Influence of Bi addition on dynamic recrystallization and precipitation behaviors during hot extrusion of pure Mg

Journal Pre-proof Influence of Bi addition on dynamic recrystallization and precipitation behaviors during hot extrusion of pure Mg Jongbin Go, Jong Un Lee, Hui Yu, Sung Hyuk Park

PII:

S1005-0302(20)30061-X

DOI:

https://doi.org/10.1016/j.jmst.2019.10.036

Reference:

JMST 1928

To appear in:

Journal of Materials Science & Technology

Received Date:

17 August 2019

Revised Date:

11 October 2019

Accepted Date:

22 October 2019

Please cite this article as: Go J, Un Lee J, Yu H, Park SH, Influence of Bi addition on dynamic recrystallization and precipitation behaviors during hot extrusion of pure Mg, Journal of Materials Science and amp; Technology (2020), doi: https://doi.org/10.1016/j.jmst.2019.10.036

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Research Article Influence of Bi addition on dynamic recrystallization and precipitation behaviors during hot extrusion of pure Mg

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Jongbin Goa, Jong Un Leea, Hui Yub, Sung Hyuk Parka,*

School of Materials Science and Engineering, Kyungpook National University, Daegu

41566, Korea b

School of Materials Science and Engineering, Hebei University of Technology, Tianjin

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300130, China

*Corresponding author.

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E-mail address: [email protected] (S.H. Park).

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Abstract

Low material cost and high extrudability for ensuring price competitiveness with Al

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alloys, as well as excellent mechanical properties, are essential for expanding the application range of Mg extrudates. Bi is a promising alloying element for developing

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extruded Mg alloys that satisfy such requirements. Bi is inexpensive, exhibits a high solubility limit, and forms a thermally stable Mg3Bi2 phase, which improves the commercial viability and enables the high-speed extrusion of Mg–Bi alloys. In this study, the effects of Bi addition on the dynamic recrystallization (DRX) and dynamic precipitation behaviors during hot extrusion of a pure Mg and the resultant microstructure and mechanical properties of the extruded materials were investigated. 1

The addition of 6 wt% and 9 wt% Bi to a pure Mg yielded numerous fine Mg3Bi2 precipitates during the early stage of hot extrusion. Consequently, the area fraction of dynamic recrystallized (DRXed) grains decreased because of DRX-behavior suppression by the Zener pinning effect. However, the DRXed grain size was substantially reduced through the grain-boundary pinning effect. The size and number of undissolved Mg3Bi2 particles in the homogenized billets increased when the Bi

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content was increased, which resulted in increased DRX fractions owing to the enhanced levels of particle stimulated nucleation. Bi addition yielded considerable

strength improvement of the extruded pure Mg. However, the extruded Mg–Bi binary

materials were less ductile than the extruded pure Mg material. This lower ductility

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resulted from the cracking at twins formed in the coarse unDRXed grains of the Mg-6Bi

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material and the cracking at large undissolved Mg3Bi2 particles in the Mg-9Bi material.

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Keywords: Magnesium; Bi addition; Extrusion; Recrystallization; Precipitation

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1. Introduction

Interest in the weight reduction of vehicles as a means of reducing carbon dioxide

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emissions and increasing fuel efficiency [1] has recently increased in the transportation industry. In this regard, Mg alloys have attracted attention because of their low density

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and high specific strength [2,3]. Mg alloys are the lightest metals among the structural metal materials and exhibit excellent castability; hence, they have been used predominantly as cast products in the electrical and automobile industries. However, because of the poor mechanical properties of cast Mg alloys, their use has been limited to interior parts subjected only to low loads [4]. Therefore, in recent decades, researchers have focused on wrought Mg alloys, which exhibit considerably greater 2

strength and ductility than cast Mg alloys, with the aim of expanding the application range of Mg alloys [5,6]. Rolling and extrusion are representative metal forming processes for manufacturing wrought materials. Rolled materials (e.g., sheets and plates) are generally produced through repeated operations of hot rolling and intermediate heat treatment, whereas extruded materials (e.g., sheets, plates, rods, beams, and tubes) are manufactured through only single-pass processes.

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Despite the advantages of the extrusion process, the application of extruded Mg alloys is limited by three main issues: mechanical properties, commercial viability, and

corrosion resistance [7]. Regarding mechanical properties, extruded Mg alloys exhibit

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lower strength than extruded Al alloys as well as lower formability and higher tensilecompression yield anisotropy [8]. In Mg with a hexagonal close-packed crystal structure

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and c/a ratio of 1.624, only two independent basal slip systems are predominantly activated during room temperature (RT) plastic deformation. Satisfying the von Mises

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criterion, where at least five independent slip systems are required for homogeneous deformation in polycrystalline metals, is therefore difficult; hence, Mg alloys exhibit

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poor RT ductility and formability. In addition, extruded Mg alloys have a strong basal texture with basal planes aligned parallel to the extrusion direction (ED). This texture

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leads to high tensile-compression yield anisotropy because of the difference in the active deformation modes (i.e., dislocation slip under tension and deformation twinning

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under compression). Commercial viability is related to the maximum extrusion speed at which a sound extrudate can be manufactured without hot cracking. Commercial highstrength Mg alloys such as AZ80 and ZK60 generally have maximum extrusion speeds of 0.5–4.0 m/min, which is substantially lower than that (>40 m/min) of commercial high-strength Al alloys. Therefore, the productivity of extruded Mg alloys is remarkably

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low, resulting in high costs of the final products, which substantially reduces the competitiveness of extruded Mg alloys in the automobile industry [9]. In commercially used Mg–Al and Mg–Zn alloy systems, Mg17Al12 and MgZn2 phases, which have melting temperatures less than 440 °C, are formed during extrusion. The formation of these thermally unstable second phases induces local melting or hot cracking on the surface of the extruded material. The susceptibility of Mg–Al and Mg–

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Zn alloys to hot cracking and the maximum extrusion speed employed for these alloys can be improved by reducing the amount of main alloying elements (i.e., Al or Zn) [9,10]. For example, Mg–Al- and Mg–Zn-based alloys, such as Mg-0.30Al-0.21Ca-

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0.47Mn, Mg-1.58Zn-0.52Gd, and Mg-0.21Zn-0.30Ca-0.14Mn (wt%), containing small

amounts of alloying elements (<2 wt%) have an extremely high maximum extrusion

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speed of 60 m/min, which is comparable to that of Al alloys. This excellent extrudability is attributed to the suppression of the formation of undesirable Mg17Al12

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and MgZn2 phases and to the high solidus temperature of the -Mg matrix [11–15]. However, these dilute Mg–Al- and Mg–Zn-based alloys exhibit low tensile strengths

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after extrusion. These low strengths are attributed to diminished solid-solution, precipitation, and/or grain boundary strengthening effects because of insufficient

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alloying elements added to the material. Simultaneous improvements in the mechanical properties and extrudability are essential for broadening the application range of

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extruded Mg alloys. In this regard, highly alloyed Mg-Sn-based alloys for extrusion have attracted increasing attention in recent years because of their good extrudability and high strength [16,17]. The solubility limit of Sn in Mg is 14.5 wt% at 561 °C and decreases to ~0.49 wt% at 200 °C. Therefore, a considerable amount of second phase can be formed during hot extrusion in high-Sn-containing Mg alloys. In contrast to Mg–

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Al and Mg–Zn alloys, a thermally stable Mg2Sn phase with a high melting point of 770 °C is formed in Mg–Sn alloys, thereby ensuring both the suppression of hot cracking during extrusion and the precipitation hardening effect after extrusion [18–20]. A Mg-7Sn-1Al-1Zn (wt%) (TAZ711) alloy was successfully extruded (i.e., without hot cracking) at a high extrusion speed of 27 m/min, and the extruded alloy is stronger than the commercial high-strength AZ80 alloy extruded under the same conditions [21].

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However, Mg–Sn-based alloys still exhibit lower extrudability and strength than Al alloys; in addition, Mg2Sn particles formed during extrusion act as cathodic sites, which

lead to substantial deterioration of the corrosion resistance of the extruded material [22]. Therefore, new Mg alloy systems with high strength and excellent extrudability are

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desired.

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The maximum solubility of Bi in Mg is 9.0% at a temperature of 551 °C, and the melting temperature of the Mg3Bi2 phase formed in Mg–Bi alloys is also very high

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(823 °C). This melting temperature is higher than that of the Mg2Sn phase (770 °C) formed in Mg–Sn alloys. Accordingly, Mg–Bi alloys can be considered a highly alloyed

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Mg system with the potential for simultaneously achieving high-speed extrusion and high strength through the formation of thermally stable Mg3Bi2 particles. Recently,

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Sasaki et al. [23] reported that a Mg-6.4Bi-1.3Zn (wt%) alloy exhibited an outstanding aging hardening effect after an isothermal aging treatment because numerous rod-type

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Mg3Bi2 precipitates were uniformly formed on the prismatic plane during aging. Remennik et al. [24] reported that extruded Mg-5Bi-1Ca and Mg-5Bi-1Si (wt%) alloys manufactured using rapidly solidified ribbons exhibited high RT tensile elongations of >40%. They attributed these elongations to the ultrafine grain structure that accommodates deformation through the grain-boundary sliding mechanism. Somekawa

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and Singh [25] reported that an ultrafine-grained extruded Mg-2.5Bi (wt%) alloy exhibited an extremely high elongation of 170% under RT tensile deformation at a strain rate of 1 × 10−3 s−1. This result indicates that superplasticity at RT can occur in the Mg–Bi alloy system. Furthermore, Meng et al. [26] recently developed an extruded Mg8Bi-1Al-1Zn (wt%) alloy with a good strength–elongation balance, which was attributed to a fully dynamic recrystallized (DRXed) microstructure and numerous

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dynamically precipitated Mg3Bi2 particles. Mg–Bi alloys can ensure excellent extrudability through the formation of thermally stable second phases, high strength through the formation of numerous fine precipitates, and high ductility through the

formation of fine DRXed grain structures. However, fundamental research on Bi-

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induced changes in the deformation behavior during hot extrusion and the resultant

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microstructure and mechanical properties of extruded Mg material is lacking. Therefore, in the present study, the effect of Bi addition on the dynamic recrystallization (DRX)

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and dynamic precipitation behaviors during extrusion was investigated. In addition, the microstructural characteristics and tensile properties of extruded pure Mg were

(wt%) alloys.

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systematically investigated through comparison with extruded Mg-6Bi and Mg-9Bi

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2. Experimental procedure

A pure Mg and two Mg–Bi binary alloys, Mg-6Bi and Mg-9Bi (wt%), were used

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in the present study. The added Bi contents of 6 and 9 wt% were determined to be smaller and larger, respectively, than the maximum Bi solubility in Mg (7.7 wt%). Cast billets for extrusion were prepared via conventional mold casting under a mixed-gas atmosphere of CO2 and SF6. A pure Mg ingot and Bi granules with a purity of 99.99% were melted in a carbon crucible, maintained at 20 min at 720 °C for stabilization, and

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poured into a steel mold preheated to 210 °C. The chemical composition of the cast billets was measured using an inductively coupled plasma (ICP) analyzer (PerkinElmer, Optima 7300DV). The compositions, Mg-5.93Bi and Mg-8.61Bi (wt%), were found to be similar to the nominal values of Mg-6Bi and Mg-9Bi alloys, respectively (Table 1). Using an electric furnace, the cast billets were subjected to a 24-h homogenization heat treatment at 500 °C under an inert gas atmosphere containing a mixture of Ar and SF6.

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Afterward, the billets were water-quenched to prevent static precipitation of the Bi dissolved in the α-Mg matrix during cooling. The homogenized billets were machined

to a diameter and length of 80 mm and 200 mm, respectively, for extrusion. The resulting cylindrical billets were preheated at 350 °C for 1 h in a resistance furnace and

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then subjected to a direct extrusion process (extrusion temperature: 350 °C, extrusion

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ratio: 20.3, and ram speed: 1 mm/s). A 300-ton extrusion machine and flat-faced extrusion die with a rectangular hole were employed, and 38 mm (width) × 6.5 mm

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(thickness) extruded plates were obtained.

Microstructures of the homogenized billets were observed by optical microscopy

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(OM; OLYMPUS, JP / BX53) and field-emission scanning electron microscopy (FESEM; Hitachi, SU8220). For microstructural observations, billet samples were

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mechanically polished using progressively finer grades of emery paper. The samples were etched with an acetic picral solution composed of 10 mL acetic acid, 3.0 g picric

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acid, 10 mL distilled water, and 100 mL ethanol. The average grain size of the homogenized pure Mg, Mg-6Bi, and Mg-9Bi billets was measured using the linear intercept method applied in wide areas of 283.4, 44.3, and 44.3 mm2, respectively, in the optical micrographs. The presence of second phases in the homogenized billets and extruded materials was confirmed via X-ray diffraction (XRD) analysis (Rigaku, D /

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Max-2500) at a scan speed of 2°/min. Furthermore, the area fraction of the undissolved second-phase particles remaining in the homogenized billets was measured in a 2.25 mm2 area of the SEM micrographs using the ImageJ software (National Institutes of Health). The number and size of the particles were measured on the basis of an equivalent circle diameter using the IMT isolution DT software (IMT i-solution Inc.). Electron backscatter diffraction (EBSD) analysis of the extruded materials was

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performed on the ED–transverse direction (TD) plane at the mid-thickness and midwidth of the extruded materials by FE-SEM using an apparatus (Hitachi, SU-70) equipped with an EBSD detector. The TexSEM Laboratories Orientation Imaging

Microscopy (TSL OIM) 7.0 software was used to analyze the obtained data with a

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confidence index greater than 0.1. Moreover, the longitudinal cross-section micrograph

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and fracture surface of the fractured tensile specimens associated with the extruded materials were also observed by FE-SEM.

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For tensile testing, dog-bone-shaped specimens (gauge dimensions: 25 mm length, 6 mm width, and 3 mm thickness) were machined from the extruded materials along the

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ED, in accordance with standards ASTM E8/E8M. Tensile tests were conducted at RT using an Instron 8516 universal testing machine operated at a strain rate of 1 × 10−3 s−1.

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To ensure repeatability and confirm the consistency of the results, tensile tests for each material were performed in triplicate and the average values were used for the sake of

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simplicity. Thermal properties of the Mg–Bi binary alloys were analyzed with the ascast Mg-6Bi and Mg-9Bi billets using a differential scanning calorimeter (DSC; HCT-2) at a heating rate of 10 °C/min and were compared with the properties of commercial AZ80 alloy. 3. Results and discussion

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3.1. Incipient melting temperature of Mg–Bi binary alloys Fig. 1 shows the DSC thermograms of the as-cast Mg-6Bi and Mg-9Bi alloys and the commercial AZ80 alloy; the total alloying content of the AZ80 alloy (8.5 wt%) is similar to that of the Mg-9Bi alloy (9.0 wt%). The incipient melting temperature of the AZ80 alloy is relatively low, 433 °C (Fig. 1), which is attributed to the presence of a Mg17Al12 phase with a relatively low melting temperature of 438 °C [10]. As a result,

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hot shortness can readily occur under high-speed extrusion conditions, leading to poor extrudability of the alloy. Indeed, Park et al. [17] have reported that, when AZ80 alloy is extruded at an exit speed of 6 m/min at 350 °C, severe hot cracking occurs during hot

extrusion. By contrast, the incipient melting temperatures of the Mg-6Bi and Mg-9Bi

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alloys are 550 °C and 551 °C, respectively, which are substantially greater than that of

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the AZ80 alloy (433 °C), indicating that the Mg–Bi alloys exhibit high thermal stability compared with the AZ80 alloy. A previous study demonstrated that the TAZ711 alloy

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has a high incipient melting temperature of 545 °C; thus, this alloy can be extruded at a high extrusion speed of 27 m/min at 350 °C without cracking [17]. Given that the Mg–

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Bi binary alloys have slightly higher incipient melting temperatures than the TAZ711 alloy and that the melting temperature of the Mg3Bi2 phase is substantially higher than

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that of the Mg2Sn phase, the Mg–Bi-based alloy can exhibit excellent extrudability, with a maximum extrusion speed of 30 m/min or more.

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3.2. Bi-induced variations in the microstructural evolution of homogenized billets Fig. 2 shows optical and SEM micrographs of the homogenized billets. The

average grain sizes of the pure Mg, Mg-6Bi, and Mg-9Bi billets are 1282 μm, 249 μm, and 283 μm, respectively (Fig. 2(a)–(c)). These results indicate that the addition of Bi to pure Mg leads to considerable grain refinement of the cast material, consistent with

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previously reported results [26,27]. In the homogenized pure Mg billet, a few undissolved particles, identified as a Mg–Si phase via EDS analysis, are present along the grain boundaries (Fig. 2(d)). These particles can be formed by Si impurities contained in the pure Mg ingot (Table 1). However, these particles are small and few in number (area fraction: 0.1%); their influence on the deformation behavior during hot extrusion is therefore negligible. The number of undissolved particles remaining in the

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homogenized billet increases with increasing added-Bi content (Fig. 2(e) and (f)). The particles present in the homogenized Mg–Bi binary billets are a Mg3Bi2 phase, which is

commonly observed in Bi-containing Mg alloys [23,26]. The area fraction of the

undissolved Mg3Bi2 particles in the Mg-9Bi billet (2.4%) is ~3.4 times greater than that

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in the Mg-6Bi billet (0.7%). Moreover, the XRD results of the homogenized billets (Fig.

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3(a)) reveal that the peaks associated with the Mg3Bi2 phase in the pattern of the Mg9Bi billet are substantially more intense than those associated with the Mg3Bi2 phase in

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the pattern of the Mg-6Bi billet; this difference in intensity indicates that the former sample contains more undissolved Mg3Bi2 particles than the latter sample.

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Relatively large (size >1 μm) second-phase particles in a billet can promote recrystallization during hot forming processes [28]. Fig. 4 shows the size distribution

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and number density of the undissolved particles with sizes larger than 1 μm (size taken as an equivalent circular diameter) of the homogenized Mg-6Bi and Mg-9Bi billets. The

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average particle size of the Mg-9Bi billet (4.7 μm) is larger than that of the Mg-6Bi billet (3.3 μm). Similarly, the number density of the particles in the former (1186 mm−2) is 1.32 times greater than that in the latter (901 mm−2). This difference in the size and number density of the undissolved particles in the billets can be explained from the equilibrium phase diagram. Fig. 5 shows the equilibrium phase diagram for the Mg–xBi

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(0 ≤ x ≤ 12 wt%) binary system calculated using the FactSage software. The diagram reveals that, at a homogenization temperature of 500 °C, the Mg-6Bi and Mg-9Bi alloys both lie in a two-phase region consisting of α-Mg and Mg3Bi2. This observation indicates that the coarse Mg3Bi2 phase formed during solidification in the casting process is only partially dissolved during this homogenization heat treatment. Because the solubility limit of Bi at 500 °C is 4.7 wt%, only 4.7 wt% Bi of the added-Bi content

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dissolves into the α-Mg matrix; the residual Mg3Bi2 phase remains in the billet after homogenization. Bi contents of 1.3 wt% and 4.3 wt% for the Mg-6Bi and Mg-9Bi

alloys, respectively, exceed the solubility limit. Hence, a substantially larger number of undissolved Mg3Bi2 particles are present in the homogenized Mg-9Bi billet than in the

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homogenized Mg-6Bi billet.

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3.3. Bi-induced variations in microstructural evolution of extruded material Fig. 6 shows the inverse pole figure maps and grain size distributions of the

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extruded materials. The average grain sizes of the extruded materials (15.7–51.6 μm) are substantially smaller than those of the homogenized billets (249–1282 μm) because

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of grain refinement via DRX during hot extrusion. During the hot extrusion process, the dislocation density in a material increases rapidly because of large plastic deformation;

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the high strain energy accumulated in the material is subsequently relieved via DRX, resulting in a fine-grained extruded material. As shown in Fig. 6, the extruded pure Mg

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material is characterized by a fully DRXed structure consisting of equiaxed DRXed grains. By contrast, the extruded Mg-6Bi and Mg-9Bi materials are characterized by a partially DRXed structure consisting of fine, equiaxed DRXed grains and coarse, distorted unDRXed grains. The area fractions of the DRXed grains (i.e., the DRX fraction) in the extruded Mg-6Bi and Mg-9Bi materials are 76.4% and 89.9%,

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respectively. These fractions are considerably lower than the DRX fraction of the extruded pure Mg material. Collectively, these results indicate that the addition of Bi retards the DRX behavior of pure Mg during hot extrusion despite the presence of numerous undissolved particles that can promote DRX in the Mg–Bi binary alloys. The average grain size of the extruded material decreases with increasing amount of added Bi increases (Fig. 6). The addition of 6 wt% Bi leads to a relatively small size

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reduction (from 51.6 μm to 41.3 μm), whereas the addition of 9 wt% Bi leads to substantial grain refinement to 15.7 μm. This decrease in the average grain size of the

extruded material results from Bi-induced variations in the size and area fraction of the

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unDRXed and DRXed grains. The inverse pole figure maps and average grain sizes of

the unDRXed and DRXed regions of the extruded Mg-6Bi and Mg-9Bi materials are

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shown in Fig. 7. The average size of the DRXed grains (8.8 μm) in the extruded Mg-6Bi material is substantially smaller than that (51.6 μm) of the extruded pure Mg material

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(Fig. 7(a)). However, the extruded Mg-6Bi material contains many (area fraction: 23.6%) coarse unDRXed grains (average size: 147 μm, maximum width: ~300 μm; Fig.

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7(b)). As a result, the average grain size (41.3 μm) of this material is slightly smaller than that (51.6 μm) of the extruded pure Mg material. The average size (7.5 μm) of the

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DRXed grains in the extruded Mg-9Bi material is 15% smaller than that (8.8 μm) of the extruded Mg-6Bi material (Fig. 7(c)). Furthermore, the area fraction, average size, and

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maximum width of the unDRXed grains comprising the extruded Mg-9Bi material (10.1%, 89 μm, and ~150 μm, respectively) are almost half those (23.6%, 147 μm, and ~300 μm, respectively) of the extruded Mg-6Bi material. That is, when the added-Bi content increases from 6 wt.% to 9 wt.%, the area fraction of the DRXed grains increases (from 76.4% to 89.9%) and the average sizes of both the unDRXed and

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DRXed grains decrease. These lead to a substantial reduction (from 41.3 μm to 15.7 μm) in the average grain size of the extruded material. The area fraction of the DRXed grains strongly affects the average grain size of the extruded material. This fraction decreases from 100% to 76.4%, when 6 wt% Bi is added to the pure Mg; however, the addition of another 3 wt% Bi to the Mg-6Bi alloy leads to an increase (from 76.4% to 89.9%) in the area fraction (Fig. 6). These changes (decrease and increase) in the DRX fraction of the

behavior during hot extrusion (see Section 3.5). 3.4. Bi-induced dynamic precipitation and texture evolution

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extruded Mg-6Bi and Mg-9Bi materials result from Bi-induced variations in the DRX

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The SEM micrographs of the extruded materials are presented in Fig. 8, which

shows the fine precipitates dynamically formed during extrusion. A few undissolved

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Mg–Si particles are observed in the precipitate-free extruded pure Mg material (Fig. 8(a)). By contrast, numerous fine precipitates and relatively coarse undissolved Mg3Bi2

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particles are present in the extruded Mg-6Bi and Mg-9Bi materials (Fig. 8(b) and (c)). In the Mg-Bi binary alloys, fine Mg3Bi2 precipitates (size: 100–400 nm) are formed

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during hot extrusion by consuming the Bi atoms supersaturated in the α-Mg matrix. The area fractions of relatively coarse (size > 1 μm) undissolved Mg3Bi2 particles in the

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extruded Mg-6Bi and Mg-9Bi materials (0.9% and 2.2%, respectively) are nearly identical to those of the corresponding homogenized billets (0.7% and 2.4%,

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respectively). The undissolved particles present in the homogenized billets are merely fragmented and rearranged (rather than being newly formed) during extrusion because of the large strain imposed; accordingly, their area fraction remains unchanged after extrusion. Moreover, fine Mg3Bi2 precipitates uniformly distributed throughout the materials comprise similar area fractions (11.8% and 12.2%, respectively) of the

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extruded Mg-6Bi and Mg-9Bi materials despite the 3 wt% difference in their Bi content. The Mg–Bi binary phase diagram (Fig. 5) shows that 4.7 wt% of Bi can be dissolved into the α-Mg matrix at the homogenization temperature of 500 °C. The dissolved-Bi contents of the homogenized Mg-6Bi and Mg-9Bi billets are therefore the same, although the Mg-9Bi billet contains a greater number of undissolved Mg3Bi2 particles. At the extrusion temperature of 350 °C, the solubility of Bi in Mg is only 0.5 wt% in the

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equilibrium phase diagram (Fig. 5). Hence, numerous Mg3Bi2 precipitates are formed during hot extrusion because of the large difference in the Bi solubility at the

homogenization and extrusion temperatures and of the high strain (3.01) and strain rate

(0.23 s−1) imposed during extrusion. Moreover, the Mg3Bi2 precipitates form in equal

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numbers during extrusion of the Mg-6Bi and Mg-9Bi materials because of the identical

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contents of Bi dissolved in the α-Mg matrix of the two homogenized billets. Because a high density of these precipitates is formed during extrusion, the XRD peak intensities

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of the Mg3Bi2 phase comprising the extruded Mg-Bi materials are considerably greater than those of the Mg3Bi2 phase comprising the homogenized Mg-Bi billets (Fig. 3(b)).

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In addition, the peak intensities of the Mg3Bi2 phase are greater in the XRD pattern of the Mg-9Bi material, which contains a larger number of undissolved Mg3Bi2 particles,

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than in the pattern of the Mg-6Bi material. Fig. 9 shows (0001) pole figures of the DRXed, unDRXed, and entire regions

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measured on the normal direction (ND)–TD plane of the rectangular-shaped extruded materials produced in this study. The texture of extruded Mg alloy sheets or plates is strongly affected by the width-to-thickness ratio of the extruded material [29]. At a width-to-thickness ratio of 5.85:1 in the present study, a large compressive stress along the ND and a small compressive stress along the TD are simultaneously applied to the

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material during extrusion. As a result, a fiber basal texture composed of a strong ND texture and a weak TD texture is generated for all of the extruded materials (Fig. 9(a)– (c)). The maximum texture intensities (13.2 and 20.5) of the unDRXed region in the extruded Mg-6Bi and Mg-9Bi materials, which have a partially DRXed structure, are substantially higher than those (6.3 and 6.1) of the DRXed region. Moreover, the unDRXed region is characterized by an ND texture, unlike the DRXed region, which is

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characterized by both ND and TD textures (Fig. 9(d)–(g)). This difference in texture results from the fact that unDRXed grains are continuously deformed without the formation of strain-free grains via DRX and their basal planes are rotated perpendicular to the principal stress direction (i.e., ND) during the entire extrusion process. The

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extruded Mg-6Bi and Mg-9Bi materials consist of strongly textured unDRXed grains;

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however, the corresponding maximum texture intensities (7.9 and 8.7) are lower than that (10.8) of the extruded pure Mg material. This difference is attributed to the fact that

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the texture intensities (6.3 and 6.1) of the DRXed grains in the former materials are lower than the texture intensity (10.8) of the latter material. As shown in Fig. 6, in the

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pure Mg, DRX is generated throughout the material during extrusion, whereas in the Bicontaining alloys, inhomogeneous deformation occurs during extrusion. Fine DRXed

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grains are therefore predominantly formed in the locally strain-concentrated regions. Furthermore, DRXed grains formed in localized high-strain regions (e.g., shear bands,

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deformation bands, twin bands, and near second-phase particles) have a weaker texture than those formed at the grain boundaries of an initial billet [28,30–32]. Therefore, the addition of Bi induces changes in the main nucleation sites of DRX during extrusion, thereby leading to a weakening of the texture characterizing the DRXed grains. As a result, the overall textures of the extruded Mg-6Bi and Mg-9Bi materials are weaker

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than that of the extruded pure Mg material. The Bi-induced changes in the DRX mechanisms are discussed in Section 3.5. 3.5. Bi-induced variation in the DRX behavior during extrusion As previously mentioned, the extruded pure Mg material is characterized by a fully DRXed grain structure, whereas the extruded Mg–Bi binary materials have a partially DRXed grain structure (Fig. 6). These observations indicate that the DRX during hot

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extrusion is suppressed by the addition of Bi. The dominant DRX mechanisms activated during hot deformation of Mg alloys vary with the deformation temperature; discontinuous DRX (DDRX) is the main mechanism operating under deformation at a

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relatively high temperature of 350 °C, i.e., the extrusion temperature employed in the present study [8]. DDRX generally occurs as follows: Initial grains are elongated along

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the metal flow and the dislocation density increases gradually because of the large strain applied during extrusion. In particular, dislocations are piled up around the grain

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boundaries and this localized strain is relieved by strain-induced grain-boundary migration, which causes grain-boundary bulging. As a result, new strain-free grains are

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nucleated along the grain boundaries, and subsequent growth of these newly formed grains (i.e., the DRXed grains) lowers the internal strain energy of the material [33].

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The boundaries act as nucleation sites for DRXed grains formed through DDRX. Therefore, the DRX fraction of the extruded materials increases with decreasing grain

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size of the initial billet because of the increased number of DRX nucleation sites [34,35]. In this study, the average grain sizes of the homogenized Mg–Bi billets are substantially smaller than that of the homogenized pure Mg billet (Fig. 2(a)–(c)); that is, more nucleation sites for DRX occur in the Mg–Bi billets than in the pure Mg billet. Nevertheless, the DRX fractions of the extruded Mg–Bi materials are smaller than that

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of the extruded pure Mg material (Fig. 6), indicating that DDRX is substantially inhibited by other factors induced by the Bi addition. To investigate the DRX behavior during extrusion, we analyzed the microstructures of the extrusion butts associated with the extruded pure Mg and Mg-9Bi materials by water-quenching the reaming part of the billet immediately after extrusion. The water-quenched samples were subsequently evaluated by OM, EBSD, and FE-SEM

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(Fig. 10). The analyzed extrusion butts and the locations of the observed microstructures on the cross-sectional area of the butt are shown in Fig. 10(a), where the

points denoted as A and B indicate the positions observed in the pure Mg and Mg-9Bi

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butts, respectively. An optical micrograph of the pure Mg butt reveals an elongated

unDRXed grain with a highly serrated boundary (Fig. 10(b)). In the early stage of

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extrusion, DRXed grains are formed primarily along the initial grain boundaries, thereby resulting in a necklace-type DRXed grain structure. With further deformation,

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additional DRXed grains are formed around these grains; this formation of DRXed grains occurs continuously from the initial grain boundaries of the billet toward the

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grain interior [8,36]. Accordingly, unDRXed grains remaining at the center of the initial grains are surrounded by the DRXed grains formed through boundary bulging, which

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results in a rugged boundary of the unDRXed grains. The grain boundaries and precipitates in the extruded Mg-9Bi butt were examined

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using an inverse pole figure map showing the grain structure and an FE-SEM image showing the distribution of Mg3Bi2 precipitates. The map and image both correspond to the point denoted B in the right side of Fig. 10(a) and are superimposed (Fig. 10(c)). Compared with point A, point B is closer to the extrusion die exit (Fig. 10(a)), indicating that greater strain is imposed in the observed region of the Mg-9Bi material

17

than in that of the pure Mg material. Nevertheless, more unDRXed grains are present in the Mg-9Bi material observed at position B than in the pure Mg material observed at position A (Fig. 10(b) and (c)). In addition, the boundaries of these unDRXed grains are smoother than those of the pure Mg material because of the fine Mg3Bi2 precipitates formed during extrusion of the Mg-9Bi material. The distribution of fine particles is well known to substantially retard DRX during hot deformation [37]. Via the boundary

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pinning effect, these particles lead to a sizeable reduction in the mobility of grain boundaries. Consequently, DRX is retarded because the boundary bulging required for the nucleation of new DRXed grains is inhibited; this metallurgical phenomenon is

referred to as the Zener pinning effect [32]. A high-magnification micrograph of the

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extrusion butt corresponding to the Mg-9Bi material shows that the boundary of initial

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grains in the billet is pinned by the numerous fine precipitates; consequently, the grain boundary is only mildly serrated (Fig. 10(d)). The size and number of precipitates

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observed in the extrusion butt are nearly identical to those of precipitates in the extruded material (Figs. 8(c) and 10(d), whereas the area fraction of DRXed grains (24.4%) in the

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extrusion butt is substantially lower than that (89.9%) in the extruded material (Fig. 6(c) and 10(c)). This difference in the area fraction of DRXed grains indicates that dynamic

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precipitation occurs substantially faster than DRX during hot extrusion. Therefore, in the Mg–Bi materials, numerous fine Mg3Bi2 precipitates are formed throughout the

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material at the early stage of hot extrusion and the DRX behavior is suppressed by the Zener pinning effect of these precipitates. This suppression of DRX eventually leads to a substantial decrease in the DRX fraction of the extruded material. In terms of grain growth, the fine Mg3Bi2 precipitates inhibit the growth of DRXed grains through the grain-boundary pinning effect, thereby resulting in considerable grain refinement of the DRXed grains. As a result, the average sizes of the DRXed grains of the extruded Mg– 18

Bi materials are substantially smaller than that of the DRXed grains of the extruded pure Mg material. However, the role of the undissolved particles in the reduction of the DRXed grain size is inconsequential because of their large sizes and low number density. Because the homogenized Mg-6Bi and Mg-9Bi billets have the same number of Bi solute atoms supersaturated in the α-Mg matrix, the number of Mg3Bi2 precipitates

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formed during extrusion is approximately the same (Fig. 8). However, the DRX fraction of the extruded Mg-9Bi material (89.9%) is higher than that of the extruded Mg-6Bi material (76.4%) (Fig. 6), which is attributed to a larger number of undissolved Mg3Bi2

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particles in the homogenized Mg-9Bi alloy billet. Small second-phase particles (size < 0.5 μm) such as the Mg3Bi2 precipitates can suppress DRX behavior through the Zener

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pinning effect (see Section 3.4). For large particles (size > 1 μm), large strain energy accumulates around the particles during hot deformation, thereby resulting in the

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formation of DRXed grains on the periphery of the particles; that is, relatively large particles (>1 μm) serve as the nucleation sites for DRX. The recrystallization

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phenomenon caused by such particles is referred to as particle-stimulated nucleation (PSN) [28,32]. The occurrence of DRX by PSN of such large particles during hot

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extrusion and the resultant increase in the DRX fraction of extruded material have been widely reported in various Mg alloy systems, including the Mg–Al [38], Mg–Zn [39],

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Mg–Sn [40], and Mg–Bi [26] systems. In the present study, the formation of DRXed grains by PSN is also confirmed from the EBSD micrograph observed in the extrusion butt of the Mg-6Bi material (Fig. 11(a)). Many fine DRXed grains are formed around the coarse undissolved particles (see white cycle region in Fig. 11(a)), which demonstrates that the undissolved Mg3Bi2 particles effectively act as nucleation sites for

19

DRX during hot extrusion. By contrast, in the region where the undissolved particles are not present, no DRX occurs even at the grain boundaries (Fig. 11(a)); this behavior is in contrast to the case of the pure Mg material, in which considerable DRX occurs in the extrusion butt (Fig. 10(b)). As previously mentioned, in the Mg–Bi alloys, the DDRX mechanism occurring at the grain boundaries is suppressed because the Mg3Bi2 precipitates interfere with the nucleation of DRXed grains by pinning the grain

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boundaries. This effect is confirmed from the SEM micrograph of the extrusion butt of the Mg-6Bi material, which shows numerous Mg3Bi2 precipitates lying along the initial grain boundary of the billet (Fig. 11(b)).

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In the homogenized Mg-9Bi billet, the undissolved Mg3Bi2 particles (sizes: ~1–10

μm) are distributed throughout the material (Fig. 2(f)) and their size and number density

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are higher than those of the homogenized Mg-6Bi billet (Fig. 4). The particles in the Mg-9Bi billet are more effective (than those in the Mg-6Bi billet) in promoting DRX

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during extrusion because the stress accumulated at particles during deformation increases with increasing particle size [32]. Consequently, greater levels of PSN,

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induced by the undissolved Mg3Bi2 particles during extrusion, occur in the Mg-9Bi material than in the Mg-6Bi material, leading to a higher DRX fraction of the extruded

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Mg-9Bi material. Specifically, when 6 wt% Bi is added to the pure Mg, DRX behavior is suppressed because of the Zener pinning effect of the fine precipitates; thus, the DRX

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fraction of the extruded material decreases by 23.6% (from 100% to 76.4%). Adding an additional 3 wt.% Bi to the Mg-6Bi alloy leads to a substantial increase in the size and number of undissolved particles remaining in the homogenized billet. These undissolved particles result in enhanced levels of PSN during extrusion, consequently leading to an increase in the DRX fraction (from 76.4% to 89.9%) of the extruded

20

material. In addition, because the Mg-9Bi billet contains more nucleation sites for DRX than the Mg-6Bi billet as a consequence of its greater number of undissolved particles, the growth inhibition of the DRXed grains induced by interferences between them is more pronounced in the former than in the latter. As a result, DRXed grains in the extruded Mg-9Bi material are smaller than those in the extruded Mg-6Bi material. 3.6. Bi-induced improvement in the tensile strength of extruded materials

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The tensile stress–strain curves and tensile properties of the extruded materials, which are shown in Fig. 12, reveal that the tensile strengths of the extruded Mg–Bi

materials are considerably higher than that of the extruded pure Mg material. When 6 wt%

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Bi is added to the pure Mg, the tensile yield strength (TYS) and ultimate tensile strength

(UTS) of the extruded material increase by 46.6% and 10.7%, respectively. This

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improvement in strength is attributed to the enhancing effects of grain-boundary hardening, precipitation hardening, and dispersion hardening. These hardening

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mechanisms are manifested through grain refinement (from 51.6 μm to 41.3 μm), formation of fine Mg3Bi2 precipitates (from 0% to 11.8%), and an increase in the

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amount of undissolved particles (from 0.1% to 0.9%), respectively (Table 2). In addition, the texture hardening effect during tension is also enhanced by the Bi addition.

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Fig. 13(a) shows the variation in the area fraction of grains as a function of the c-

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axis deviation angle from the ED (θ) in the DRXed and unDRXed regions of the extruded Mg-6Bi material. The unDRXed grains are continuously deformed without forming new grains with relatively random orientations during hot extrusion. Therefore, their lattice rotation via dislocation slips during extrusion is more pronounced than that of the DRXed grains. As a result, the unDRXed region (average θ: 83.0°) exhibits a stronger texture and a higher θ value than the DRXed region (average θ: 79.2°, see Figs. 21

9 and 13(a)). For the DRXed and unDRXed regions of the extruded materials, the Schmid factor (SF) distribution for basal slip under tension along the ED is shown in Fig. 13(b)–(d). With an increase in the θ from 45° to 90°, the SF for basal slip under tension along the ED decreases from ~0.5 to 0 [41]. Accordingly, the SF values of the unDRXed region, which has a higher θ, are lower than those of the DRXed region because its orientations disfavor basal slip under tension along the ED (Fig. 13(c) and

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(d)). The average SF of the extruded pure Mg material (0.20), which consists of only DRXed grains, is higher than those of the extruded Mg-6Bi and Mg-9Bi materials (0.16 and 0.15) because of the presence of unDRXed grains in the Mg-6Bi and Mg-9Bi

materials (Fig. 13(b)–(d)). This result indicates that, compared with the pure Mg

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material, the Mg–Bi materials exhibit an enhanced texture hardening effect under

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tension along the ED.

Strain hardening and solid-solution hardening effects are also enhanced by the Bi

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addition; this enhancement partially contributes to the improvement in the strength of the extruded materials. Fig. 14 shows the kernel average misorientation (KAM) maps of

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the DRXed and unDRXed regions of the extruded materials. Because of the high strain imposed during extrusion, a high density of statically stored dislocations (SSDs) with a

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net Burgers vector of zero and geometrically necessary dislocations (GNDs) with nonzero Burgers vectors are formed in the material [42]. In the KAM map, the variation

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from blue to red indicates an increase in the density of SSDs and GNDs accumulated in the measured area. The average KAM values of the DRXed region in the extruded Mg– Bi materials (1.05 and 1.02 for the Mg-6Bi and Mg-9Bi, respectively) are higher than that of the DRXed region in the extruded pure Mg material (0.54). This difference is attributable to dislocation multiplication and accumulation by the dynamic precipitates

22

and undissolved particles that act as barriers to dislocation movement. Moreover, in the extruded Mg-6Bi and Mg-9Bi materials, the average KAM values of the unDRXed region, which is continuously deformed during extrusion, are 2.99 and 3.11, respectively. Because of the high density of accumulated dislocations in this region, these values are substantially greater than those of the DRXed region. Accordingly, the average KAM values of the Mg-6Bi and Mg-9Bi materials (1.51 and 1.23) are 2.8 and

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2.3 times greater, respectively, than that of the extruded pure Mg material (0.54). These greater average KAM values indicate that the strain hardening effect induced by the internal strain energy in the Mg-Bi materials is substantially greater than the effect in

the pure Mg material. Furthermore, compared with the extruded Mg-9Bi material, the

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extruded Mg-6Bi material contains a higher area fraction of unDRXed grains with a

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high dislocation density; therefore, the strain hardening effect is more pronounced in the extruded Mg-9Bi material.

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The Bi solubility at an extrusion temperature of 350 °C is 0.5 wt% (Fig. 5); hence, approximately 0.5 wt.% Bi remains as solute atoms in the α-Mg matrix after extrusion

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despite the extrusion conditions being in a non-equilibrium state. The remaining Bi leads to a solid-solution hardening effect of the extruded materials. However, because

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the atomic radii of Mg and Bi are approximately the same (145 pm and 143 pm for Mg and Bi, respectively), the contribution of this effect to the Bi-induced strength

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improvement may be inconsequential. Therefore, the strengthening mechanisms (i.e., grain-boundary hardening, precipitation hardening, dispersion hardening, texture hardening, strain hardening, and solid-solution hardening) affecting the TYS of the material are more effective in the extruded Mg-6Bi and Mg-9Bi materials than in the extruded pure Mg material. The precipitation hardening effect is likely the largest

23

contributor to the strength improvement of the extruded Mg-Bi materials, as suggested by the homogeneous distribution of numerous fine Mg3Bi2 precipitates (the extruded pure Mg material is precipitate-free) (Fig. 8). When the added-Bi content is increased from 6 wt% to 9 wt%, the TYS of the extruded materials increases from 129 MPa to 141 MPa (Fig. 12). The number of Mg3Bi2 precipitates, average SF for basal slip under tension along the ED, and Bi content dissolved in the α-Mg matrix are approximately

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the same for the extruded Mg-6Bi and Mg-9Bi materials (Table 2, Figs. 5, 8, and 13). The aforementioned increase in the added-Bi content leads to an increase in the area

fraction of the undissolved Mg3Bi2 particles comprising the extruded materials, thereby enhancing the dispersion hardening effect. However, the decrease in the average KAM

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value (Table 2 and Fig. 14), which is attributed to the reduction in the area of the

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unDRXed region, weakens the strain-hardening effect. More importantly, the average grain size of the extruded material decreases substantially (from 41.3 μm to 15.7 μm)

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when the added-Bi content increases from 6 wt% to 9 wt%. The corresponding enhancement in the grain-boundary hardening effect therefore plays a major role in

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improving the TYS of the extruded Mg-9Bi material. Notably, the addition of 6 wt% Bi leads to a 41-MPa TYS improvement (from 88 MPa to 129 MPa) of the extruded

ur

material, whereas the addition of another 3 wt.% Bi results in only a 12-MPa (from 129 MPa to 141 MPa) improvement; the increase in TYS values per 1 wt.% added Bi are 6.8

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MPa and 4 MPa, respectively. These results indicate that the addition of more than 6 wt% Bi is less effective at improving the strength of the extruded material than the addition of smaller amounts of Bi, as evidenced by the lack of additional fine Mg3Bi2 precipitates. 3.7. Bi-induced deterioration in tensile ductility of extruded materials

24

Fig. 15 shows the cross-sectional SEM images and fractographs of the fractured tensile specimens of the extruded materials. Many twins are observed in the grains near the fracture surface of the pure Mg specimen after tensile testing (Fig. 15(a)). This observation indicates that the fracture mechanism of the extruded pure Mg material is strongly correlated with the activation of deformation twinning during tension. When extruded Mg alloys with ED-parallel basal planes are tensioned along the ED at RT,

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intragranular {10-11} contraction twins and {10-11}–{10-12} double twins are generally formed. In addition, microcracking occurs along these twins because of the

local concentration of deformation at the twins, which in turn leads to fracture [43]. The cross-sectional microstructure of the fractured pure Mg specimen reveals that a straight

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fracture line is formed along the intragranular twin (see red dashed circle in Fig. 15(a)).

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This cracking at the tension-induced twin bands leads to the formation of many riverpatterned cleavage planes in the fracture surface (Fig. 15(d)). Although the grain

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structure comprising the Mg–Bi specimens is difficult to observe in SEM images, the cross-sectional SEM micrograph in Fig. 15(b) shows that a large portion of the straight

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fracture line is formed in the Mg-6Bi specimen, as in the case of the pure Mg specimen. Moreover, large cleavage planes with river patterns (typical of fracture surfaces caused

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by cracking at twins) are observed in the fracture surface of both specimens (Fig. 15(e)). In addition, the width of the cleavage planes associated with the pure Mg specimen (57

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μm) is consistent with the average grain size of the material (51.6 μm). The extruded pure Mg material has a fully DRXed grain structure with a relatively uniform grain size distribution; hence, contraction and double twins are formed throughout the specimen during tension. In the case of the Mg-6Bi specimen, the width of the cleavage planes (111 μm) is similar to the average size of the unDRXed grains comprising the material (147 μm). 25

The stress required to activate deformation twinning (i.e., twinning stress) increases dramatically with decreasing grain size of a material [44]; this behavior means that deformation twins are more easily formed in coarser grains than in finer grains. The DRXed grains in the extruded Mg-6Bi material are small (average size: 8.8 μm), whereas the unDRXed grains are large (average size: 147 μm; see Fig. 7(a) and (c)). During tension along the ED of the Mg-6Bi specimen, the formation of deformation

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twins is suppressed in the DRXed grains because of the grains’ small size. By contrast, contraction and double twins, which function as initiation sites for microcracks, are

easily formed in the unDRXed grains during the early stage of deformation as a consequence of the reduced twinning stress (caused by the large size of these grains),

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which eventually leads to premature fracture. Because the average size of the unDRXed

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grains of the Mg-6Bi material (147 μm) is 2.8 times greater than the average grain size of the pure Mg material (51.6 μm), the tensile elongation of the Mg-6Bi material (4.1%)

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is substantially smaller than that of the pure Mg material (8.6%). Consequently, the addition of 6 wt% Bi to pure Mg suppresses DRX behavior during extrusion, which

na

leads to the formation of coarse unDRXed grains and, accordingly, a significant reduction in the tensile ductility of the extruded material. When the added-Bi content is

ur

increased from 6 wt% to 9 wt%, the DRX behavior is promoted by enhanced levels of PSN at the undissolved particles, which leads to a decrease in the size and area fraction

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of the unDRXed grains comprising the extruded material. As a result, the fracture mode changes from quasi-cleavage fracture with large flat facets formed by cracking at twins for the pure Mg and Mg-6Bi specimens to ductile fracture with numerous small dimples for the Mg-9Bi specimen (Fig. 15(f)). However, the tensile elongation of the Mg-9Bi material (5.6%) is considerably lower than that of the pure Mg material (8.6%) but is slightly higher than that of the Mg-6Bi material (4.1%). This low ductility of the Mg26

9Bi material results from the cracking of the coarse undissolved Mg3Bi2 particles during tension. In the extruded Mg-9Bi material, many undissolved particles are homogeneously distributed throughout the material; these particles are easily broken during tensile deformation because of their brittle nature and large size. Indeed, cracked particles are observed in the fracture surface of the Mg-9Bi specimen (Fig. 15(f)). The corresponding cross-sectional micrograph shows the propagation of the cracks from the

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initiation sites (i.e., particles) to the surrounding α-Mg matrix (Fig. 15(c)). The undissolved particles in the homogenized Mg-9Bi billet contribute substantially to the increase in the DRX fraction of the extruded material, thereby suppressing the formation of undesirable twins during tension. However, these particles act as crack

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initiation sites during tensile deformation, leading eventually to deterioration of the

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material’s ductility.

Bi is an inexpensive alloying element that is highly soluble in Mg (solubility limit:

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~9 wt%) and forms a very thermally stable phase (Mg3Bi2) during extrusion. The formation of numerous Mg3Bi2 precipitates can substantially improve the extrudability

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and mechanical properties of extruded products without dramatically increasing their cost because this approach avoids the addition of expensive rare-earth elements. In the

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present work, the variations in the metallurgical phenomena that occur during hot extrusion (i.e., DRX and dynamic precipitation behaviors) were studied on a

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fundamental basis. The microstructural characteristics of the extruded material (e.g., grain size, texture, precipitates, and internal strain energy), and the mechanisms related to the mechanical properties (i.e., changes in the strengthening and fracture mechanisms), which are all caused by Bi addition to pure Mg, were also investigated. The results obtained from this study provide useful information for the development of

27

new, low-cost multi-component Mg–Bi-based alloys that simultaneously exhibit excellent extrudability and outstanding mechanical properties, which are major factors required for expanding the application range of extruded Mg alloys. The development of such materials necessitates further in-depth research on the addition of other alloying elements and on controlling the extrusion process conditions. The acquired knowledge could be used to develop extruded Mg–Bi-based materials with a fully DRXed grain

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structure without large particles that can induce cracking during tensile deformation; this grain structure can lead to a substantial improvement in the microstructural

homogeneity and strength of the extruded material and prevent the deterioration of the

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tensile ductility.

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4. Conclusion

This study demonstrates that the addition of Bi to a pure Mg strongly affects its

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DRX and dynamic precipitation behaviors during hot extrusion. The microstructure and mechanical properties of the extruded material are also greatly affected. The addition of

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6 wt% and 9 wt% Bi results in the formation of numerous fine Mg3Bi2 precipitates during the early stage of hot extrusion, thereby retarding DRX through the Zener

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pinning effect. As a result, the extruded pure Mg material is characterized by a fully DRXed grain structure, whereas the extruded Mg–Bi binary materials are composed of

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a partially DRXed bimodal grain structure, which consists of fine DRXed grains and coarse unDRXed grains. When the Bi content increases from 6 wt% to 9 wt%, the DRX fraction of the extruded material increases from 76.4% to 89.9%. This increase in the DRX fraction results from enhanced levels of PSN associated with an increase in the size and number of undissolved Mg3Bi2 particles. The Bi addition results in

28

considerable grain refinement of the DRXed grains through grain-boundary pinning by fine precipitates, thereby leading to a decrease in the average grain size of the extruded material. The tensile strengths of the extruded Mg–Bi materials are considerably greater than those of the extruded pure Mg material, which is attributed to the combined enhancement effect of several strengthening mechanisms. These mechanisms include grain-boundary hardening, precipitation hardening, dispersion hardening, texture

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hardening, strain hardening, and solid-solution hardening, among which the precipitation hardening effect caused by the numerous fine precipitates is likely the most pronounced. However, the tensile ductility of the extruded material deteriorates with Bi addition. This deterioration is induced by cracking at twins formed in the coarse

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unDRXed grains of the Mg-6Bi material and by cracking at large undissolved Mg3Bi2

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particles in the Mg-9Bi material. The aforementioned results indicate that further studies are required for producing extruded Mg–Bi-based materials with excellent mechanical

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properties achieved by suppressing the formation of coarse unDRXed grains and large undissolved particles. This suppression can be realized through control of the Bi content,

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addition of alloying elements, and/or optimization of extrusion process parameters.

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Acknowledgments

This work was supported financially by the National Research Foundation of

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Korea (NRF) Grants funded by the Korean government (MSIP, South Korea) (Nos. 2017R1A4A1015628 and 2019R1A2C1085272).

29

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[39] H. Yu, S.H. Park, B.S. You, Y.M. Kim, H.S. Yu, S.S. Park, Mater. Sci. Eng. A 583 (2013) 25–35. [40] S.H. Kim, J.G. Jung, B.S. You, S.H. Park, Mater. Sci. Eng. A 657 (2016) 406–412. 32

[41] S.W. Lee, S.H. Kim, W.K. Jo, W.H. Hong, W. Kim, B.G. Moon, S.H. Park, J. Alloys Compd. 791 (2019) 700–710. [42] S.I. Wright, M.M. Nowell, D.P. Field, Microsc. Microanal. 17 (2011) 316–329. [43] D. Ando, J. Koike, Y. Sutou, Mater. Sci. Eng. A 600 (2014) 145–152.

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[44] M.A. Meyers, O. Vöhringer, V.A. Lubarda, Acta Mater. 49 (2001) 4025–4039.

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Fig. 1. DSC curves of as-cast Mg–Bi binary alloys and commercial AZ80 alloy.

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Fig. 2. (a–c) Optical and (d–f) SEM micrographs of homogenized (a, d) pure Mg, (b, e) Mg-6Bi, and (c, f) Mg-9Bi billets; davg, fMg-Si and fMg3Bi2 denote the average grain size and area fractions of undissolved Mg–Si and Mg3Bi2 particles, respectively.

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Fig. 3. XRD patterns of (a) homogenized billets and (b) extruded materials.

Fig. 4. Size distribution and number density of undissolved particles in homogenized

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Mg-6Bi and Mg-9Bi billets; Davg and N denote the average diameter and

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number density of particles, respectively.

Fig. 5. Equilibrium phase diagram for Mg–xBi (0 ≤ x ≤ 12 wt%), as calculated using 35

FactSage software; Thomo. and Text. denote the homogenization temperature

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(500 °C) and extrusion temperature (350 °C), respectively.

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Fig. 6. Inverse pole figure maps and grain size distributions of extruded (a) pure Mg, (b) Mg-6Bi, and (c) Mg-9Bi materials; fDRX denotes the area fraction of

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recrystallized grains.

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Fig. 7. Inverse pole figure maps of (a, c) DRXed region and of (b, d) unDRXed region

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of extruded (a, b) Mg-6Bi and (c, d) Mg-9Bi materials; AGS denotes the

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average grain size.

Fig. 8. SEM micrographs showing precipitates and undissolved particles of extruded (a) pure Mg, (b) Mg-6Bi, and (c) Mg-9Bi materials; fppt and fundis denote the area 37

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fractions of Mg3Bi2 precipitates and undissolved Mg3Bi2 particles, respectively.

Fig. 9. (0001) pole figures corresponding to (a–c) entire region, (d, e) DRXed region, and (f, g) unDRXed region of extruded (a) pure Mg, (b, d, f) Mg-6Bi, and (c, e,

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g) Mg-9Bi materials.

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Fig. 10. (a) Photos showing remaining billet after extrusion (i.e., extrusion butt) and

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extrusion butt cross-sectioned for microstructural observation; (b) Optical micrograph obtained at position B (see (a)) in extrusion butt of pure Mg

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material; (c) Microstructure observed at position A (see (a)) in extrusion butt of Mg-9Bi material; this image is obtained by superimposing an inverse pole

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figure map showing the grain structure onto a SEM micrograph showing the Mg3Bi2 precipitates; (d) High-magnification view of region C in (c); Blue

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arrows in (d) indicate the fine Mg3Bi2 precipitates that pin initial grain

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boundary.

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Fig. 11. EBSD and SEM micrographs in extrusion butt of Mg-6Bi material: (a)

superimposed EBSD image of inverse pole figure map and image quality map;

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(b) SEM image showing Mg3Bi2 precipitates that pin initial grain boundary of

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billet.

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Fig. 12. (a) Tensile stress–strain curves and (b) tensile properties of extruded pure Mg,

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Mg-6Bi, and Mg-9Bi materials.

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Fig. 13. (a) Variation in area fraction of DRXed and unDRXed regions as a function of

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c-axis deviation angle from extrusion direction of extruded Mg-6Bi material. Distribution of Schmid factor for basal slip under tension along extrusion

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direction of extruded (b) pure Mg, (c) Mg-6Bi, and (d) Mg-9Bi materials.

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Fig. 14. EBSD kernel average misorientation maps of DRXed and unDRXed regions in

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extruded (a) pure Mg, (b) Mg-6Bi, and (c) Mg-9Bi materials. KAMDRX and KAMunDRX denote the average KAM values of DRXed and unDRXed regions,

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respectively.

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Fig. 15. SEM images showing (a–c) cross-section and (d–f) fracture surface of fractured

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tensile specimens of extruded (a, d) pure Mg, (b, e) Mg-6Bi, and (c, f) Mg-9Bi

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materials.

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Table 1 Chemical compositions of materials used in this study (wt%). Bi

Si

Mn

Fe

Mg

Pure Mg

-

0.027

0.021

0.002

Bal.

Mg-6Bi

5.93

0.028

0.019

0.001

Bal.

Mg-9Bi

8.61

0.031

0.018

0.002

Bal.

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Material

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Table 2 Microstructural characteristics and mechanical properties of extruded materials. Microstructural characteristics* Material

fDRX

davg

dDRX

Mechanical properties** dunDRX

fundis.

fppt

SFbasal

TYS

UTS

EL

(MPa)

(MPa)

(%)

KAM

(%)

(µm)

(µm)

(µm)

(%)

(%)

Pure Mg

100

51.6

51.6

-

0.1

-

0.20

0.54

88

177

8.6

Mg-6Bi

76.4

41.3

8.8

147

0.9

11.8

0.15

1.51

129

196

4.1

Mg-9Bi

89.9

15.7

7.5

89

2.2

12.2

0.16

1.23

141

203

5.1

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Notes: *fDRX, davg, dDRX, dunDRX, fundis., fppt, SFbasal, and KAM denote the area fraction of DRXed grains, average grain size of entire grains, average grain size of DRXed grains, average grain size of unDRXed grains, area fraction of undissolved Mg3Bi2 phases, area fraction of Mg3Bi2 precipitates, Schmid factor value for basal slip under tension along the extrusion direction, and average kernel average misorientation value, respectively. **TYS, UTS, and EL denote the tensile yield strength, ultimate yield strength, and elongation, respectively.

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