recrystallization textures. Recrystallization and related annealing phenomena.
2nd edition, Elsevier, (2004), pp. 379-413
[48]
C. Wang, R. Xin, D. Li, B. Song, M. Wu, Q. Liu.
Enhancing the age-hardening response of rolled AZ80 alloy by pre-twinning deformation.
Mater Sci Eng: A, 680 (2017), pp. 152-156
[49]
W. Ostwald.
Lehrbuch Der allgemeinen chemie.
Engelmann, (1887),
[50]
J. Li, C. Guo, Y. Ma, Z. Wang, J. Wang.
Effect of initial particle size distribution on the dynamics of transient Ostwald ripening: a phase field study.
Acta Mater, 90 (2015), pp. 10-26
[51]
E. Orowan.
Symposium on internal stresses in metals and alloys.
Institution of mechanical engineers, organizer, (1948), pp. 70-71
[52]
W.J. Lai, Y.Y. Li, Y.F. Hsu, S. Trong, W.H. Wang.
Aging behavior and precipitate morphologies in Mg–7.7Al–0.5Zn–0.3Mn (wt.%) alloy.
J Alloy Compd, 476 (2009), pp. 118-124
[53]
S.H. Kim, J.U. Lee, Y.J. Kim, J.H. Bae, B.S. You, S.H. Park.
Accelerated precipitation behavior of cast Mg-Al-Zn alloy by grain refinement.
J Mater Sci Technol, 34 (2018), pp. 265-276
[54]
S.I. Wright, M.M. Nowell, D.P. Field.
A review of strain analysis using Electron backscatter diffraction.
Microsc Microanal, 17 (2011), pp. 316-329
[55]
Z. Trojanová, P. Lukáč.
Compressive deformation behavior of magnesium alloys.
J Mater Process Technol, 162–163 (2005), pp. 416-421
[56]
G. Bajargan, G. Singh, D. Sivakumar, U. Ramamurty.
Effect of temperature and strain rate on the deformation behavior and microstructure of a homogenized AZ31 magnesium alloy.
Mater Sci Eng: A, 579 (2013), pp. 26-34
[57]
M.R. Barnett.
Twinning and the ductility of magnesium alloys: part II. “Contraction” twins.
Mater Sci Eng: A, 464 (2007), pp. 8-16
[58]
D. Ando, J. Koike, Y. Sutou.
Relationship between deformation twinning and surface step formation in AZ31 magnesium alloys.
Acta Mater, 58 (2010), pp. 4316-4324
[59]
L. Lu, T. Liu, Y. Chen, Z. Wang.
Deformation and fracture behavior of hot extruded Mg alloys AZ31.
J Miner Mater Charact Eng, 67 (2012), pp. 93-100
[60]
A. Ghaderi, M.R. Barnett.
Sensitivity of deformation twinning to grain size in titanium and magnesium.
Acta Mater, 59 (2011), pp. 7824-7839
[61]
X. Chen, F. Pan, J. Mao, J. Wang, D. Zhang, A. Tang, et al.
Effect of heat treatment on strain hardening of ZK60 Mg alloy.
Mater Des, 32 (2011), pp. 1526-1530
[62]
I.A. Yakubtsov, B.J. Diak, C.A. Sager, B. Bhattacharya, W.D. MacDonald, M. Niewczas.
Effect of heat treatment on microstructure and tensile deformation of Mg AZ80 alloy at room temperature.
Mater Sci Eng: A, 496 (2008), pp. 247-255
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