Achieving Superplasticity in Ultrafine-Grained Mg-9Li Alloy via Dual-Phase Microstructure Optimization

Abstract

In this study, high toughness and superplastic deformability were achieved in Mg-9Li alloys through dual-phase microstructure optimization. Solid solution (SS) and equal channel angular pressing (ECAP) treatments were employed to refine the alloy’s microstructure. The effects of these treatments on room-temperature and low-temperature high-strain-rate superplasticity were systematically investigated under varying microstructural conditions. Results demonstrate that the SS-ECAP alloy exhibits outstanding superplasticity at room temperature and remarkable high-strain-rate deformation capability, achieving a maximum fracture elongation of 602.1%. Grain refinement and reduced dislocation density promote uniform void nucleation under high strain. Calculations of the strain rate sensitivity index (m-value) and activation energy (Q) reveal that the superplastic behavior in the SS-ECAP state is predominantly governed by grain boundary sliding facilitated by grain boundary diffusion. These findings provide critical insights into advancing the superplastic forming technology of Mg-9Li alloys.

1. Introduction

Magnesium (Mg) and its alloys have emerged as promising candidates for lightweight structural applications in aerospace, automotive, and advanced engineering fields owing to their exceptional combination of low density, high specific strength, and favorable plasticity [1,2,3]. Among these, Mg-Li alloys stand out as next-generation materials due to their superior plasticity, positioning them as ideal choices for high-performance applications [4,5,6]. However, it should be noted that they also exhibit poor corrosion resistance. In practical applications, anti-corrosion processes (such as surface treatment techniques) need to be incorporated to enhance their corrosion resistance. For instance, in the aerospace field, Mg-Li alloys are used in rocket hatch covers, waveguides, heat-insulation panels, etc. [7]; in the military field, the United States applied Mg-Li alloys to the disk of the aiming device in TOW (optical—tracking wire-guided missile generator tube), and when the temperature changes, this disk can both ensure the contact gap and reduce the mass [8,9].

The phase composition of Mg-Li alloys changes significantly with varying lithium content. When the lithium content is below 5.7 wt.%, the alloy adopts a single-phase α-Mg structure (hexagonal close-packed, HCP). At intermediate lithium levels (5.7–10.3 wt.%), a dual-phase microstructure comprising α-Mg and β-Li (body-centered cubic, BCC) emerges. Beyond 10.3 wt.% Li, the alloy transitions to a single-phase β-Li structure. Notably, the β-Li phase exhibits a lower critical resolved shear stress (CRSS), and increasing β-Li content significantly enhances plasticity, particularly in dual-phase systems [10,11,12].

However, as-cast Mg-9Li alloys still fall short of meeting practical mechanical requirements. Additionally, they exhibit poor high-temperature performance, as the alloy softens at relatively low temperatures, which restricts their application in areas with significant thermal exposure (e.g., engine components, and aerospace systems). This necessitates microstructural optimization through advanced processing techniques [13]. For instance, Peng et al. [14] demonstrated that homogenization at 200 °C for 4 h induced minimal microstructural changes in Mg-7.7Li-2.93Al-0.36Y alloy. Conversely, Ji et al. [15] reported that solution treatment at 350 °C for 4 h maximized the strength of Mg-8Li-3Al-2Zn-0.5Y alloy (YS = 226 MPa, UTS = 292 MPa) but resulted in poor ductility. In contrast, Maurya et al. [16] observed complete dissolution of the Al-Li phase in Mg-9Li-7Al-1Sn alloy after 400 °C solution treatment and subsequent aging at 100 °C, achieving remarkable hardness improvements (61% for α-Mg, 136% for β-Li). These findings collectively indicate that optimal heat treatment temperatures differ between multicomponent (200 °C) and binary Mg-9Li (400 °C) alloys.

While solution treatment enhances compositional homogeneity, its capacity to improve mechanical properties remains limited. To address this, Liu et al. [17] applied equal channel angular pressing (ECAP) to Mg-8Li-1Al alloy, refining both α-Mg and β-Li grains to ~500 nm and significantly enhancing mechanical performance. Achieving superplasticity in Mg alloys requires stable ultrafine grains (<10 µm), abundant high-angle grain boundaries, and distributed precipitates [18,19,20,21,22,23,24,25]. ECAP, as a severe plastic deformation technique, is particularly effective in fulfilling these microstructural criteria [26,27]. For magnesium alloys, superplasticity can be exhibited when the microstructure contains stable fine equiaxed grains (<10 µm), a high density of large-angle grain boundaries, and abundant precipitate [28].

Previous studies have demonstrated that Mg-Li alloys exhibit superplastic behavior at 250–400 °C. Although light, Mg-Li alloys generally have lower mechanical strength and hardness compared to other magnesium alloys. They are more prone to plastic deformation under stress and may not withstand heavy structural loads. For example, Chen et al. [29] achieved 306.6% elongation in ECAP-processed Mg-10.73Li-1.49Al-0.52Y alloy at 350 °C. Mehrabi et al. [30] further reduced grain size to 4 µm via multi-directional forging, enabling superplasticity at 225–300 °C with strain rates of 3.3 × 10⁻³ to 1.3 × 10⁻¹s⁻¹. According to Langdon’s criteria [31], true superplasticity (elongation ≥ 400%) requires a strain rate sensitivity index m = 0.5, while m = 0.3 corresponds to near-superplastic behavior. Taleff et al. [32] reported m = 0.5 and activation energy Q = 65 kJ/mol in rolled Mg-9Li, confirming grain boundary sliding controlled by boundary diffusion as the dominant mechanism. Critically, superplasticity in fine-grained materials demands grain sizes < 10 µm, temperatures > 0.5 Tm (where Tm is the melting point), and strain rates of 10⁻⁴–10⁻² s⁻¹ [33], with grain refinement enabling high-strain-rate superplasticity (HSRSP) [24].

Despite these advances, no systematic study has investigated the low-temperature (273 K–373 K) high-strain-rate superplasticity of solution-treated and ECAP-processed Mg-9Li alloys, particularly regarding deformation mechanisms. This work therefore aims to optimize the dual-phase microstructure of Mg-9Li through combined solution treatment and ECAP, characterize its superplastic behavior under room-temperature and low-temperature conditions, and elucidate the underlying deformation mechanisms through microstructural and mechanical analyses.

2. Experimental Details

The raw material Mg-9Li alloy ingots were purchased from Jiangsu Limagnesium Aviation Materials Co., Ltd. (Wuxi, China). The preparation of Mg-9Li alloy ingots involved first mixing high-purity magnesium and high-purity lithium at a mass ratio of 91:9 for melting. A suitable proportion of protective gas was introduced to prevent contamination by oxygen, water vapor, and other impurities under high-temperature and high-vacuum conditions. Sufficient stirring was performed during the alloy smelting process to ensure no gas bubbles formed within the ingot, thereby obtaining high-quality Mg-Li alloy billets. The molten Mg-9Li alloy was then cast into square ingots. The composition of the Mg-9Li alloy ingots was analyzed using a GNR S3 spark emission spectrometer (GNR, Milan, Italy). The results are summarized in

Table 1. Analyzed composition results of Mg-9Li alloy.

Elemental	Mg	Li	Zn	Si	Ce	Fe	Mn
content (wt.%)	Bal.	8.809	0.014	0.012	0.011	0.010	0.024

. Based on the mass ratio of magnesium to lithium, the alloy was confirmed as Mg-9Li (wt.%) and is hereafter referred to as Mg-9Li.

The as-cast Mg-Li alloy was cut into blocks (20 mm × 20 mm × 44 mm) using an electrical discharge wire-cutting machine and subsequently subjected to solution treatment. The samples were heated to 848 K, held at this temperature for 4.5 h, and then rapidly quenched in an ice–water mixture. A schematic of the solution treatment process is illustrated in Figure 1a.

The alloy was further processed via equal channel angular pressing (ECAP, RD-ECAP technique). The ECAP equipment is independently set up (Nanjing, China). As shown in Figure 1b, the ECAP mold featured a 90° extrusion channel angle with a square cross-sectional area (20 mm × 20 mm). The as-cast alloy was processed by ECAP at 473 K to produce the Cast-ECAP alloy. The solution-treated Mg-9Li alloy underwent multiple ECAP passes at room temperature to obtain the SS-ECAP alloy.

Microstructural characterization was performed on samples extracted from the center of the as-cast and solution-treated alloys. For the Cast-ECAP and SS-ECAP alloys, microstructures were observed along the transverse plane (TP, i.e., ED × ND). Fracture morphology and post-tensile surfaces of Mg-9Li tensile samples were examined using optical microscopy (OM) and scanning electron microscopy (SEM). In OM and SEM images, the x-axis corresponds to the extrusion direction (ED). To assess deformation uniformity, tensile-tested samples were placed on a black contrast surface, arranged by strain rate, and photographed to evaluate necking behavior.

Phase composition and internal microstructure were analyzed using a Bruker D8 X-ray diffraction (XRD) apparatus (Bruker, Billerica, MA, USA). The test conditions for XRD were as follows: Cu target Kα radiation with a wavelength λ = 1.5406 Å, tube current of 40 mA, voltage of 40 kV, scanning spot area diameter of 2 mm, scanning range from 25° to 90°, and scanning speed of 5°/min. The internal microstructure was analyzed using an FEI Tecnai G2 transmission electron microscope (TEM) (Thermo Fisher Scientific, Waltham, MA, USA), respectively. Both XRD and TEM characterizations were aligned along the ED.

Tensile tests were conducted on a Suns UTM4294X electronic universal testing machine (SUNS, Shenzhen, China). Superplasticity test specimens had a gauge length of 1.5 mm and a thickness of 2 mm (Figure 1c). Testing conditions included strain rates of 1 × 10⁻³, 7.5 × 10⁻⁴, 5 × 10⁻⁴, 2.5 × 10⁻⁴, and 1 × 10⁻⁴ s⁻¹ at room temperature (298 K), as well as strain rates of 1 × 10⁻³, 2.5 × 10⁻³, 5 × 10⁻³, 7.5 × 10⁻³, and 1 × 10⁻² s⁻¹ under low-temperature heating (323 K, 348 K, and 373 K). To ensure accuracy, at least three parallel samples were tested for each condition.

3. Results

3.1. Microstructural Evolution of Mg-9Li Alloy

As shown in Figure 2a, scanning electron microscopy (SEM) images of the as-cast Mg-9Li alloy reveal a dual-phase microstructure consisting of a gray α-Mg matrix (hexagonal close-packed, HCP) and black β-Li phase (body-centered cubic, BCC), which is further confirmed by X-ray diffraction (XRD) patterns in Figure 2e and the prior literature [3]. The α-Mg phase in the as-cast state exhibits a coarse needle-like morphology, with average grain dimensions of approximately 400 µm in length and 50 µm in width. After solution treatment (SS) at 848 K for 4.5 h (Figure 2c), the α-Mg phase retains its needle-like structure but undergoes significant refinement, with most grains reduced to lengths below 80 µm and widths less than 10 µm.

Subsequent equal channel angular pressing (ECAP) processing further modifies the microstructure. For the Cast-ECAP alloy (Figure 2b), α-Mg grains become more thoroughly intermixed with the β-Li phase, though the degree of phase uniformity remains lower compared to the solution-treated alloy. The α-Mg grains in Cast-ECAP also exhibit enhanced alignment along the extrusion direction (ED). In contrast, the SS-ECAP alloy (Figure 2d) demonstrates superior grain refinement and microstructural homogeneity, with both α-Mg and β-Li phases uniformly distributed along the ED.

3.2. Superplastic Behavior at Room Temperature

Figure 3 and

Table 2. UTS, TYS, UEL, and TEL of Mg-9Li alloys during tensile testing at 298 K.

Alloys	ε (s−1)	UTS (MPa)	TYS (MPa)	UEL (%)	TEL (%)
Cast	1.0 × 10−4	69.5 ± 4.2	52.3 ± 3.1	13.1 ± 0.7	95.0 ± 5.7
	2.5 × 10−4	88.8 ± 6.3	65.7 ± 4.6	15.1 ± 1.0	87.2 ± 6.1
	5.0 × 10−4	105.2 ± 7.4	75.0 ± 5.4	13.8 ± 0.9	66.9 ± 4.7
	7.5 × 10−4	114.7 ± 8.1	81.5 ± 5.4	16.3 ± 1.2	58.3 ± 4.1
	1.0 × 10−3	120.5 ± 9.4	86.8 ± 6.9	19.2 ± 1.5	50.4 ± 4.0
Cast-ECAP	1.0 × 10−4	97.4 ± 5.0	75.4 ± 3.7	14.9 ± 0.8	110.2 ± 5.6
	2.5 × 10−4	128.6 ± 7.6	100.3 ± 5.9	18.1 ± 1.1	95.3 ± 5.5
	5.0 × 10−4	157.7 ± 9.5	120.1 ± 7.1	22.3 ± 1.3	90.8 ± 5.6
	7.5 × 10−4	177.0 ± 10.8	141.1 ± 8.6	17.3 ± 1.1	82.1 ± 5.0
	1.0 × 10−3	191.3 ± 13.2	154.4 ± 10.5	22.8 ± 1.4	70.3 ± 4.7
SS	1.0 × 10−4	70.1 ± 2.9	52.4 ± 2.1	10.1 ± 0.4	100.2 ± 4.1
	2.5 × 10−4	90.9 ± 4.2	71.3 ± 3.7	14.2 ± 0.7	89.3 ± 4.4
	5.0 × 10−4	107.7 ± 5.4	87.9 ± 4.4	17.0 ± 1.0	79.2 ± 4.0
	7.5 × 10−4	117.2 ± 6.3	96.1 ± 4.9	20.2 ± 1.0	69.7 ± 3.6
	1.0 × 10−3	123.2 ± 7.3	104.1 ± 6.3	15.0 ± 0.9	59.6 ± 3.5
SS-ECAP	1.0 × 10−4	104.8 ± 3.2	68.5 ± 2.1	29.1 ± 0.9	144.2 ± 4.5
	2.5 × 10−4	140.9 ± 5.5	105.9 ± 3.9	17.8 ± 0.7	130.3 ± 5.1
	5.0 × 10−4	175.4 ± 7.0	137.0 ± 5.3	21.2 ± 0.8	122.6 ± 4.9
	7.5 × 10−4	198.8 ± 8.2	159.3 ± 6.1	33.9 ± 1.4	113.9 ± 4.7
	1.0 × 10−3	217.1 ± 10.6	192.7 ± 9.4	21.0 ± 1.0	101.9 ± 5.8

summarize the tensile properties of Mg-9Li alloys at 298 K under varying strain rates (1 × 10⁻⁴ to 1 × 10⁻³ s⁻¹). The SS-ECAP alloy exhibits the highest mechanical performance, achieving an ultimate tensile strength (UTS) of 104.8 MPa and a total elongation (TEL) of 144.2% at a strain rate of 1 × 10⁻⁴ s⁻¹. As the strain rate increases to 1 × 10⁻³ s⁻¹, the UTS increases to 217.1 MPa, while the average TEL reduces to 101.9%. In comparison, the Cast-ECAP alloy shows inferior performance, with UTS values ranging from 97.4 MPa (1 × 10⁻⁴ s⁻¹) to 191.3 MPa (1 × 10⁻³ s⁻¹) and average TEL values between 70.3% and 110.2%. This limitation is attributed to rapid dislocation accumulation and saturation during tensile testing, which restricts work-hardening capacity.

For Figure 3a–e, generally, the strength of the four samples increases while the elongation decreases with the increase in strain rate. Scientifically, this reveals the influence law of strain rate on the room-temperature mechanical properties of Mg-9Li alloys, providing a basis for selecting appropriate strain rates in practical applications. Take Figure 3e as an example, it shows the stress–strain curves of four samples (Cast, SS, Cast-ECAP, SS-ECAP) at a strain rate of 1 × 10⁻³ s⁻¹ under 298 K. At the same strain rate, the strength and elongation of SS-ECAP are significantly improved, indicating its superior comprehensive mechanical properties.

The big differences in Figure 3f arise from significant disparities in microstructural states due to distinct processing techniques. Magnesium alloys, with their HCP structure, have poor symmetry and are highly texture sensitive. For example, high-pressure torsion (HPT) produces a basal texture, while equal-channel angular pressing (ECAP) generates a fiber texture. These textures significantly affect mechanical properties. These differences can be primarily attributed to the significant disparities in microstructural states induced by distinct processing techniques. Magnesium alloys, with their HCP structure, exhibit poor symmetry and are highly influenced by texture. For instance, the high-pressure torsion (HPT) process tends to produce a basal texture, while the equal-channel angular pressing (ECAP) process generates a fiber texture. These different textures significantly affect the mechanical properties, highlighting the importance of processing techniques in regulating alloy properties. The enhanced superplasticity of the SS-ECAP alloy is mechanistically linked to its refined grain structure (~360 nm) and reduced dislocation density, which promotes grain boundary sliding and uniform void nucleation under strain.

3.3. High Strain Rate Superplastic Behavior Under Heating Conditions

The tensile behavior of Mg-9Li alloys was systematically evaluated at elevated temperatures (323 K, 348 K, and 373 K) under strain rates ranging from 1 × 10⁻³ to 1 × 10⁻² s⁻¹ (Figure 4, Figure 5 and Figure 6, Tables S1 and S2).

For Figure 4a–e, Figure 5a–e and Figure 6a–e,g,h, generally, the strength of the four samples decreases while the elongation increases with the rise in temperature. Scientifically, it clarifies the influence of temperature on the high-temperature mechanical properties of Mg–9Li alloys, offering theoretical support for high-temperature applications. As shown in Figure 4a, Figure 5a and Figure 6a, which display the stress–strain curves of four samples (Cast, SS, Cast-ECAP, SS-ECAP) at a strain rate of 1 × 10⁻³ s⁻¹ under 323 K, 348 K, and 373 K, respectively. Different from the room-temperature (298 K) condition, compared with the other three states of alloys, the strength of SS-ECAP decreases significantly and the elongation increases remarkably at the same strain rate, indicating its excellent superplasticity at high temperatures, which has guiding significance for developing high—temperature superplastic forming processes.

The SS-ECAP alloy demonstrates exceptional HSRSP, achieving a maximum average TEL of 602.1% at 373 K and 1 × 10⁻³ s⁻¹ (

Table 3. UTS, TYS, UEL, and TEL of Mg-9Li alloys during tensile testing at 373 K.

Alloys	ε (s−1)	UTS (MPa)	TYS (MPa)	UEL (%)	TEL (%)
Cast	1.0 × 10−3	17.9 ± 1.1	12.8 ± 1.0	9.6 ± 0.6	162.3 ± 9.9
	2.5 × 10−3	24.7 ± 1.7	20.0 ± 1.4	16.2 ± 1.1	146.6 ± 10.1
	5.0 × 10−3	30.7 ± 2.1	25.6 ± 1.8	13.5 ± 0.9	135.3 ± 9.5
	7.5 × 10−3	35.0 ± 2.5	27.9 ± 2.0	22.3 ± 1.6	125.1 ± 8.9
	1.0 × 10−2	38.2 ± 3.0	32.0 ± 2.5	11.3 ± 1.0	106.9 ± 8.4
Cast-ECAP	1.0 × 10−3	22.6 ± 1.2	16.2 ± 0.8	24.7 ± 1.3	212.2 ± 10.8
	2.5 × 10−3	31.5 ± 1.9	24.6 ± 1.5	17.9 ± 1.1	194.9 ± 11.5
	5.0 × 10−3	40.2 ± 2.4	30.8 ± 1.8	24.7 ± 1.5	185.4 ± 11.1
	7.5 × 10−3	46.2 ± 2.8	35.7 ± 2.2	18.1 ± 1.1	172.5 ± 10.5
	1.0 × 10−2	50.8 ± 3.5	40.4 ± 2.8	16.5 ± 1.1	151.1 ± 10.4
SS	1.0 × 10−3	16.4 ± 0.7	12.4 ± 0.7	26.4 ± 1.1	173.1 ± 7.1
	2.5 × 10−3	23.1 ± 1.1	19.7 ± 1.0	14.8 ± 0.8	155.2 ± 7.6
	5.0 × 10−3	29.0 ± 1.5	24.3 ± 1.2	14.3 ± 0.7	145.3 ± 7.3
	7.5 × 10−3	33.0 ± 1.7	28.2 ± 1.4	17.0 ± 0.9	133.9 ± 6.8
	1.0 × 10−2	36.9 ± 2.2	32.6 ± 1.9	18.5 ± 1.1	125.3 ± 7.4
SS-ECAP	1.0 × 10−3	13.7 ± 0.4	8.1 ± 0.3	75.2 ± 2.3	602.1 ± 18.7
	2.5 × 10−3	23.6 ± 0.9	14.3 ± 0.6	98.4 ± 3.8	472.1 ± 18.4
	5.0 × 10−3	33.3 ± 1.3	24.2 ± 1.0	52.0 ± 2.1	395.2 ± 15.8
	7.5 × 10−3	39.7 ± 1.6	31.1 ± 1.3	42.2 ± 1.7	295.2 ± 12.1
	1.0 × 10−2	44.7 ± 2.2	35.4 ± 1.7	46.0 ± 2.3	255.0 ± 12.5

). This transition reflects enhanced dynamic recrystallization (DRX) and dislocation annihilation mechanisms at higher temperatures. The Cast-ECAP alloy shows moderate improvements, reaching a maximum average TEL of 212.2% at 373 K (Table 3). However, its performance remains inferior to the SS-ECAP alloy due to coarser grains and less uniform phase distribution. Figure 4f, Figure 5f, and Figure 6f similar to Figure 3f, present comparisons with the previous literature. Once again, it emphasizes that the differences in microstructures and textures caused by different processing techniques affect the properties, reinforcing the previous viewpoint. Critical analysis of strain rate sensitivity (m-value) and activation energy (Q) reveals that the SS-ECAP alloy achieves true superplasticity (m > 0.5, Q = 61.55 kJ/mol) at 348–373 K, governed by grain boundary sliding controlled by boundary diffusion.

3.4. Fracture Surface Morphology and Void Evolution

SEM analysis of fracture surfaces (Figure 7 and Figure 8) provides insights into deformation mechanisms. The as-cast alloy exhibits ductile fracture characteristics at low strain rates (1 × 10⁻⁴ s⁻¹), with equiaxed dimples and tear ridges (Figure 7a–d). High-magnification images (Figure 7b,d) reveal dynamically recrystallized spherical grains (~1 µm in size) on fracture surfaces. At higher strain rates (1 × 10⁻³ s⁻¹) or lower temperatures (298 K), as demonstrated in regions such as Figure 7c,d,k,l, the fracture morphology exhibits a prevalent mixed mode. This arises from the combined effect of plastic deformation and stress concentration within the alloy. The coexistence of tearing ridges and dimples indicates that the alloy undergoes simultaneous shearing and stretching. For the SS-ECAP alloy, fracture morphology evolves with strain rate: samples tested at 1 × 10⁻⁴ s⁻¹ (Figure 7m,n) display fine dimples and tearing ridges, indicative of a more ductile deformation mode. This is attributed to the sufficient time available for dislocation movement and grain boundary sliding at lower strain rates, which are characteristic of ductile deformation processes.

Similarly, in Figure 8e,g, the fracture surface presents a complex combination of features, implying mixed-mode deformation under higher strain rates. For the SS-ECAP alloy, fracture morphology varies with strain rate: samples tested at 1 × 10⁻⁴ s⁻¹ (Figure 7m,n) show fine dimples and tear ridges, while those at 1 × 10⁻³ s⁻¹ (Figure 7o,p) exhibit coarser features. At 373 K (Figure 8m–p), intergranular separation dominates, generating abundant voids that accommodate large strains. This uniform void distribution, combined with ultrafine grains (~360 nm), prevents premature coalescence and enables the alloy’s unprecedented average TEL of 602.1%.

4. Discussion

4.1. Analysis of Superplastic Deformation Mechanism of Mg-9Li Alloy Based on m and Q Values

Based on the changes in the tensile curves of Mg-9Li alloys and the study of microstructural evolution during superplastic deformation, the strain rate sensitivity m value and the deformation activation energy Q value during superplastic deformation were calculated and compared. This provides an in-depth understanding of the superplastic deformation mechanisms of as-cast and solution-treated Mg-9Li alloys.

The tensile peak stress values of the alloy at different strain rates under the same temperature serve as the flow stress values for superplastic deformation. The relationship between these stress values and the tensile strain rate is used to calculate the strain rate sensitivity m value during the superplastic deformation process. Additionally, the flow stress values at different temperatures but the same strain rate are used as a function of the inverse temperature to calculate the deformation activation energy Q-value. Both m and Q c can be calculated using the Arrhenius equation, as shown in the following Formula (1) [28]:

$$\dot{\mathsf{\epsilon }}=\mathrm{A}{\mathsf{\sigma }}^{\mathrm{n}}{{\mathrm{e}}^{-}}^{{\displaystyle \frac{\mathrm{Q}}{\mathrm{RT}}}}$$

(1)

where A is the constant of the material, σ is the value of the rheological stress on the tensile curve of the alloy, n is the value of the stress exponent of the alloy (n = 1/m), R is the gas constant of the material with a value of about 8.31 J·mol⁻¹K⁻¹, and T is the absolute temperature of the material. The m value can be directly calculated using Formula (2):

$$\mathrm{m}=\partial \ln \mathsf{\sigma }/\partial \ln \mathsf{\epsilon }$$

(2)

where ε is the corresponding tensile strain rate for each sample. Therefore, the value of m can be derived by fitting the value of the rheological stress as a function of the tensile strain rate, the slope of which is the m value. Once the m value is calculated, the Q value can be calculated by the following Equation (3) [36]:

$$\mathrm{Q}=\mathrm{nR}\left({\displaystyle \frac{\partial \ln \mathsf{\sigma }}{\partial \left(1000/\mathrm{T}\right)}}\right)$$

(3)

Among them, $\partial \ln \mathsf{\sigma }/\partial (1000/\mathrm{T})$ is calculated from the fitted values of the rheological stress values as a function of the temperature inverse.

Based on the changes in the tensile curves of Mg-9Li alloy in four states—Cast, SS, Cast-ECAP, and SS-ECAP—and the study of microstructural evolution during superplastic deformation, a detailed analysis was conducted on the strain rate sensitivity (m value) and activation energy (Q value) of the alloy during superplastic deformation. This analysis aimed to explore the superplastic deformation mechanisms of the four states (Cast, SS, Cast-ECAP, and SS-ECAP) of the Mg-9Li alloy.

The curves for calculating the m-value of Mg-9Li alloy in the Cast, SS, Cast-ECAP, and SS-ECAP states at 298 K, 323 K, 348 K, and 373 K are shown in Figure 9. The calculated numerical results are shown in

Table 4. Values of m for different strain rates.

T/K	ε (s−1)	m
		Cast	SS	Cast-ECAP	SS-ECAP
298 K	1.0 × 10−4	0.3163	0.3206	0.3350	0.4280
	2.5 × 10−4	0.2990	0.3004	0.3239	0.4160
	5.0 × 10−4	0.2910	0.2994	0.3110	0.4115
	7.5 × 10−4	0.2810	0.2893	0.3080	0.4010
	1.0 × 10−3	0.2710	0.2744	0.3042	0.3590
323 K	1.0 × 10−3	0.3320	0.3444	0.3580	0.5000
	2.5 × 10−3	0.3260	0.3374	0.3450	0.4360
	5.0 × 10−3	0.3110	0.3234	0.3380	0.4160
	7.5 × 10−3	0.3080	0.3204	0.3310	0.4090
	1.0 × 10−2	0.2930	0.3054	0.3140	0.4030
348 K	1.0 × 10−3	0.3350	0.3445	0.3600	0.5600
	2.5 × 10−3	0.3280	0.3370	0.3550	0.4590
	5.0 × 10−3	0.3150	0.3310	0.3490	0.4210
	7.5 × 10−3	0.3090	0.3280	0.3410	0.4180
	1.0 × 10−2	0.2950	0.3220	0.3290	0.4020
373 K	1.0 × 10−3	0.3400	0.3620	0.3720	0.6530
	2.5 × 10−3	0.3340	0.3510	0.3690	0.5850
	5.0 × 10−3	0.3260	0.3470	0.3550	0.4550
	7.5 × 10−3	0.3200	0.3390	0.3530	0.4450
	1.0 × 10−2	0.3070	0.3340	0.3240	0.4310

. With the increase in tensile temperature, the general trend of the alloy’s m-value is an increase, which to some extent indicates that the alloy becomes more prone to achieving superplasticity at higher temperatures. The m-values of SS-ECAP alloy are 0.4103, 0.4329, 0.4520, and 0.5135, with all m-values exceeding 0.3, and at 373 K, the m-value exceeds 0.5. This is consistent with the alloy’s observed elongation results.

As shown in Figure 10, existing research suggests that, when the strain rate sensitivity (m-value) is in the range of approximately 0.0 to 0.3, the elongation of magnesium alloys is generally small and less than 100%, making it difficult to achieve superplasticity. This is because the deformation mechanisms involved in tensile testing are typically dominated by the generation and slip of dislocations. During dislocation movement and slip, dislocations are prone to pile up, and once the pile-up reaches a certain extent, stress concentration occurs, leading to crack formation and premature fracture. When the strain rate sensitivity is in the range of 0.3 to 0.5, the tensile elongation of magnesium alloys can generally reach over 100%, thus achieving near-superplasticity. When the strain rate sensitivity exceeds 0.5, true superplasticity can be achieved. The m-values of the four microstructural states at 323 K, 348 K, and 373 K under different strain rates (Figure 10) fall into different stages, mainly focusing on the true superplasticity stage. The SS-ECAP alloy can achieve true superplasticity at 348 K and 373 K under low strain rates.

Based on the listed flow stress values and the corresponding temperatures during tensile testing, the curves for calculating the activation energy (Q-value) of the cast, SS, Cast-ECAP, and SS-ECAP alloys at different temperatures and strain rates are plotted and shown in Figure 11 and

Table 5. Q-values for different strain rates.

T/K	ε (s−1)	Q/(kJ/mol)
		Cast	SS	Cast-ECAP	SS-ECAP
298 K	1.0 × 10−3	84.22	85.39	86.17	88.37
323 K	1.0 × 10−3	82.56	80.84	77.61	73.32
	2.5 × 10−3	81.95	79.33	77.06	67.50
	5.0 × 10−3	81.52	78.84	76.64	64.99
	7.5 × 10−3	80.95	78.36	76.16	63.35
	1.0 × 10−2	80.48	77.72	75.56	62.16
348 K	1.0 × 10−3	82.40	79.35	76.12	70.23
	2.5 × 10−3	81.79	77.87	75.58	64.65
	5.0 × 10−3	81.36	77.39	75.16	62.25
	7.5 × 10−3	80.79	76.91	74.69	60.68
	1.0 × 10−2	80.32	76.29	74.11	59.53
373 K	1.0 × 10−3	79.07	76.00	74.52	61.55
	2.5 × 10−3	78.49	74.59	74.0	56.91
	5.0 × 10−3	78.08	74.12	73.59	54.80
	7.5 × 10−3	77.53	73.67	73.12	53.42
	1.0 × 10−2	77.08	73.07	72.56	52.41

.

For Figure 11, generally, for the four different states of samples, as the strain rate increases, the slope K of the fitting curves in the figure gradually decreases. Scientifically, as the strain rate increases, the deformation process becomes more rapid, leading to a change in the dominant deformation mechanism. The decrease in the K value indicates that the sensitivity of the activation energy to temperature reduces with the increase in strain rate. This reflects the intrinsic relationship between the activation energy and the strain rate, helping to understand how different strain rates affect the deformation mechanism of the alloy. It provides a theoretical basis for optimizing the processing parameters to achieve the desired mechanical properties. According to the fitting results and the previous discussion of the m-values for the cast and solid-solution alloys, the average Q-values for the cast and solid-solution Mg-9Li alloys are calculated to be 80.27 kJ/mol, 76.91 kJ/mol, 75.08 kJ/mol, and 61.55 kJ/mol, respectively [19].

In Mg-Li alloys [37], when the strain activation energy (Q-value) is less than 60 kJ/mol, it indicates that grain boundary self-diffusion occurs during superplastic deformation; when the Q-value is between 60 and 100 kJ/mol, it suggests that grain boundary diffusion is dominant during superplastic deformation; when the Q-value is between 100 and 130 kJ/mol, it indicates that lattice diffusion occurs in β-Li; and when the Q-value exceeds 130 kJ/mol, it implies that dislocation creep plays a key role in the superplastic deformation process

As shown in Figure 12, the superplastic deformation of Cast, SS, Cast-ECAP, and SS-ECAP Mg-9Li alloys is primarily governed by grain boundary sliding controlled by grain boundary diffusion. As the strain rate increases, the Q-value gradually decreases, and similarly, as the temperature increases, the Q-value also decreases. Different from the other three states (Cast, SS, Cast-ECAP), the changing trend of SS-ECAP with the variation in strain rate and temperature is more significant. Scientifically, the SS-ECAP alloy exhibits a more uniform microstructure and smaller grain size, which leads to a higher superplastic elongation. (The finer microstructure makes the grain boundary diffusion and sliding more active and sensitive to external conditions). At medium to low strain levels, the superplastic deformation mechanism of SS-ECAP alloys is controlled by grain boundary self-diffusion, where grain boundary sliding can facilitate dislocation motion inside grains and also help coordinate smaller strains through the formation of substructures between neighboring grains. On the other hand, after a certain degree of grain boundary sliding, the deformation is further accommodated by the formation of cavities (Figures S3 and S4), where the generation of cavities is mainly driven by the separation of grain boundaries. The more significant change in trend of SS-ECAP highlights its unique superplastic deformation mechanism, indicating that its microstructure enables more complex interactions between grain boundary sliding, dislocation movement, and diffusion under different conditions. This is of great significance for understanding the essence of its excellent superplasticity and has guiding value for its application in practical processing.

4.2. Analysis of Enhancing Plastic Deformation Ability Based on Ultrafine Grains

The SEM microstructure of Mg-9Li alloy before and after solid solution treatment shows significant differences (Figure 2), and the XRD patterns after ECAP treatment (Figure 2e) also show changes in peak height. However, unlike the solid-solution alloy, the diffraction peaks shift slightly to the left and the peak width increases. These changes in peak height reflect the different crystal planes revealed during the test, focusing on different textures. The shift and broadening of the peaks indicate that, after ECAP processing, the microstructure of the alloy has been significantly refined. Due to the large amount of plastic deformation during ECAP, a considerable amount of strain is accumulated within the alloy, leading to severe lattice distortion and an increase in dislocation density. These factors affect the interplanar spacing in the alloy, thus influencing the peak width and position in the XRD pattern.

To analyze the grain size of the samples more clearly, Figure 13a–c presents the TEM images of the Cast-ECAP alloy. The results of single-crystal electron diffraction calibration show that the phase with dark contrast in the images is α-Mg grains. The grain size distribution in Figure 13d indicates an average grain size of about 550 nm. The fitting curve (black) of the frequency distribution histogram is based on the Log Normal function from Peak Functions, with the iterative algorithm being Levenberg–Marquardt [38]. The computed u_g(q₀) the geometric mean, represents the average grain size (with standard deviation). The cumulative frequency distribution curve (red) is based on the Boltzmann function from Origin Basic Functions, also using the Levenberg–Marquardt iterative algorithm. The calculated d_i corresponding grain size when calculating the grain percentage at i%.

The grain shapes are nearly equal. Low-pass ECAP processing generally involves the progressive evolution of Mg grains and β-Li phases from coarse structures to submicron or even nanoscale structures. High-pass ECAP processing, on the other hand, tends to gradually reduce the strain in the alloy, leading to grain growth. In cast alloys, the boundary between low-pass and high-pass ECAP processing is typically around 12 passes. After 12 passes of ECAP processing, the recrystallization process in the alloy is nearly complete, resulting in nearly fully recrystallized grains.

Figure 13d–f further demonstrates the TEM microstructure of SS-ECAP alloy. Compared to the Cast-ECAP alloy, the grain size in this alloy is significantly smaller, the grains are more uniform, and they exhibit a nearly equiaxed shape. From the enlarged TEM image of the alloy, it is clear that the grain morphology is equiaxed. In the grain size distribution plot (Figure 13h), the average grain size is about 360 nm, with the calculation method being the same as described earlier. Smaller grain sizes, more grain boundaries, and higher grain boundary energy facilitate deformation.

By combining SEM and TEM microstructural images, models of the different microstructural states of Cast, Cast-ECAP, SS, and SS-ECAP Mg-9Li alloys are illustrated in Figure 14a–d. These models provide an intuitive analysis of the superplastic deformation mechanisms. Comparing the microstructural models of Cast, SS, Cast-ECAP, and SS-ECAP samples allows a direct assessment of how the dual-phase microstructure of Mg-Li alloys influences mechanical properties. The differences in grain size, morphology, and distribution are significant. Solid-solution treatment plays a critical role in achieving a more uniform distribution of the α-Mg phase, while ECAP processing significantly refines the grain size. Among the four states, the SS-ECAP alloy exhibits the smallest grain size, reduced to an ultrafine level of approximately 360 nm.

Smaller grain sizes, with a higher fraction of grain boundaries, effectively strengthen the material through the Hall-Petch relationship [39,40]: σ_y = σ₀ + kd^−1/2. Here, grain boundaries act as barriers to dislocation motion, thereby enhancing the strength. On the other hand, changes in grain size also affect the material’s ability to undergo plastic deformation. Superplasticity occurs under specific deformation conditions due to the activation of internal deformation mechanisms, including grain boundary sliding, dynamic recrystallization, and the annihilation of dislocations. A unified model of creep and superplasticity based on grain size differences is shown in Figure 14e,f. In this model, d represents the grain size, and λ denotes the average subgrain size. Outside the superplasticity range, dislocation climb occurs, leading to subgrain formation within grains. When d < λ (Figure 14e), dislocations transfer from one grain to another without forming subgrains. Dislocations are easily generated and eliminated, resulting in minimal grain size changes during deformation, characteristic of superplastic deformation. Conversely, when d > λ (Figure 14f), dislocations transferring between grains lead to the formation of subgrains. Subgrain formation causes dislocation pile-up, generating internal stress, which alters grain orientation and shape [31]. This increases flow stress, indicating hardening behavior rather than superplasticity.

Figure 14g,h illustrates the relationship between grain size and flow behavior under superplastic conditions [41]. Figure 14g shows the relationship between strain rate and flow stress during superplastic deformation, with the superplasticity region occurring at moderate strain rates. For different grain sizes, the range of flow stress remains comparable, which is essential for achieving significant superplasticity. However, materials with finer grains demonstrate superplasticity at higher strain rates. Regarding superplastic elongation (Figure 14h), materials with smaller grains achieve higher elongation at elevated strain rates. This effect is due to the accelerated attainment of peak elongation as grain size decreases.

For the SS-ECAP alloy, the reduction in grain size to approximately 360 nm leads to an increased fraction of grain boundaries. Under low-temperature heating conditions, the deformation activation energy Q of the SS-ECAP alloy is 61.55 kJ/mol, indicating that grain boundary diffusion dominates the grain boundary sliding mechanism, significantly enhancing plastic deformation. Additionally, the shortened development and growth time of cavities further contributes to achieving higher elongation rates at these elevated strain rates.

5. Conclusions

The SS-ECAP Mg-9Li alloy demonstrates exceptional superplastic performance, achieving a maximum total elongation (TEL) of 602.1% at 373 K and a strain rate of 1 × 10⁻³ s⁻¹. This breakthrough is attributed to the synergistic effects of dual-phase microstructure optimization and ultrafine grain refinement (~360 nm) achieved through sequential solution treatment and equal channel angular pressing (ECAP). Key findings include:

• The SS-ECAP alloy exhibits an average TEL of 144.2% at 298 K, with uniform deformation and no necking. This is mechanistically linked to reduced dislocation density and enhanced grain boundary sliding enabled by ultrafine grains.
• At 373 K, the alloy achieves 602.1% elongation under a strain rate of 1 × 10⁻³ s⁻¹, exhibiting excellent deformation capability.
• Strain rate sensitivity (m > 0.5) and activation energy (Q = 61.55 kJ/mol) calculations (Table 4 and Table 5) reveal that grain boundary sliding controlled by grain boundary diffusion is the primary deformation mechanism. The α-Mg/β-Li dual-phase interface further enhances strain accommodation by suppressing void coalescence (Figure 7 and Figure 8).
• TEM analysis (Figure 13) confirms that the SS-ECAP process refines grains to ~360 nm, significantly smaller than Cast-ECAP (~550 nm) while promoting equiaxed morphology and grain-size uniformity.

These results establish a novel pathway for designing Mg-Li alloys with low-temperature HSRSP capabilities, critical for applications in aerospace and automotive lightweighting. Future work should explore scaling up the SS-ECAP process and evaluating fatigue resistance under cyclic loading.