Microstructure and Tensile Property Tailoring of Al–5Mg (wt%) Alloy Combined with Continuous Rheo-Extrusion and Sc Modification

Abstract

In this work, Al–Mg alloys fabricated by combining continuous rheo-extrusion (CRE) and Sc modification were proposed for producing Al–Mg alloys with high efficiency and superior mechanical performance. The microstructural evolution and mechanical property response of the CREed Al–5Mg alloy with Sc modification were investigated. The grain refinement and strengthening mechanisms induced by nanoscale Al₃Sc-phase particles in the alloy were discussed. The results showed that an obvious grain refinement effect was achieved in the CREed Al–5Mg alloy as the Sc content increased from 0 to 0.5 wt%, and the average grain size decreased from 52.6 μm to 2.4 μm, respectively. The primary Al₃Sc-phase particles formed during solidification behaved as heterogeneous nucleation sites for the α-Al matrix, while the nanoscale Al₃Sc-phase particles achieved during CRE enhanced the driving force of continuous dynamic recrystallization and the Zener drag force. As a result, a superior grain refinement effect was observed. The ultimate tensile strength, yield strength, and hardness of the alloy were enhanced as the Sc content increased from 0 to 0.5 wt%. Grain boundary strengthening, second-phase strengthening, and dislocation strengthening were the main strengthening mechanisms of the CREed Al–Mg–Sc alloys.

1. Introduction

As a typical medium-strength aluminum (Al) alloy, Al–Mg alloys have a wide array of applications in the aerospace, transportation, and marine industries owing to their superior specific strength and good comprehensive performance [1,2,3]. The strength of Al–Mg alloys is improved due to the solid solution strengthening effect induced by the substantial solid solubility of Mg in the α-Al matrix [4]. Thus, Al–Mg alloys with high Mg content are expected to broaden the applications in the above fields [5,6]. However, high-Mg-containing Al–Mg alloys prepared through the conventional casting method tend to form divorced eutectic structures that are coarse dendrites and Al₃Mg₂-phase particles during solidification, leading to the deformation difficulty and high brittleness of the alloys [7]. Meanwhile, the formation of A1₃Mg₂-phase particles decreases the solid solution content of Mg in the α-Al matrix. Therefore, a weakened solid solution strengthening effect of Mg and insufficient mechanical properties are achieved in Al–Mg alloys. The exhibited mechanical properties are closely related to the microstructures [8]. Thus, it is essential to tailor the microstructure and obtain improved mechanical properties of Al–Mg alloys.

Adjusting the main alloying elements, utilizing heat treatments, and adopting deformation techniques are the main methods to improve the microstructure and thereby artificially regulate the mechanical performance of metal materials [9,10]. Based on extensive studies and given their non-heat-treatable characteristics, the windows for the former two methods are quite narrow for tailoring Al–Mg alloys. Microalloying and deformation have become promising strategies to further modify the microstructure and improve the mechanical properties [11,12]. It is well-known that rare earth (RE) elements and nanomaterials have superior advantages of melt purification and grain/dendrite refinement even in a small amount, thus improving the microstructure and mechanical properties of non-ferrous alloys [13,14]. It is worth mentioning that scandium (Sc) has been referred to as the most effective microalloying RE element for Al–Mg alloys due to the introduced Al₃Sc-phase particles of L1₂-type dispersoids with a similar lattice size to the α-Al matrix [15]. Furthermore, the primary Al₃Sc-phase particles preferentially formed during solidification act as heterogeneous nucleation sites due to the low lattice mismatch and good coherency with the α-Al matrix [16]. In addition, the continuous segregation and the growth of brittle A1₃Mg₂-phase particles are inhibited by the Al₃Sc-phase particles, contributing to the refined and discontinuous Al₃Mg₂-phase particles [17]. Thus, significant grain refinement and enhanced mechanical properties are achieved in Al–Mg–Sc alloys. Zhou et al. [18] reported that the typical dendritic microstructures in the as-cast Al–5Mg (wt%) alloy were eliminated and changed into fine and equiaxed grain microstructures as the Sc content reached 0.4 wt%. As a result, the ultimate tensile strength (UTS) and yield strength (YS) were enhanced from 215.4 and 86.1 MPa to 280.2 and 153.6 MPa, respectively, followed by a decline in elongation (EL) from 22.2% to 14.8% [18]. Moreover, Al₃Sc-phase particles are considered one of the most effective precipitation hardening particles compared with the other common particles in Al alloys [19]. However, Li et al. [20] pointed out that the grain refinement effect was limited upon further increasing the Sc content, which was attributed to the coarse and blocky Al₃Sc-phase particles distributed near the grain boundaries playing no significant role in grain refinement and reducing the effective Sc content in the as-cast Al–3.8Cu–1.2Li–0.4Mg (wt%) alloy. Thus, the application amount of Sc in Al–Mg alloys is limited. Further increasing the Sc addition and forming a large amount of small-sized Al₃Sc-phase particles with uniform distribution for grain refining and precipitation hardening seem to be effective approaches for developing Al–Mg–Sc alloys with high performance.

On the other hand, the shear stress induced by the inhomogeneous flow during deformation contributes to refinement of secondary-phase particles [21]. Meanwhile, the movement of dislocations and grain boundaries is hindered by the refined particles, accelerating the occurrence of recrystallization; hence, the grains are further refined, and the mechanical performance is improved [22]. However, the size of secondary-phase particles tends to be still micron-sized after one-pass deformation. Li et al. [23] investigated the microstructural evolution of an Al–6.0Mg–1.1Mn–0.19Er–0.1Zr (wt%) alloy during hot extrusion and found that the size of phase particles in the hot-extruded alloy was about 5 μm. Thus, severe plastic deformation (SPD) is expected to achieve refined secondary-phase particles and grains with high efficiency [24]. However, the common SPD methods, such as equal-channel angular pressing (ECAP) [25], high-pressure torsion (HPT) [26], and reciprocating extrusion (RE) [27], exhibit long manufacturing processes and limited production size. Thus, a low-cost, high-efficiency, and green deformation method combined with casting and SPD, referred to as continuous rheo-extrusion (CRE), is proposed to tailor the microstructure and mechanical properties of high-Mg-containing Al–Mg alloys that go from liquid to solid in one step [28]. During the CRE process, the high cooling rate during solidification is conducive to the refinement of primary grains/dendrites and secondary-phase particles, while the subsequent high shear strain contributes to the elimination and recrystallization of dendrites and grains, followed by the further refinement of secondary-phase particles to nanoscale [29]. However, few studies on Al–Mg alloys prepared through CRE and Sc modification have been reported. Therefore, the microstructural evolution of CREed Al–Mg alloys with different Sc content requires further investigation.

In this work, the Al–5Mg–xSc (x = 0.1, 0.3, and 0.5 wt%) alloys were manufactured through CRE. The microstructural evolution of CREed Al–5Mg alloys with different Sc additions was investigated. The strengthening mechanisms induced by nanoscale Al₃Sc-phase particles on mechanical properties of the alloys were discussed. The aim of this work is to offer valuable insights into producing high-Sc-containing Al–Mg–Sc alloys through a high-efficiency processing technique and obtaining superior mechanical performance.

2. Materials and Methods

Al–5Mg–xSc (x = 0.1, 0.3, and 0.5 wt%) alloys considered in this study were prepared using high-purity Al (>99.994 wt%) ingots, Al–10.6Mg (wt%), and Al–2.08Sc (wt%) master alloys (provided by Beijing Ryubon New Material Technology, Ltd., Beijing, China). The alloy melt with different Sc content was produced in an intermediate-frequency furnace with high-purity argon protection. After all of the required elements were added during the smelting process, the melt (3 kg of each individual experiment) was heated up to 720 °C, followed by degassing and deslagging with C₂Cl₆, and then it was thoroughly stirred for composition uniformity and finally poured into the entrance of a wheel-shoe gap of a 300 series CRE device. The device is mainly composed of melting tundish, a melt flow channel, a rotational extrusion shoe, a fixed extrusion wheel, and a cooling system. The diameters of the extrusion wheel and the prepared alloy rods were 300 and 9.5 mm, respectively. Details of the device were described in recent studies [28,29]. The alloy rods were water-quenched after the CRE process. The actual chemical composition results of the alloy rods determined using a PerkinElmer Avio500 (Waltham, MA, USA) inductively coupled plasma-optical emission spectrometer (ICP-OES) are displayed in Table 1. The selected compositions were designed to increase, from the top to the bottom of the table, to promote the formation of Sc-containing phase particles in the alloys.

Specimens for microstructural characterization were cut along the extrusion direction (ED). The sectional dimension was 9.5 × 10 mm². In detail, specimens for optical metallography observation were ground and polished using alumina sandpapers and diamond pastes, respectively, according to standard metallographic techniques. A Leica DMI 5000 M (Wetzlar, Germany) optical microscope (OM) was utilized to characterize the optical microstructure of the specimens under polarized light mode. The average grain sizes were measured via Image Pro Plus software version 6.0. The morphology, distribution, and composition of the secondary-phase particles were analyzed using a Zeiss Ultra Plus (Zeiss, Oberkochen, Germany) field emission-scanning electron microscope (SEM) equipped with a back-scattered electron (BSE) detector and an energy dispersive X-ray spectrometry (EDS) detector. X-ray diffraction measurements were obtained using the Bruker D8 (Billerica, MA, USA) Advance X-ray diffractometer (XRD) with Cu-Kα radiation for the phase analysis. The rated voltage and current of the XRD equipment were 60 kV and 80 mA, respectively. Furthermore, the scanning range and the speed utilized in this work were 20°~90° and 2 °/min, respectively. FEI G² 20 transmission electron microscopy (TEM) was performed to further observe the nanoscale secondary-phase particles in the alloys. The TEM specimens with a diameter of 3 mm were punched and ground to a final thickness of 60 μm and then electro-polished using the Tenupol-5 Struers twin-jet polisher (Struers, Ballerup, Denmark) in a solution of 30 vol% nitric acid and 70 vol% methanol at a temperature of −25 °C.

Specimens for mechanical performance testing were cut from the middle of the alloys. The specification of specimens was smooth and dog-bone-shaped with a circular cross section, and the gauge size was 6 mm in diameter and 30 mm in length. A Shimadzu AG-Xplus (Kyoto, Japan) electronic universal tensile machine was utilized for tensile testing at a strain rate of 10^–3 s^–1 under an ambient temperature of 25 °C. Three specimens for each condition were examined for reproducibility. A SIOMM HVD-1000MP Vickers hardness tester (Shanghai Jvjing Precision Instrument Manufacturing Co., Ltd., Shanghai, China) was used for Vickers hardness testing at an applied load of 200 gf and a dwell time of 10 s. The average value of nine indentations for each condition was considered as the final result.

3. Results

3.1. Microstructural Characterization

The typical optical microstructures and statistical average grain sizes of the CREed AM alloys modified by different Sc additions are displayed in Figure 1. On the contrary, no obvious dendrites but fine and equiaxed grains are observed in the CREed AM alloy, and the average grain size of the alloy is 52.6 ± 3.8 μm, as shown in Figure 1a. Compared with the AM alloy prepared through CRE, the grains are further refined in the AM alloy fabricated by combining CRE and Sc modification. As the Sc content increases to 0.1 wt%, the average grain size of the alloy decreases to 34.3 ± 2.4 μm, as presented in Figure 1b. An obvious grain refinement effect is achieved in the alloy upon further increasing the Sc addition to 0.3 and 0.5 wt%, and the average grain sizes are 11.2 ± 1.7 and 2.4 ± 0.6 μm, which are 78.7% and 95.4% declines in comparison to that of the CRE AM alloy, respectively, as displayed in Figure 1c,d. In addition, the statistical average grain size result of the CREed AM–xSc alloys is shown in Figure 1e.

Backscattered SEM micrographs of the CREed AM(–0.5Sc) alloys are displayed in Figure 2 for investigating the microstructural variation of the alloy after Sc modification. No obvious secondary-phase particles are precipitated in the CREed AM alloy, as presented in Figure 2a, which is consistent with the XRD analysis results in Figure 3. Furthermore, the EDS results of point A in the alloy are shown in Figure 2c, indicating that Mg atoms tend toward a solid solution in the α-Al matrix during the CRE process [30]. On the contrary, refined white secondary-phase particles with short rod-like or spherical morphologies are observed in the CREed AM–0.5Sc alloy, as shown in Figure 2b. The EDS results of point B are presented in Figure 2d, indicating that the white secondary-phase particles mainly consist of Al and Sc elements. Combined with the XRD analysis results in Figure 3, the Al–Sc-phase particles in the CREed AM–0.5Sc alloy are confirmed to be Al₃Sc-phase particles [19].

Figure 4 presents the micrographs obtained through TEM of the AM(–0.5Sc) alloys manufactured through CRE. No obvious nanoscale secondary-phase particles are observed in the CREed AM alloy, as shown in Figure 4a, which is consistent with the conclusion above that Mg atoms are a solid solution in the α-Al matrix. Meanwhile, nanoscale Al₃Sc-phase particles with typical petaloid-like morphology show a dispersed distribution in the CREed AM alloy with 0.5 wt% Sc modification, as presented in Figure 4b. The high-resolution transmission electron microscope (HRTEM) image indicates that the Al₃Sc-phase particles have a completely coherent interface relationship with the α-Al matrix, which agrees with the previous studies [31,32].

3.2. Mechancial Properties

The variations in strength, ductility, and hardness have been measured to investigate the mechanical property response of the CREed AM alloys with different Sc additions. Figure 5 displays the engineering stress–strain curves of CREed AM alloys modified with different Sc content. The ultimate tensile strength (UTS), yield strength (YS), and elongation (EL) of the CREed AM alloy are 257.6 MPa, 99.8 MPa, and 52.8%, respectively. Meanwhile, the hardness of the CREed AM alloy is 69.8 HV. It is noticed that the UTS, YS, and hardness of the CREed AM alloy are enhanced upon increasing the Sc content, while the EL displays a tendency to decline. When the content of Sc reaches 0.1 wt%, the UTS, YS, and hardness of the CREed AM alloy are 331.3 MPa, 133.5 MPa, and 101.4 HV, respectively, followed by a decrease in EL to 41.8%. It is worth mentioning that an obvious improvement in mechanical properties is achieved in the CREed AM alloy with 0.5 wt% Sc modification. The UTS, YS, and hardness of the CREed AM–0.5Sc alloy are 456.4 MPa, 220.1 MPa, and 143.7 HV, which are 17.3%, 21.4%, and 13.6% enhancements compared to those of the CREed AM–0.3Sc alloy, respectively. Meanwhile, the EL of the CREed AM–0.5Sc alloy decreases to 24.8%, which is a 22.5% decrement in EL of the CREed AM–0.3Sc alloy. Table 2 lists the tensile properties and hardness results of the AM alloys prepared through CRE and Sc modification.

4. Discussion

4.1. Grain Refinement Induced by Sc Modification During Solidification

Compared with the CREed AM alloy, a superior grain refinement effect is achieved in the CREed AM alloy modified by Sc. It is well-known that undercooling and heterogeneous nucleation are the main methods for microalloying elements to refine the microstructures of Al alloys during solidification [33,34]. Owing to the chilling effect induced by the extrusion wheel and shoe during CRE, the cooling rate of the AM(–xSc) alloy melt is enhanced, thus contributing to the refinement of primary microstructures of the alloys. Moreover, the total undercooling (∆T) at the dendrite tips is described as follows [35]:

∆T = ∆T_t + ∆T_c + ∆T_r + ∆T_k(1)

where ∆T_t and ∆T_c are thermal and solutal undercooling and ∆T_r and ∆T_k are curvature and kinetic undercooling, respectively. Due to the large atomic radius and low solubility of Sc in the α-Al matrix, Sc atoms tend to enrich in front of the solid–liquid interface during solidification [19]. Thus, solute redistribution occurs and solutal undercooling is enhanced upon increasing the amount of Sc, contributing to the reduction in the melting point, promoting the grain nucleation, and refining the grains of the AM–xSc alloys [18]. In addition, the growth restriction factor (Q_GRF) is utilized to describe the inhibition effect of Sc element enrichment on the grain growth of Al alloys, which is expressed as follows [36]:

Q_GRF = m_Sc C_Sc (k_Sc − 1)(2)

where m_Sc is the concentration of solute Sc in the alloy melt and C_Sc and k_Sc are the liquidus slope and equilibrium partition coefficient in the Al–Sc binary alloy-phase diagram, respectively. Therefore, the growth restriction factor is enhanced upon increasing the Sc content. Furthermore, the Sc element’s enrichment at the solid/liquid interface tends to affect the diffusion of Al, inhibit the supplementation of Al in the liquid phase, and behave as a barrier to defer the growth of the primary α-Al matrix [37].

On the other hand, recent studies reported [11,19] that no compounds were formed with Mg and Sc, while Al₃Sc-phase particles tended to form at a eutectic reaction temperature in Al alloys due to the low solid solubility of Sc in the α-Al matrix. In addition, the solubility of Sc in solid solution is decreased by the existence of Mg in Al alloys [16]. Thus, Al₃Sc-phase particles tend to form in the alloy melt as the solidification temperature decreases. Based on the solidification theory [20], the heterogeneous nucleation potency of secondary-phase particles is closely related to the interfacial energy of the particle/matrix interface. The smaller the lattice misfit between particles and the α-Al matrix, the lower the interfacial energy, and stronger heterogeneous nucleation potency of particles is achieved [19]. The degree of lattice misfit (δ) is proposed to analyze the lattice misfit between the α-Al matrix and nucleated particles, as follows [38]:

δ = (α₁ − α₀)/α₀(3)

where α₁ and α₀ are lattice constants of secondary-phase particles and the α-Al matrix, respectively. Superior and moderate heterogeneous nucleation effects are obtained with a degree of lattice misfit below 6% and between 6% and 12%, respectively [39]. It is noticed that the lattice constant of Al₃Sc-phase particles is close to that of the α-Al matrix with a similar crystal structure and consistency in orientation [19]. Thus, significant grain refinement is achieved due to the formation of primary Al₃Sc-phase particles. However, primary Al₃Sc-phase particles tend to form only when Sc addition reaches eutectic composition [18]. Thus, a limited grain refinement effect is achieved in the CREed AM alloy with Sc In addition, further grain refinement effect is obtained in the CREed AM–0.5Sc alloy, which is also related to the recrystallization process during CRE.

4.2. Grain Refinement Induced by Al₃Sc-Phase Particles During Recrystallization

After solidification, the alloys undergo three stages of shear deformation during CRE, which are friction shear deformation, angular shear deformation, and extended shear deformation. The accumulative shear strain increases as the shear deformation proceeds. The total equivalent strain is calculated to be 3.8, and the primary Al₆(Mn, Fe)-phase particles at the microscale are cracked and refined to the nanoscale during CRE [29]. Therefore, the high equivalent strain during CRE contributes to the uniform distribution and refinement of Al₃Sc-phase particles in the AM–0.5Sc alloy, as presented in Figure 4. Moreover, continuous dynamic recrystallization (CDRX) is regarded as the main grain refinement mechanism during CRE [40]. The driving force of CDRX (F_d) is in direct proportion to the dislocation density (ρ) in Al alloys, which is presented as follows [41]:

F_d = 0.5Gb²ρ(4)

ρ = ρ_Al + ρ_sp = ρ_Al + 6f_vS/d_spb(5)

where G and b are the shear modulus and Burger’s vector of the α-Al matrix and ρ_Al and ρ_sp are the dislocation density of the α-Al matrix and the dislocation density increment induced by secondary-phase particles during deformation, respectively. f_v and d_sp are the volume fraction and the average diameter of secondary-phase particles, respectively. S is the shear strain during deformation. With the accumulative shear deformation during CRE, Al₃Sc-phase particles in the AM–0.5Sc alloy are cracked and refined into nanoscale, contributing to the enhancement of f_v and S. Meanwhile, d_sp is decreased due to the formation of nano-sized Al₃Sc-phase particles. Furthermore, the nano-sized Al₃Sc-phase particles distributed within the grains have a strong inhibition effect on the dislocations in motion during deformation, which are beneficial for the enhancement of the driving force of CDRX and the refinement of grains in the alloy.

In addition, the Zener drag force (F_z), induced by the secondary-phase particles, is crucial to the migration velocity of grain boundaries and the grain size of Al alloys [42]. Thus, it is essential to analyze the Zener drag force generated by the nano-sized Al₃Sc-phase particles, which is described as follows [41]:

Fz = 3f_vγ/d_sp(6)

where γ is the grain boundary energy of the α-Al matrix. It is noticed that nano-sized Al₃Sc-phase particles have strong impeding and pinning effects on the migration and growth of grain boundaries of Al alloys [19]. Meanwhile, the Zener drag force is enhanced, which is attributable to the f_v increment and d_sp decrement of Al₃Sc-phase particles from microscale to nanoscale during CRE. Thus, a superior grain refinement effect is achieved in the CRE AM–0.5Sc alloy in comparison to the alloy without Sc modification.

4.3. Strengthening Mechanisms of the CREed AM Alloy with Sc Modification

Refined grains with enhanced mechanical properties are achieved in the CREed AM alloy with Sc modification. Thus, it is imperative to investigate the relationship between YS enhancement and microstructural variation and reveal the strengthening mechanisms of the CREed AM(–Sc) alloys. The yield strength of AM alloys is determined by the externally applied stress, which is crucial for driving dislocations to overcome various obstacles, such as Mg solutes, grain boundaries, nanoscale secondary-phase particles, and dislocation structures within the alloys [23]. Moreover, the Al matrix is strengthened by Mg solutes by directly pinning moving dislocations and indirectly dragging dislocations by diffusing to dislocation cores to form solute clusters [24]. Furthermore, the strength increment (∆σ_SS) induced by the above interaction approaches is expressed as follows [43]:

∆σ_SS = k_SSCⁿ(7)

where k_SS and n are material-related parameters and C is the Mg solute content in the Al crystal lattice. In this work, similar content and a fully solid solution of Mg atoms are achieved in the CREed AM(–Sc) alloys. Hence, the improvement of YS induced by Mg solute strengthening is insignificant. In addition, the strength enhancement generated by dislocation pileups at grain boundaries (∆σ_GB) is estimated using the Hall–Petch equation [17]:

∆σ_GB = k_HPd^−0.5(8)

where k_HP and d are the Hall–Petch coefficient and the average grain size of the CREed AM(–Sc) alloys, respectively. It is noticed that the obvious decrement in the average grain size from 52.6 μm to 2.4 μm is achieved in the CREed AM–0.5Sc alloy in comparison to the CREed AM alloy. Thus, a superior grain boundary strengthening effect is obtained in the CREed AM–0.5Sc alloy.

Furthermore, second-phase strengthening (∆σ_Or) induced by the interaction between nanoscale Al₃Sc-phase particles and gliding dislocations via the Orowan mechanism is estimated using the following equation [16]:

∆σ_Or = Gb(f_v/2π)^0.5/3d_sp(9)

Hence, the second-phase strengthening effect is enhanced due to the formed Al₃Sc-phase particles from micron-sized to nano-sized in the CREed AM–0.5Sc alloy in contrast with the CREed AM alloy with no secondary-phase particle precipitation.

Meanwhile, the dislocation structures formed during CRE have an inhibition effect on the slip of mobile dislocations and provide a superior dislocation strengthening (∆σ_Dis) effect in the alloy, which is expressed via the Taylor equation, as shown below [29]:

∆σ_Or = αGbMρ^0.5(10)

where α, b, and M are the constant and the magnitudes of the Burgers vector and the Taylor factor, respectively. Massive studies have reported that nanoscale Al₃Sc-phase particles and Mg solutes have a strong pinning effect on dislocations in motion. Hence, the dislocation density in the CREed AM–0.5Sc alloy is enhanced, and the dislocation strengthening effect is improved. The physical meaning and values of the symbols utilized in the calculation Equations (7)–(10) are summarized in Table 3. The calculated YS strength increment is 131.1 MPa, which is close to the experimentally measured YS increment (120.3 MPa). The difference between the measured and calculated results may result from the interaction between strengthening mechanisms [29]. Therefore, the strengthening mechanisms of the CREed AM–0.5Sc alloy are grain boundary strengthening, second-phase strengthening, and dislocation strengthening.

During tensile loading, plastic deformation of Al alloys is realized via dislocation slip [46]. However, the nanoscale Al₃Sc-phase particles have a strong pinning effect on dislocations during motion, leading to the difficulty in dislocation slip [47]. Thus, inhomogeneous deformation still occurs between the Al₃Sc-phases and the Al matrix. With the increase in strain, the stress concentration is further accumulated. As the accumulative stress concentration and inhomogeneous deformation exceed the critical value, microcracks form between the Al₃Sc-phases and the Al matrix. The microcracks gradually grow, propagate, and connect with the neighboring microcracks. Consequently, circular cracks around Al₃Sc-phase particles are formed. Therefore, the CREed Al–5Mg–xSc alloys with nano-sized Al₃Sc-phase particles had low elongation in comparison to the CREed Al–5Mg alloys.

5. Conclusions

In this work, the microstructural evolution and mechanical property response of the AM alloy prepared through CRE and Sc modification were investigated. The grain refinement and strengthening mechanisms induced by nanoscale Al₃Sc-phase particles in the alloys were discussed. The following conclusions were drawn:

(1) An obvious grain refinement effect was achieved in the CREed AM alloy, with the average grain size reduced from 52.6 μm to 11.2 μm as the Sc content increased from 0 to 0.3 wt%. A further reduction in the average grain size was obtained as the Sc content reached 0.5 wt%, representing a 95.4% decline in comparison to that of the alloy without Sc modification.

(2) The primary Al₃Sc-phase particles formed during solidification acted as heterogeneous nucleation sites for α-Al matrix. Meanwhile, the nanoscale Al₃Sc-phase particles achieved during CRE enhanced the driving force of the CDRX and the Zener drag force, resulting in a superior grain refinement effect.

(3) The UTS, YS, and hardness were enhanced with a loss in ductility as the amount of Sc increased from 0 to 0.5 wt%. Grain boundary strengthening, second-phase strengthening, and dislocation strengthening were the main strengthening mechanisms of the CREed AM alloy modified with Sc.

Author Contributions

Conceptualization, B.Y. and M.G.; methodology, B.Y.; software, W.L.; validation, W.L.; formal analysis, X.L.; investigation, B.Y. and X.L.; resources, D.Y.; data curation, W.L.; writing—original draft preparation, B.Y.; writing—review and editing, B.Y. and M.G.; visualization, D.Y.; supervision, D.Y. and M.G.; project administration, D.Y.; funding acquisition, M.G. All authors have read and agreed to the published version of the manuscript.

Data Availability Statement

The original contributions presented in this study are included in the article. Further inquiries can be directed to the corresponding authors.

Conflicts of Interest

Authors Bowei Yang, Wenyue Liu, Xin Liu are employed by the company Ansteel Beiing Research Institute Co., Ltd. The remaining authors declare that the research was conducted in the absence of any commercial or financial relationships that could be construed as a potential conflict of interest.

Footnotes

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Figures and Tables

Figure 1. OM images and statistical average grain size result of the CREed AM alloys modified with different Sc additions: (a) AM alloy; (b) AM–0.1Sc alloy; (c) AM–0.3Sc alloy; (d) AM–0.5Sc alloy; (e) statistical average grain size result.

Figure 2. SEM micrographs and EDS results of the CREed AM(–0.5Sc) alloys: (a) AM alloy; (b) AM–0.5Sc alloy; (c) EDS results of point A marked in (a); (d) EDS results of point B marked in (b).

Figure 3. XRD analysis results of the CREed AM alloys with different Sc content.

Figure 4. TEM micrographs of (a) the AM alloy and (b) the AM–0.5Sc alloy fabricated through CRE.

Figure 5. Engineering stress–strain curves of CREed AM alloys with different Sc additions.

Table 1

Chemical compositions of the Al–5Mg–xSc alloys fabricated through CRE (in wt%).

Alloy	Mg	Sc	Al
Al–5Mg (AM)	4.94	–	Bal.
Al–5Mg–0.1Sc (AM–0.1Sc)	4.97	0.09	Bal.
Al–5Mg–0.3Sc (AM–0.3Sc)	4.92	0.31	Bal.
Al–5Mg–0.5Sc (AM–0.5Sc)	4.96	0.47	Bal.

Table 2

Tensile properties and hardness results of the AM alloys prepared through CRE and Sc modification.

Alloy	UTS (MPa)	YS (MPa)	EL (%)	Hardness (HV)
AM	257.6 ± 1.7	99.8 ± 2.1	52.8 ± 4.2	69.8 ± 0.7
AM–0.1Sc	331.3 ± 2.4	133.5 ± 1.2	41.8 ± 1.1	101.4 ± 1.5
AM–0.3Sc	389.2 ± 3.9	181.3 ± 3.2	32.0 ± 2.5	126.5 ± 1.9
AM–0.5Sc	456.4 ± 3.6	220.1 ± 2.8	24.8 ± 1.9	143.7 ± 2.2

Table 3

Physical meaning and values of the symbols utilized in the calculation Equations (7)–(10).

Symbol	Description	Values	Unit	Ref.
k SS	Constant	13.8	MPa/(wt%)	[43]
n	Constant	1	-	[43]
C	Mg solute content	4.94/4.96	wt%	This work
k HP	Hall–Petch coefficient	0.14	MPa·m−0.5	[44]
d	Average grain size	52.6/2.4	μm	This work
G	Shear modulus	28.9	GPa	[45]
b	Burgers vector	0.286	nm	[23]
d sp	Average diameter	20.5	nm	This work
f v	Volume fraction	0.065	-	This work
α	Constant	0.3	-	[29]
M	Mean orientation factor	3.06	-	[29]
ρ	Dislocation density	1.32 × 1013/9.47 × 1013	m−2	This work

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