Effects of Aging Processes on the Dynamic Impact Mechanical Behavior of Mg-Gd System Alloys

Abstract

Exploring the effect of the magnesium alloy aging process on dynamic impact performance could plays an important role in the application of magnesium alloy in automotive lightweighting. In this work, the effects of single-stage, two-stage, and reverse two-stage aging processes on the dynamic mechanical properties of Mg-8.5 Gd-3 Y-0.5 Zr alloy were studied by means of SEM analysis, hardness testing, a quasi-static compression experiment, and SHPB. The results show that the compressive strength of the materials after single-stage, two-stage, and reverse two-stage aging treatments is improved to different degrees compared with that of the alloys in the extruded state. Due to the generation of dynamic precipitation with semi-annular distribution during SHPB, the compressive strength of the reverse two-stage aging alloys reached an excellent 761 MPa, while the two-stage aging alloys had more dynamic precipitation phases at the strain rate of 3500 s⁻¹, resulting in a compressive strength of 730 MPa, which is superior to that of the aluminum alloys used in a wide range of automotive applications. The results of this study can provide a reference for the application of Mg-Gd magnesium alloys under dynamic loading.

1. Introduction

Magnesium alloy has been widely used in the field of automobile manufacturing due to its lower density, higher specific strength and other advantages [1,2,3,4]. However, magnesium alloy itself is not strong, which largely affects its application. With the in-depth study of magnesium alloys, researchers have found that adding rare earth elements to magnesium alloys can enhance the comprehensive properties of magnesium alloys [5,6,7,8,9]. Among them, Mg-Gd alloy has attracted wide attention from scholars at home and abroad due to its excellent mechanical properties. For example, Yu et al. [10] extruded Mg-Gd-Y-Zn-Zr alloy and found that the compressive strength of the alloy in its extruded state reaches 371 MPa. Mg-Gd-Y-Zn-Zr alloys were produced by Homma et al. [11] with the tensile strength of the specimens up to 542 MPa. In summary, magnesium alloys are anticipated to be a new generation of high-strength light-alloy engineering materials.

Mg-Gd alloys are typical of the aging precipitation-strengthened type of alloys. For instance, Wang et al. [12] prepared Mg-Gd-Y-Ag-Zr alloy, and the strength of the alloy after peak aging was up to 448 MPa. Mg-12 Gd-2 Y-1 Zn-Mn alloys were prepared by Su et al. [13] and, after extrusion and ageing, the alloys reached an excellent σt of 509 MPa.

However, at present, single-stage aging for harsh service environments cannot meet engineering demands. In the study of aluminum alloys, it has been found that two-stage aging can effectively improve the comprehensive mechanical properties of aluminum alloys [14,15]. At the same time, in other series of magnesium alloys, research on two-stage aging has also been attempted. For example, the mechanical properties of Mg-12 Gd-1 Er-1 Zn-0.9 Zr alloys under two-stage aging and single-stage aging was explored by Jia [16]. It was found that, after two-stage aging treatment, the material further improved its strength, with the tensile strength reaching 549 MPa, while there was no significant decrease in plasticity. Li et al. [17] investigated and explored the effects of different aging treatments on AZ63 magnesium alloy and found that, compared with single-stage aging alloys, the two-stage aging and reverse two-stage aging alloys had a greater effect on the organization properties of the alloy.

Magnesium alloys are commonly used in automotive components, and in the course of automotive use, automotive components have the probability of suffering dynamic impacts, such as emergency braking and crashes. Enhancing the dynamic impact performance of the material can increase the safety of the automotive use process. Yu et al. [18,19] found that the impact resistance of Mg-Gd-Y alloys was significantly improved after aging treatment, and their dynamic compressive strength could reach 510 MPa. At the same time, other researchers have found that the aging treatment is critical for enhancing the dynamic mechanical properties of the alloys [20,21]. Therefore, it is necessary to use two-stage aging and reverse two-stage aging processes to enhance the alloys’ properties. However, the dynamic impact behavior of Mg-Gd system alloys in the two-stage aging state and reverse two-stage aging state has not been reported up to now. In this paper, the Mg-8.5 Gd-3 Y-0.5 Zr alloy is taken as the research object to explore the impact of the aging process on its dynamic mechanical properties. The findings can serve as a reference for the practical application of this alloy under dynamic loading conditions.

2. Experimental Procedures

The composition of the experimental material used in this work was Mg-8.5 Gd-3 Y-0.5 Zr alloy; Gd, Y and Zr elements were added in the form of Mg-Gd, Mg-Y and Mg-Zr intermediate alloys, respectively, and Mg elements were added in the form of pure magnesium (≥99.96%). The raw materials were distributed according to the alloy composition, melted in an iron crucible, refined, left to stand and then poured to obtain 30 mm-diameter ingots. The actual composition of the alloy was 88.07% Mg, 8.56% Gd, 2.84% Y, 0.50% Zr (mass fraction). The extrusion process was as follows: a Mg-8.5 Gd-3 Y-0.5 Zr magnesium alloy ingot billet with the dimensions of Φ30 mm × 70 mm was used as the material, preheated at 400 °C for 30 min; the diameter of the bar was 8 mm, the extrusion cylinder was 32 mm, the extrusion ratio was 16, and the extrusion rod moved at a speed of 30 mm/min. All the test samples were made into cylindrical bars with diameters of 8 mm and heights of 4 mm by wire-cutting. The test was carried out in three aging processes, and the heat treatment was carried out in a Sx3-1-7 box-type resistance furnace. The specific aging process was as follows: single-stage aging at 225 °C × 12 h, two-stage aging at 150 °C × 8 h + 225 °C × 12 h, reverse two-stage aging at 225 °C × 12 h + 150 °C × 8 h, and the heat preservation was cooled at room temperature after the cooling conditions. The hardness test was performed using an HV-1000 micro Vickers hardness tester from Shjingmi (Shanghai, China), the loading time was 10–14 s, the load was set to 9.8 N, each specimen was measured 5–10 times, and the average value was taken. The compressive properties of the materials under different conditions were tested using a WDW-50 compression testing machine from Kangyuan (Jinan, China) at a compression rate of 2 mm/min.

In this work, a split Hopkinson pressure bar (SHPB) device was used to perform dynamic compression tests. The compression surface was polished before compression and the loading direction was parallel to the extrusion direction. The firing pin struck the input rod and generated an incident wave, which propagated to the specimen and caused it to become damaged and deformedamage and deformation. At this point, the strain gauges on the input and output rods recorded the incident, reflected and transmitted waves, and converted them using the SHPB software from Donghua (Jiangsu, China, Version 1.0). Subsequently, the fracture and microstructure morphology of the samples was observed by SEM. Before the observation, the specimens were roughly and finely ground, polished using a specimen polishing machine, and then corroded with 6% nitric acid alcohol. SEM observations were undertaken utilizing a MERLIN Compact from Zeiss (Oberkochen, Germany), and the surface and compression fracture microstructures were analyzed using an energy spectrometer (EDS) from Zeiss (Oberkochen, Germany) at an accelerating voltage of 10 kV.

3. Experimental Results

3.1. Initial Microstructure Analysis

The microstructure of Mg-Gd-Y-Zr in the extruded state is shown in Figure 1. The microstructure of the alloy mainly consisted of Mg matrix and a small amount of rare earth eutectic phase distributed at the grain boundaries. As a result of the extrusion process, the alloy grain size was not uniform, while the grain size was small, measured to be about 13.36 μm by the intercept method, while the grain shape was approximately isometric. The mass fraction ratio of Gd and Y at matrix A was approximately 8.5:3, and the element distribution was close to the nominal composition of the alloy, while the solid solubility of Zr in the Mg matrix was relatively low and aggregated into the Zr-rich particles of approximate spherical shape, as shown in B, and the mass fraction of Zr was as high as 48.15%; the second phase C was an RE-rich phase, and the mass fraction of RE was close to 90%, in which the volume of Gd and Y fraction ratio was approximately 2:1.

3.2. Aging Analysis

The aging strengthening effect of Mg-Gd-Y-Zr alloy is remarkable, and the aging temperature is usually set at 200 °C–250 °C [22]. Therefore, this paper chose to carry out aging treatment at 200, 225 and 250 °C; the aging hardening curve is shown in Figure 2. When the alloy was aging at 200 ℃, the hardness of the specimen rapidly increased within 6 h, reaching 100.8 HV, but, after this initial period, the hardness began to increase slowly. It took 16 h to reach peak aging hardness. The aging temperature for optimal hardness performance was 225 °C, with a peak aging hardness of 110.5 HV in 12 h. When the aging temperature was increased to 250 °C, the time needed to reach the peak aging hardness decreased to 8 h. However, the peak aging hardness decreased to 106.3 HV. As the aging temperature rose, the aging reaction accelerated, leading to a shorter time to reach the peak hardness. Meanwhile, the hardness showed a tendency to increase first and then go down, which was because the aging microstructure grows excessively when the aging temperature is too high. In summary, 225 °C was more suitable as the aging temperature of this magnesium alloy.

According to the literature [23], the Mg-Gd-Y-Nd-Zr alloy produces a diffuse and homogeneous precipitate when aged at 150 °C, and more precipitated phases may be obtained at 8 h. Therefore, 150 °C × 8 h was chosen as the low-temperature aging in two-stage and reverse two-stage aging, while the optimal single-stage aging process of 225 °C × 12 h was used as the high-temperature aging in two-stage and reverse two-stage aging. The aging hardening curve of magnesium alloy in different aging states is shown in Figure 2b. In the two-stage aging process, low-temperature aging at 150 °C was carried out first, resulting in the alloy hardness reaching 99.7 HV after 8 h. Subsequently, high-temperature aging at 225 °C was carried out, and the hardness increased to 117.3 HV at 20 h, which is a significant improvement over the samples with the single-stage aging. The hardness of the single-stage aging sample decreased after 12 h, while the two-stage aging sample continued to undergo low-temperature aging at 150 °C after the high-temperature aging at 225 °C, resulting in a further increase in hardness to 117.9 HV at 20 h, an increase of 6.7% compared to the single-stage aging sample, which is close to the hardness of the two-stage aging sample. Therefore, other testing methods are required to compare the differences between the two aging processes.

Figure 3 shows the microstructure of the material with different aging treatments.

Table 1. The distribution of elements at each point (wt%).

Elements	A	B	C	D	E	F
Mg	56.93	78.58	9.79	85.97	16.7	79.98
Gd	29.56	14.34	56.96	8.96	47.81	10.95
Y	13.15	6.45	31.48	4.61	34.6	8.26
Zr	0.36	0.63	1.77	0.47	0.89	0.81

shows the distribution of elements at each point. Compared with the extruded alloy, the second phase of the single-stage aging sample is uniformly dispersed and has grown significantly. After keeping the material at 225 °C for 12 h, aging precipitation occurred, in which the number of precipitated phases gradually increased and first refined and then grew, and most of them were precipitated in a discontinuous manner at the grain boundaries. As can be seen in Table 1, the precipitated phases at the grain boundaries had a higher content of rare earth elements and a larger volume, and the number density of the precipitates at the grain boundaries was larger than that in grain interior. This is due to the larger lattice distortion energy around the grain boundaries, which generates sufficient energy for the diffusion of atoms [24].

In Figure 3b, the number of second phases in two-stage aging alloys increases significantly relative to single-stage aging alloys. This is due to the fact that, in the two-stage aging process, aging at lower temperatures results in uniformly dispersed precipitates; aging at higher temperatures promotes the growth of precipitated phases [25]. Therefore, the two-stage aging process is superior to the single-stage aging process in terms of alloy aging precipitation.

After the inverse two-stage aging treatment, there was no significant increase in the number of precipitated phases in the material, but the precipitated phases appear to be aggregated. This is due to the fact that a certain amount of ageing precipitation can occur in the high-temperature ageing treatment, and the subsequent low-temperature ageing process promotes further precipitation around the ageing precipitated phases [26].

Figure 4 shows the mechanical properties of the alloy with different aging treatments. As shown in Figure 4, compared with the extruded alloy, the plasticity of the material decreased after single-stage aging treatment, and the strain rate was reduced to 10.5%, but the yield strength and compressive strength increased by 16.2% and 6.1%, respectively. In summary, the aging treatment effectively improved the mechanical properties of the alloy. The reverse two-stage aging treatment resulted in a smaller increase in plasticity, the strain rate increased to 10.8%, and there was an increase in compressive strength of 8.6%, compared to single-stage aging. Compared with single-stage aging, the compressive strength of the material after two-stage aging treatment increased by 17.3%, the plasticity also increased to some extent, and the strain rate increased to 12.2%. The two-stage aging and reverse two-stage aging treatments effectively improved the mechanical properties of the Mg-Gd-Y-Zr alloy, which reached 610 MPa and 565 MPa, respectively, and are much higher than those of the aluminum alloys that are widely used in automotive applications [27].

The stress–strain curves of alloys under various aging conditions prepared via SHPB tests are shown in Figure 5. The effects of different aging treatments on the properties of the alloys are shown in Figure 6. The stresses of both extruded and two-stage aging alloys were first elevated at high speed, and then slowed down to enter a smooth deformation stage after reaching a certain stress, until they reached the highest strength point, then the curve decreased sharply, and gradually increased with increasing strain rate. The curve of the single-stage aging and the reverse two-stage aging alloys changed with a similar trend to the extruded and the two-stage aging alloys at the high strain rate, but, at low strain rates, the stress increase did not slow down significantly, and rose at high speed until the maximum compressive strength was reached.

3.3. Dynamic Mechanical Behavior Analysis

The strength of the extruded alloy at a 1500 s⁻¹ strain rate was 356 MPa, and the strength grew to 626 MPa with the growth of the strain rate to 3500 s⁻¹. The strength of the alloy in the single-stage aging state at a strain rate of 1500 s⁻¹ was 513 MPa. With a growth in strain rate to 3000 s⁻¹, the compressive strength rose to 708 MPa. The compressive strength of the alloy in the two-stage aging state increased with the increasing strain rate and reached 730 MPa at a 3500 s⁻¹ strain rate. The reverse two-stage aging alloy had the highest compressive strength at a 2500 s⁻¹ strain rate, reaching an impressive 761 MPa. However, as the strain rate grew further to 3500 s⁻¹, its compressive strength decreased to 692 MPa. At a strain rate of 3500 s⁻¹, the extruded-state alloy had the best plastic properties with a strain of 0.53, which was significantly higher than the other alloys under aging conditions. This resulted in the extruded alloy not fracturing at a strain rate of 3500 s⁻¹, while all other specimens fractured.

In summary, in the extruded state and the two-stage aging state, the patterns of change in the compressive strength of the alloy were similar. In these two cases, the compressive strength increased with the strain rate. However, in single-stage aging and reverse two-stage aging, the compressive strength initially increased and then decreased with the increase in strain rate. Among the different aging processes, the reverse two-stage aging showed the best dynamic compressive strength, reaching an excellent 761 MPa at a strain rate of 2500 s⁻¹, which is superior to that of 6061 aluminum alloy, which is widely used in automotive applications [28].

4. Discussion

4.1. The Effects of Aging Processes on Dynamic Impact Behaviors

The SEM and EDS images of the extruded-state alloy after dynamic impact are displayed in Figure 7. The EDS results showed that the particles at B were RE-rich phases, with the largest weight fraction of Gd elements, more than 49%, and 38% Y elements. Figure 7d shows that the smaller particles surrounded by it were about 20% elemental Gd and 15% elemental Y. According to the study by Tang [24], the EDS results of the smaller particles are similar with the data for dynamically precipitated particles. During the deformation process, the hindering effect of the dynamically precipitated particles on the dislocation motion is beneficial. There is a greater hindering effect and superior macroscopic mechanical properties as their number increases. In summary, this indicates that dynamic precipitation occurs in extruded alloys during dynamic impact, and the second phase and the surrounding dynamically precipitated phases can effectively hinder deformation and enhance the macroscopic mechanical properties of the materials.

The extruded alloy has a certain number of second-phase particles at a 1500 s⁻¹ strain rate, and the fine second-phase particle phase surrounds the larger square phase, which is due to the hindering effect of the second-phase dislocations, bringing about dislocations aggregated in the surrounding area. With increasing strain rate, the dynamic precipitation phase of the alloy significantly increases at a strain rate of 2500 s⁻¹ with increasing strain rate. However, at a strain rate of 3500 s⁻¹, no dynamic precipitation occurs due to the high strain rate, and the fine second-phase particles are almost invisible.

The SEM and EDS images of the single-stage aging state alloy after dynamic impact are shown in Figure 8. As shown in Figure 8e, the results of the energy spectrum at A show that the Gd element was about 20%, and the Y element was about 15%, indicating that dynamic precipitation also occurred in the single-stage aging state alloy. The alloy with single-stage aging state had a good precipitation effect under 2500 s⁻¹ strain rate conditions, with numerous smaller particles scattered around the Zr-rich particles. The quantity of second-phase particles was notably higher compared to the extruded-state alloy. However, under 3500 s⁻¹ strain rate conditions, the precipitated phase diminished, and dynamic precipitation hardly occurred. This phenomenon leads to a decrease in the compressive strength of the alloy.

The SEM and EDS images of the two-stage aging state alloy after dynamic impact are shown in Figure 9. The dynamic precipitation of the two-stage aging state alloy increased significantly compared with the single-stage aging state alloy, showing a trend of growth with increasing strain rate, and also occurred under the condition of 3500 s⁻¹ strain rate, while dynamic precipitation of the single-stage aging and extrusion-state alloy did not occur in this condition, which made the compressive strength of the two-stage aging state alloy in this condition significantly better than the other samples. At the same time, the dynamic precipitation was distributed around the grain boundaries and tended to be concentrated.

The SEM and EDS images of the reverse two-stage aging state alloy after dynamic impact are shown in Figure 10. The reverse two-stage aging state alloy was similar to the single-stage aging state alloy in that the number of dynamic precipitation phase particles increased and then decreased with increasing strain rate. There was only a small amount of dynamic precipitation at the 1500 s⁻¹ strain rate, while with the increase in strain rate, there was significant dynamic precipitation at the 2500 s⁻¹ strain rate, where fine dynamic precipitation phases were seen to converge around grain boundaries, while the number of particles increased dramatically and had a semi-annular distribution. This is due to the presence of hard particles, which can increase the dislocation movement resistance and gradually generate a dislocation line around the hard particles bending, and the excellent organizational properties of the reverse two-stage aging state alloys make the dislocation line at the generation of dense dynamic precipitation phases, further hindering dislocation movement and ultimately enhancing the compressive strength of the alloy.

During the dynamic impact process, the vast majority of the dynamic energy created is released as thermal energy, with a small portion stored in the material [29]. The adiabatic temperature rise during dynamic impact can be calculated from the absorbed energy $W$ per unit volume of the alloy as shown in Equation (1) [30]:

$$W=\int \sigma d\epsilon$$

(1)

where $\sigma$ is the stress and $\epsilon$ is the strain during dynamic compression.

The impact absorption characteristics of the alloys under various treatment conditions were analyzed based on the stress–strain curves, as presented in

Table 2. The absorption energy of alloys with different heat treatment conditions.

Treatment Condition	Extruded State			Single-Stage Aging			Two-Stage Aging			Reverse Two-Stage Aging
	1500 s−1	2500 s−1	3500 s−1	1500 s−1	2500 s−1	3500 s−1	1500 s−1	2500 s−1	3500 s−1	1500 s−1	2500 s−1	3500 s−1
$W$(MJ·m−3)	50.5	156.8	242.9	53.2	195.1	293.1	76.8	194.4	306.9	49.4	209.7	275.8

. The adiabatic temperature rise in a dynamic impact is calculated as shown in Equation (2) [31].

$$∆T={\displaystyle \frac{\beta }{\rho C_v}}\int \sigma d\epsilon$$

(2)

In this paper, where $\beta$ represents the conversion factor, $\beta$ = 1; where $\rho$ represents the density, $\rho$ = 1900 kg m⁻³; and where $C_v$ represents the specific heat capacity, $C_v$ = 10³ J kg⁻¹ °C⁻¹ [32].

Table 3. The adiabatic temperature rise of alloys with different heat treatment conditions.

Treatment Condition	Extruded State			Single-Stage Aging			Two-Stage Aging			Reverse Two-Stage Aging
	1500 s−1	2500 s−1	3500 s−1	1500 s−1	2500 s−1	3500 s−1	1500 s−1	2500 s−1	3500 s−1	1500 s−1	2500 s−1	3500 s−1
$∆T$(°C)	26.6	82.5	127.8	28	102.7	154.3	40.4	102.3	161.5	26	110.2	145.1

displays the calculated adiabatic temperature rise of the alloy for different treatment cases. In Table 3, $∆T$ increases as the strain rate increases. The overall increase in strength observed in the two-stage aging was higher compared to other state alloys; the reverse two-stage aging-state alloys exhibited the greatest rise and achieved the highest values when subjected to dynamic compression at a strain rate of 2500 s⁻¹. The good adiabatic temperature rise of the two-stage aged alloys resulted in good overall dynamic precipitation, while the volume fraction of the dynamically precipitated particles precipitated from the reverse two-stage aging-state alloy was higher than that of the other specimens at a strain rate of 2500 s⁻¹, which led to better compressive properties than the other specimens.

4.2. Fracture Mechanism

Under 3500 s⁻¹ strain rate conditions, no fracture occurred due to the better plastic properties of the alloys in the extruded state, while the rest of the samples showed fracture phenomena, and their fracture micro-morphology is illustrated in Figure 11. The figure reveals that the fracture surface consists of cleavage planes and dimples, and the fracture morphology of the specimens shows the characteristics of ductile–brittle mixed fracture dominated by ductile fracture. Under high-speed impact conditions, fracture was generated on the disintegration surfaces of different layers, followed by a gradual increase, which ultimately led to the obvious disintegration fracture of the alloy.

The fracture of the single-stage aging state sample mainly consists of dimples and cleavage planes. The dimples are distributed on the cleavage planes in the form of a network, and the depth of the dimples is shallow, with an obvious smooth tearing ridge. In the two-stage aging state, the alloy has a large piece of cleavage plane around the dimples, and the tearing ridges are significantly reduced. The size and depth of the dimples are larger than those of the single-stage aged alloy, and the number of dimples is higher, indicating that the plasticity of the alloy has been improved after two-stage aging. White second-phase particles can be found in the dimples. The fracture of the reverse two-stage aging alloy consists of tearing ridges, cleavage planes, and dimples distributed around the tearing ridges. The number of dimples is high, but they show a small and shallow morphology, and the dimples are smaller. The number and size of the tearing ridges are obviously increased compared with those of the previous two kinds of alloys, and they are spread over the entire fracture, which indicates that the fracture mode of the reverse two-stage aging alloy is mainly cleavage fracture. At the same time, the distribution of the dimples around the tearing ridge indicates that there is also a certain degree of ductile fracture during the fracture process.

5. Conclusions

This study primarily investigated the effect of different aging treatment conditions on the dynamic impact behavior of extruded Mg-Gd-Y-Zr alloy. In summary, the following conclusions can be drawn:

(1)

The mechanical properties of the Mg-Gd-Y-Zr alloy can be improved by single-stage, two-stage, or reverse two-stage aging processes. The compressive strength of the material in the single-stage aging state reached 520 MPa, and the compressive strength of the material in the reverse two-stage aging state increased by 8.6%. After two-stage aging treatment, the material had an optimum compressive strength of 610 MPa.

(2)

Dynamic impact properties can be improved by the aging process. The compressive strength of the material after single-stage aging treatment was significantly improved compared with that of the extruded alloy, reaching up to 708 MPa. The compressive strength of the alloy in the two-stage aging state improved with the strain rate up to 730 MPa. The reverse two-stage aging generated a superior strength of 761 MPa.

(3)

Dynamic precipitation occurred in the extruded state alloys and single-stage aging alloys at low strain rates but not at high strain rates, while dynamic precipitation occurred in both two-stage aging and reverse two-stage aging at 1500 s⁻¹–3500 s⁻¹ strain rates, in which the reverse two-stage aging showed a semi-annular distribution of fine dynamic precipitation phases at a 2500 s⁻¹ strain rate, which effectively improved the dynamic impact properties of the material.

(4)

Under 3500 s⁻¹ strain rate conditions, the extruded alloys did not fracture due to their better plastic properties, while the rest of the samples fractured, and the fracture morphology of the specimens shows the characteristics of ductile–brittle mixed fracture dominated by ductile fracture. Among the samples, the size of the dimples of the alloys in the two-stage aging state increased significantly.