Revealing the Interaction Between Dislocations and LPSO-Precipitates Structure in a Mg-Y-Al Alloy at Different Temperatures

Abstract

Precipitation strengthening represents a crucial strengthening approach in the realm of metals, with particular significance for magnesium. In this study, a complex LPSO–precipitate structure, which is constituted of the principal secondary phases in Mg rare earth (RE) alloys, namely the Long-Period Stacking Ordered (LPSO) phase and the aging precipitate, was successfully fabricated within a Mg-11Y-1Al alloy. Subsequently, an in-depth investigation was conducted regarding the interaction between dislocations and this LPSO–precipitate structure under varying temperature conditions. The findings revealed that, at room temperature (RT), the aging precipitates effectively hindered the movement of basal dislocations, and the activation of non-basal dislocations is rather difficult, resulting in the alloy’s high strength and low plasticity. When the temperature was elevated to 200 °C, although non-basal slip could be initiated, the LPSO–precipitate structure was capable of blocking both basal and non-basal slips. Consequently, the alloy still demonstrated high strength and low plasticity. As the temperature further increased to 250 °C, dislocations could cut through the aging precipitate particles, and the interior of the grains could provide partial deformation. Hence, the tensile elongation of the alloy was significantly enhanced, increasing from 4% to 12% as the temperature was elevated from 200 °C to 250 °C. These results suggest that the LPSO–precipitate structure still exerts a remarkable strengthening effect at 200 °C. When the temperature reaches 250 °C, the plasticity of the alloy is improved but its strength decreases. The research outcomes presented in this paper offer a novel perspective for the precise tailoring of mechanical properties through precipitation strengthening within Mg-RE alloys.

1. Introduction

Magnesium alloy is the lightest structural metal and has excellent energy-saving potential [1,2]. However, its low strength and poor ductility restrict the application area of Mg alloys. Alloying is the most simple and effective strengthening and toughening method for magnesium alloys [3,4]. Among magnesium alloy systems, rare earth (RE) magnesium alloys have attracted much attention because of their superior properties to non-RE alloys [5,6]. The alloying of RE elements could significantly improve the mechanical performance of magnesium alloys at both room temperature (RT) and elevated temperatures [7,8]. The mechanical properties of Mg-RE-based alloys are highly correlated to the distribution, morphology, and phase structure of precipitates. However, more in-depth and detailed research is needed to determine the addition and existence form of alloying elements in magnesium alloys according to the actual demand.

Precipitates are the products of “excess” alloying. By controlling the rare earth concentration in the matrix of Mg-RE-based alloys, different types of precipitates with different densities and distribution uniformities can be formed during the treatment process at different isothermal temperatures. When isothermally treated at relatively high temperatures (usually between 400 °C and 570 °C), the LPSO phase parallel to the basal plane can be formed in the ternary system of magnesium-rare earth-transition metals. These LPSO phases are classified into many types, such as 18R, 14H, 12R, 10H, etc. [9,10,11], according to the presence of a crystal structure; when isothermally treated at relatively low temperatures (usually between 125 °C and 250 °C), the age-hardening prismatic precipitates perpendicular to the basal plane will be formed in magnesium-rare earth-based magnesium alloys. These age-hardening precipitates are divided into many types, such as β″, β′, β, β_e, etc. [12,13,14,15], according to different age-hardening stages, among which the β′ phase along the prismatic plane is generally recognized as having the best strengthening effect [7,13,16].

Previous studies have shown that in magnesium alloys, the critical resolved shear stress (CRSS) for the activation of the basal slip system at room temperature is significantly higher than that of the non-basal slip system, and the gap between them decreases at high temperatures, allowing the activation of the basal slip to occur. The age-hardening precipitates can effectively prevent the basal dislocation slip at room temperature, and the resistance to the non-basal slip weakens at high temperatures, while the LPSO phase has a limited capacity to prevent dislocations at room temperature and exhibits a blocking effect on the non-basal slip at high temperatures [17,18]. Although the strengthening effects of the LPSO phase and the aging precipitates have been discussed in many previous studies and the general trend of the change of the difference in CRSS between the basal and the non-basal slip of the actual alloy with temperature can be calculated, it is still difficult to determine which phase has a better strengthening effect by calculating the specific CRSS values. On the other hand, although the individual strengthening effects of the LPSO phase and the aging precipitates can be roughly calculated, the actual effect of their composite structure on dislocations, especially the influencing behavior when the non-basal slip can be activated at high temperatures, remains unclear.

In order to more accurately evaluate the actual strengthening effects of the LPSO–precipitate structure at room temperature and higher temperatures and their influences on the deformation behavior in actual alloys, we used mechanical testing and TEM characterization to study the mechanical properties of T4 and T6 Mg-11Y-1Al alloys at different temperatures, as well as the corresponding interaction behaviors between dislocations and the LPSO–precipitate structure. The Mg-Y-Al system was chosen for this study because its as-cast fine-grained structure and good performance show great application potential [6,19]. In addition, the blocking behaviors of the LPSO–precipitate structure on dislocations at different temperatures are still unclear. Our work has discovered three completely different interaction behaviors of this composite structure on dislocations at room temperature, 200 °C, and 250 °C. Specifically, at room temperature, the precipitates solely block the basal dislocations; at 250 °C, the precipitates are directly cut through by the basal dislocations; and at 200 °C, the LPSO–precipitate structure blocks both the basal dislocations and the non-basal dislocations simultaneously. The results of this paper provide new insights into regulating the mechanical properties of Mg-RE alloys by preparing precipitates.

2. Materials and Methods

The specific process for preparing the alloy with the nominal composition of Mg-11Y-1Al by gravity casting is as follows: pure magnesium, Mg-25Y (mass percentage) master alloy, and pure aluminum are added into the crucible in different proportions. A mixed protective gas of carbon dioxide and sulfur hexafluoride is introduced. The melt is kept at 750 degrees Celsius for 30 min and then cast into a preheated mold. The actual chemical composition detected by Inductive Coupled Plasma Emission Spectrometry (ICP) is Mg-10.83Y-1.09Al (weight percentage). The casting methodology and heat treatment parameters of the Mg-11Y-1Al alloy can be sourced from a previous study [20]. In the context of the Mg-11Y-1Al alloy, the homogenization treatment parameters (as shown in Figure 1) were configured to maintain a temperature of 520 °C for a duration of 8 h (hereinafter referred to as T4), whereas the peak aging parameters were set to 225 °C for 10 days (hereinafter denoted as T6). The solution treatment was carried out in a box furnace, while the aging treatment was performed in an oil bath furnace. After being taken out, the samples were immediately cooled in water. The selection of the solution treatment parameters was based on previous studies [18,20]. These parameters can obtain a large amount of LPSO phase in the microstructure of this alloy. The aging treatment parameters were the peak aging parameters.

Tensile tests were conducted by employing dog-bone specimens with a diameter of 5 mm and a gauge length of 30 mm at a strain rate of ~5 × 10⁻⁴. At each temperature point, these tests were replicated on three individual samples. Energy Dispersive X-ray Spectroscopy (EDS) analysis and Scanning Electron Microscopy (SEM) examinations were carried out on a Phenom ProX apparatus. The microstructure morphology was photographed in the backscattered electron mode, while the fracture morphology was photographed in the secondary electron mode. The metallographic samples for SEM-EDS were sanded with 400, 1200, 3000, and 7000 mesh sandpaper, then polished with magnesium oxide suspension. The Transmission Electron Microscopy (TEM) samples were initially sliced into sheets with an approximate thickness of 0.5 mm and subsequently delicately polished to around 70 μm. Thereafter, specimens were prepared by punching out 3 mm disks from these sheets. These disks were then subjected to twin-jet electro-polishing within a 4% perchloric acid alcohol solution. Following the twin-jet electro-polishing procedure, the specimens were further processed by ion-milling at an energy level of 1 keV for a period of 30 min. The TEM observations were performed using a JEOL-2100F TEM operating at an accelerating voltage of 200 kV.

3. Results and Discussion

3.1. Microstructures of Mg-11Y-Al Alloys in T4 and T6 States

Figure 2 illustrates the microstructures of the alloys in T4 and T6 states. The grain sizes of both T4 and T6 Mg-11Y-Al alloy samples are approximately 40 μm. This alloy possesses fine grains of around 40 μm in the as-cast state and exhibits good thermal stability, with the grain size remaining unchanged after heat treatment [20]. As shown in Figure 2a (in the Backscattered Electron, i.e., BSE mode), the main second phases in the T4 state Mg-11Y-Al alloy microstructure consist of bright particle phases ranging from several microns to dozens of microns, coarse laths near the grain boundaries, and fine laths in the grain. Figure 2b reveals that the microstructure of the T6 state Mg-11Y-Al alloy is similar to that of the T4 state.

The point analysis results of the EDS in Figure 2c demonstrate that the particle phase is Al₂Y particles formed during the casting process. In addition, the lath phases with different thicknesses are all LPSO phases, which should be formed during the solid solution treatment process [21,22]. The XRD diffraction spectra of the alloys in the T4 and T6 states also indicate that the β′ phase will appear in the alloy microstructure after aging treatment. The precipitates generated during the aging process are hardly observable under the SEM. However, the TEM observation results in Figure 2e indicate that these precipitates are evenly distributed in the matrix. Moreover, the precipitates within the intervals of the LPSO laths are restricted by these laths, thereby forming a composite structure of LPSO and precipitates. The diffraction pattern in Figure 2f also confirms that these precipitates within the intervals of the 18R type LPSO laths are β′ phase [20].

3.2. Mechanical Properties of Mg-11Y-Al Alloys in T4 and T6 States

Figure 3 illustrates the tensile curves of T4 specimens at RT and T6 specimens at different temperatures.

Table 1. The strength and elongation of T4 and T6 state Mg-11Y-1Al samples obtained from Figure 3.

	Yield Strength/MPa	Tensile Strength/MPa	Elongation/%
T4-RT	151	225	4.1
T6-RT	198	275	3.9
T6-200 °C	172	262	4.1
T6-250 °C	152	256	11.5

provides the data of yield strength, tensile strength, and elongation corresponding to the tensile curves in Figure 3. Specifically, the T4 specimens possess a yield strength of 151 MPa and a tensile strength of 225 MPa at room temperature, with an elongation of merely around 4%. When the precipitate is introduced by aging treatment, the yield and tensile strength of the T6 samples are enhanced while the elongation has no obvious change. When the temperature is raised to 200 °C, the strength of T6 samples is slightly reduced while the elongation is still only about 4%. However, the elongation is significantly increased to 12% and the strength is reduced to 225 MPa when the tensile temperature of T6 samples is raised to 250 °C.

3.3. Fracture Morphology of the T4 and T6 State Mg-11Y-Al Alloy

With the increase in temperature, the elongation of the T6 sample increases significantly at 250 °C and similar traces should be seen from the fracture behavior as confirmation. Thus, we observed the tensile fracture in different states of Mg-11Y-1Al alloys to find some clues, with the corresponding SEM fractographic images presented in Figure 4. Figure 4a shows the fracture morphology of the T4 sample deformed at RT and Figure 4b–d show the fracture morphology of T6 samples deformed at RT, 200 °C, and 250 °C, respectively. The Al₂Y particles, which still maintained their original morphology, could be observed at all fractures, and the characteristics of cleavage fracture are shown between these particles and the matrix. The fracture surface with such planar features is regarded as the characteristic of cleavage fracture and is generally considered to be a brittle fracture [23,24]. From the observation of fracture morphology, Al₂Y particles seem to cause cracking. The hard Al₂Y phase formed at the grain boundaries of Mg-11Y-1Al alloys would cause stress between grains to concentrate more easily, thus becoming the origin of cracks [18].

In addition, there is an obvious disintegration plane accompanied by tearing edges (as shown in Figure 4a–c), which are brittle fractures and indicate their poor ductility. However, the fracture morphology shows tearing edges and several dimples, and the Al₂Y particles can still be observed in Figure 4d. Correspondingly, the fracture morphology of the T4 samples increased significantly when the temperature was above 200 °C, and that of the T6 samples increased significantly at 250 °C. In general, the fracture morphology of the samples corresponded well with the elongation.

3.4. TEM Observation of the Mg-11Y-Al Alloy in T4 and T6 States and Discussion

Figure 5 shows the TEM observation results of the T6 state Mg-11Y-1Al samples after tensile testing under different dual-beam conditions at room temperature. Among them, Figure 4a,c show the observations of dislocations near the same precipitate particle using different diffraction vectors, and Figure 5b,d show the corresponding enlarged views. According to the diffraction contrast vector criterion, all the dislocations observed in Figure 5b with $\overrightarrow{\mathit{g}}=(\mathbf{10}\overline{\mathbf{1}}\mathbf{0})$ are basal <a> dislocations. It can be seen that there are several dislocations piled up on both sides of the precipitate particle, without any obvious signs of passing through, and the morphology of the precipitate particle is unbroken. These characteristics indicate that the aging precipitates block the basal slip. There is no obvious dislocation contrast with $\overrightarrow{\mathit{g}}=(\mathbf{0002})$ in Figure 5b,d, which means that the non-basal slip has not been activated.

The above results demonstrate that the dominant slip mode of the T6 state Mg-11Y-1Al alloy is the basal slip at RT. Since the impeding ability of the aging precipitates to the basal slip is very strong, the yield strength of Mg-11Y-1Al alloys is high. However, it is difficult for the basal <a> dislocations to pass through the influence area of the precipitates, and the non-basal slip has not been activated; so it is difficult for the interior of the grains to contribute a large amount of deformation. At this time, the hard second-phase Al₂Y particles on the grain boundaries are difficult to deform in coordination due to the large difference in modulus from the matrix, which would lead to rapid fracture and thus result in an obvious cleavage plane. Therefore, the T6 state Mg-11Y-1Al alloy exhibits low plasticity.

Figure 6 also shows the TEM observation results of the T6 state Mg-11Y-1Al samples after tensile testing at 200 °C under different dual-beam conditions. Figure 6a,d represent the observations obtained using $\overrightarrow{\mathit{g}}=(\mathbf{10}\overline{\mathbf{1}}\mathbf{0})$ and $\overrightarrow{\mathit{g}}=(\mathbf{0002})$, respectively. Corresponding magnified observations were made on two of the precipitate particles. It can be seen from Figure 6b that a large number of dislocations are piled up on the right side of this precipitate particle, and there is a relatively obvious contrast, indicating that there is obvious deformation caused by the basal slip on the right side. In addition, Figure 6e shows that although there are some dislocations on the right side, there are only a few dislocations on the left side. Judging from the dislocation traces, these few dislocations may have passed through the precipitate particle through cross-slip [25]. After bypassing the precipitate particle, the slip traces of this non-basal dislocation seem to return to the basal plane. A similar phenomenon also appears in Figure 6c,f.

Figure 6 indicates that it is still difficult for the basal <a> dislocation to cut through the β′ phase at 200 °C. At this time, due to the increase in temperature, the CRSS of the non-basal slip is reduced; thus, it can be activated [26]. If the LPSO phase is absent in the microstructure, the magnesium will deform through the non-basal slip of dislocations to bypass the precipitate and the elongation will increase more than at RT [27]. However, there are a large number of LPSO phases in the microstructure. These LPSO laths on the basal plane have an obvious hindering effect on the non-basal plane slip, making it very difficult for the movement of dislocations [28]. The interior of the grains also cannot contribute a significant amount of deformation. Therefore, the yield and tensile strengths of the T6 state Mg-11Y-1Al sample at 200 °C are slightly reduced compared to those at RT, and the elongation hardly increases. The fracture morphology is similar to that at room temperature. A large number of Al₂Y particles and many cleavage planes were still observed.

Figure 7 shows the TEM observation results of the T6 state Mg-11Y-1Al samples after tensile testing at 250 °C using $\overrightarrow{\mathit{g}}=(\mathbf{10}\overline{\mathbf{1}}\mathbf{0})$ and $\overrightarrow{\mathit{g}}=(\mathbf{0002})$. Both the basal slip and the non-basal slip are fully activated in this sample, and the dislocation density in Figure 7a,c is quite high. In addition, it can be seen from the red circles in Figure 7a,c that this precipitate particle has been cut through by dislocations. The corresponding magnified Figure 7b shows that there are obvious dislocation lines and contrasts on both the upper and lower sides of this precipitate particle, indicating that there are obvious deformations on both sides. As can be seen from Figure 7d, a large number of <c + a> dislocations directly cut through the precipitated phase particles without obvious cross-slip characteristics.

This figure provides evidence of the extensive activation of dislocations. In the T6-treated Mg-11Y-1Al sample at 250 °C, due to the softening of the precipitate, the basal <a> dislocation can directly cut the precipitate particle. Although the non-basal plane slip can also be activated, the strengthening effect of the LPSO phase is not obvious since the basal slip could be activated. Therefore, both the yield and tensile strength of the alloy decreased. Correspondingly, because the deformation becomes easier, the problem of stress concentration caused by Al₂Y at the grain boundaries is significantly alleviated, and thus, the alloy shows an obvious improvement in plasticity.

Figure 7e is a schematic diagram that summarizes the interaction between dislocations and the LPSO–precipitate structure at different temperatures. Specifically, at room temperature, the basal slip of dislocations will be blocked by the precipitated phase, while the non-basal slip cannot be activated yet; at 200 °C, the dislocations will first be blocked by the precipitate, then will attempt to bypass the precipitate for non-basal slip through cross-slip, and will finally be blocked by the LPSO phase; and at 250 °C, the dislocations will directly cut through the precipitate particles through the basal slip. The above dislocation behaviors will directly affect the deformation in the grains, thereby influencing the strength and plasticity of the alloy.

4. Conclusions

The interaction between dislocations and the LPSO–precipitate structure in the T6 state Mg-11Y-1Al alloy at different temperatures was studied in this paper. The conclusions are as follows:

(1) The strengthening effect of the LPSO–precipitate structure is approximately the same as that of the aging precipitate at room temperature. Both enhance the strength of the alloy by blocking the movement of the basal slip.

(2) The LPSO–precipitated phase structure has a stronger effect of blocking dislocations at 200 °C than the individual aging precipitate or LPSO phase. TEM observations show that it can still effectively block the movement of dislocations at 200 °C.

(3) When the temperature rises to 250 °C, due to the softening of the aged precipitated phase, dislocations can directly cut through the precipitated phase particles. Therefore, the tensile elongation of the T6 state Mg-11Y-1Al alloy is improved. When the temperature rises from 200 °C to 250 °C, its elongation increases from 4% to 12%. Meanwhile, the tensile strength of the alloy decreases slightly.