Microstructure Evolution During Preparation of Semi-Solid Billet for 7075 Aluminum Alloy by EASSIT Process

Abstract

The 7075 aluminum alloy semi-solid billet is prepared using the extrusion alloy semi-solid isothermal treatment (EASSIT) process. These findings indicate that as the isothermal time increases, there is a noticeable increase in both the average grain size (AGS) and shape factor (SF). The relationship between the AGS, SF, and isothermal temperature is complex due to the influence of grain refinement mechanisms. The HV0.2 of isothermal samples decreased with the increase in isothermal temperature, which may be related to the increase in liquid-phase composition and AGS; Cu and Si show obvious segregation at grain boundaries and within intracrystalline droplets. The segregation of Cu and Si in the initially melted solid grains leads to the creation of intracrystalline droplets. The diffraction peaks of Al₇Cu₂Fe, Al₆(Cu, Fe), Al₂CuMg, and MgZn₂ gradually decrease as the isothermal temperature increases. Due to the influence of the grain refinement mechanism and melting mechanism, the coarsening behavior of grains at high isothermal temperatures is more complicated, and the coarsening rate constant shows an increment followed by a subsequent decrease as the isothermal temperature rises. The coarsening kinetics of 7075 aluminum alloy in a semi-solid state can be described using the LSW equation of n = 3.

1. Introduction

Widely utilized in automobiles, ships, aviation, aerospace, and various other sectors, 7075 aluminum alloy is a high-quality deformation aluminum alloy [1,2,3,4]. Conventional processing techniques of deformed aluminum alloy are often fraught with cumbersome processes and waste of materials [5]. Consequently, there is a pressing need for a novel aluminum alloy processing approach to address these drawbacks. The advent of semi-solid processing offers a fresh perspective for aluminum alloy processing.

Semi-solid forming technology is one of the significant technologies in the field of metallurgical materials. Thixoforming, a type of semi-solid forming technology, eliminates the need to handle molten metal and has minimal equipment requirements compared to rheological forming. The thixoforming process typically involves the five following steps: blank preparation, cutting, remelting, heating, feeding, and forming [6]. Crucial to thixoforming is the preparation of semi-solid billets with a uniform, finely near-spherical structure. Currently, the principal methods for preparing semi-solid billets include strain-induced melting activation (SIMA) and recrystallization remelting (RAP) [7,8,9,10,11]. The SIMA and RAP processes both involve pre-deformation and semi-solid isothermal heat treatment stages, differing mainly in the temperature at which pre-deformation is carried out. SIMA typically occurs at temperatures exceeding the recrystallization temperature, while RAP is conducted at temperatures below this point.

In recent years, many different pre-deformation methods have been used in the SIMA and RAP processes. BINESH et al. [7] were the first to utilize the repeated upsetting extrusion method (RUE) in the semi-solid processing of 7075 aluminum alloy to generate a semi-solid slurry with a fine near-spherical microstructure. They investigated the impact of the RUE cycle number on the semi-solid microstructure evolution. BOLOURI et al. [8] explored the influence of the compression ratio on the microstructure evolution of semi-solid 7075 aluminum alloy prepared using the SIMA process. Their findings indicated that the AGS decreased as the compression ratio increased. However, the decreasing trend became less noticeable when the compression ratio exceeded 30%. The authors also discussed and verified the grain growth mechanism and microstructure coarsening of 7075 aluminum alloy. JIANG et al. [9] produced 7075 aluminum alloy semi-solid billets using the RAP and SIMA processes and assessed the effects of isothermal temperature and holding time on the microstructure. Their research revealed that raising the isothermal temperature shortened the spheroidizing time and that the optimal holding time for semi-solid billets prepared by the RAP and SIMA methods was 20 min. However, the complex pre-deformation process often leads to decreased efficiency and application restrictions. Considering that the deformed aluminum alloy still stores a certain degree of strain energy, the method of obtaining semi-solid billets by direct isothermal treatment of these deformed aluminum alloys has gradually attracted the interest of researchers. Wang [12] et al. directly processed cold-rolled ZL104 aluminum alloy by semi-solid isothermal treatment to obtain semi-solid billets. The optimum temperature and time for preparing semi-solid ZL104 aluminum alloy were 570C and 5 min, respectively. The AGS, SF, and hardness of the obtained alloys were 35.88 um, 0.81, and 55.24 MPa, respectively. Liu et al. [13] prepared a semi-solid square billet of 2A14 aluminum alloy by direct semi-solid isothermal treatment of deformed aluminum alloy with a 2A14 hot-rolled plate as raw material and studied the effects of the holding temperature and holding time on AGS and SF, as well as the three-dimensional microstructure evolution in different directions. Since most commercial 7075 aluminum alloys are supplied in a deformed state, extruded 7075 aluminum alloy can undergo direct semi-solid isothermal treatment to produce semi-solid billets.

This study aims to explore the 7075 aluminum alloy’s microstructure evolution mechanism during the EASSIT process. The focus is on analyzing the impact of the isothermal temperature and isothermal time on the semi-solid 7075 aluminum alloy’s microstructure, with the goal of determining the suitable EASSIT process parameters. The initial findings from this research are crucial for producing 7075 aluminum alloy semi-solid billets suitable for industrial production.

2. Materials and Methods

The experiment bar material was produced through a hot-extrusion process in this study. It has an extrusion ratio of 8, a length of 1000 mm, and a bar diameter of 10 mm.

Table 1. The proportion of elements in the experimental material (wt.%).

Zn	Mg	Cu	Si	Fe	Mn	Al
5.20	2.50	1.20	0.04	0.29	0.11	Bal

presents the results of the analysis of the material’s chemical composition, conducted using an X-ray fluorescence spectrometer (XRF, Bruker S8 Tiger).

First and foremost, it is crucial to establish the solid–liquid temperature of the experimental bar prior to preparing a semi-solid billet. Using a differential scanning calorimeter (DSC, STA449F3), the solid–liquid temperature of the experimental alloy was determined. The experimental sample (15 mg) underwent heating to 700 °C at the rate of 5 °C/s in an inert gas atmosphere while placed in an alumina crucible. The corresponding DSC curve can be seen in Figure 1b. In order to minimize errors resulting from the uneven distribution of alloying elements, six different sets of DSC experiments were conducted following the aforementioned steps, with the sampling locations indicated in Figure 1a. The DSC curves of samples in region 4 were used as the standard for analyzing the curve change trend, as shown in Figure 1b. The average solidus and liquidus temperatures are 519 °C (the standard deviation is 5.1 °C) and 655 °C (the standard deviation is 6.7 °C), respectively, according to the results of the experiment. The 7075 aluminum alloy has a wide semi-solid interval (about 136 °C), which is conducive to the smooth progress of semi-solid processing.

Figure 2 shows the preparation method of a semi-solid billet for the SIMA, RAP, and EASSIT processes. It can be found that the SIMA process is characterized by the recrystallization of the metal and the accumulation of strain energy through high-temperature plastic deformation (above the recrystallization temperature) and then microstructure adjustment during isothermal treatment. In contrast, the RAP process is the deformation treatment below recrystallization temperature, followed by isothermal treatment conducted in a semi-solid state [12,14]. Due to the need for preliminary deformation and relatively high processing costs associated with these processes, a semi-solid billet preparation process for extruded alloys is proposed in this study, the EASSIT process. In this process, extruded 7075 aluminum alloy without heat treatment is directly heated to a semi-solid temperature and held for a certain time to obtain a semi-solid billet with an equiaxed structure. Heat treatment can reduce the strain energy accumulated during extrusion, which has a negative effect on the microstructure evolution during subsequent isothermal treatment. Therefore, the direct semi-solid isothermal treatment of the extruded aluminum alloy can not only greatly reduce the processing process and save the processing cost but also help to prepare high-quality semi-solid billets.

First, the extrudated bar material of the original 7075 aluminum alloy without heat treatment was cut into cylindrical samples with a length of 10 mm and a diameter of 10 mm by wire cutting equipment. Subsequently, the experiment for isothermal treatment was carried out using a PID-controlled resistance furnace with rapid temperature-rising capabilities, with the temperature rise rate established at 10 °C/min. To measure the temperature accurately, a type k thermocouple was suspended in the center of the furnace. The temperature error was controlled below 3 °C during the experiment. After reaching the set temperature, it should be insulated for 20 min. After the temperature distribution in the heating furnace was uniform, the sample was quickly placed into the furnace. According to the solid and liquid temperatures of the experimental material, the isothermal temperature was set to 600 °C, 610 °C, 620 °C, and 630 °C, respectively, and the isothermal time was set to 10 min, 20 min, and 30 min, respectively. The sample was quenched with cold water immediately after insulation to preserve its semi-solid microstructure.

The experimental sample was ground and polished according to the standard procedure, and then the metallographic sample was etched with modified Keller etching solution (96.5 mL H₂O–2 mL HNO₃–1 mL HCl–0.5 mL HF). The microstructure was analyzed using both a field emission scanning electron microscope (SEM, Zeiss Sigma 300) and an optical microscope (OM, Nikon ECLIPSE LV 150N). The elemental composition was analyzed by an X-ray energy dispersive spectrometer (EDS), and the acceleration voltage was set to 15 kV. The phase composition was determined utilizing an X-ray diffractometer (XRD, Bruker D8 ADVANCE), and the data were analyzed using Jade software. For measuring the hardness, an automatic turret microhardness tester (HV-1000Z) was employed. It should be noted that the experimental results obtained by SEM, EDS, XRD, and the microhardness instruments are based on polished but unetched metallographic samples. Image Pro Plus software was used to form a closed graph according to the outline of the α-Al grain, and then the area and perimeter of the closed graph were calculated. Each analysis result was based on at least 120 α-Al solid grains, which were located in the middle region of the image. The AGS and SF values are determined by Equations (1) and (2), respectively [12,15].

$$D={\displaystyle \frac{1}{n}}{\displaystyle \sum }_{i=1}^n\sqrt{{\displaystyle \frac{4A_i}{\pi }}}$$

(1)

$$F={\displaystyle \frac{{{\displaystyle \sum }}_{i=1}^n\frac{4\pi A_i}{p_i^2}}{n}}$$

(2)

where $A_i$ is the area of the i-th primary solid grain, n is the number of the primary solid grain, and $p_i$ is the perimeter of the i-th primary solid grain.

3. Results and Discussion

3.1. Microstructure of the Original Sample

The SEM image of the initial 7075 extruded alloy is depicted in Figure 3. The microstructure of the alloy’s longitudinal section (Figure 3a) comprises α-Al grains elongated along the extrusion direction and interspersed with intermetallic precipitated particles. Similarly, the microstructure of the alloy’s cross-section (Figure 3b) also exhibits fine precipitated particles, albeit with no distinct orientation. The compositions of the precipitated particles were analyzed by EDS. The results show that the precipitated particles are mainly of the Al₂CuMg phase, Al₇Cu₂Fe phase, and Al₆(Cu, Fe) phase. Mg, Zn, and Cu are the key elements in 7075 aluminum alloy, significantly influencing the formation of precipitated particles. Due to the PSN effect [16,17], these precipitated particles will become the preferred nucleation points, which has an important influence on the microstructure evolution during semi-solid isothermal treatment.

3.2. Semi-Solid Microstructure Evolution

Figure 4 clearly displays the microstructure evolution of the extruded 7075 aluminum alloy at isothermal temperatures from 600 °C to 630 °C and isothermal times from 10 min to 30 min. The result shows that the semi-solid structure of extruded 7075 aluminum alloy is an obvious thixotropic structure, primarily characterized by near-spherical α-Al solid-phase grains, intracrystalline droplets, and intercrystalline liquid films. When the isothermal temperature was 600 °C and the isothermal time was 10 min, the grain boundary was obviously melted, with the solid grains enveloped by the surrounding liquid phase. Some solid grains have a tendency to grow, but at this time, the sample’s solid grain size was uneven, and their shapes were irregular. As the isothermal temperature and time increased, the composition of the liquid phase increased, and the intercrystalline liquid film became significantly thicker, but the number of intracrystalline droplets decreased instead, and large-sized intracrystalline droplets gradually appeared in the crystal, which is caused by the mutual aggregation of adjacent small droplets in the crystal [14]. The change in isothermal time and isothermal temperature also impacts grain morphology. The AGS expands with the duration of isothermal time. When the isothermal time increased from 10 min to 20 min, the grain growth trend was obvious, the grain spheroidization degree deepened, and coarse equiaxed grains appeared in the sample. With a further increase in the isothermal time to 30 min, the grain size changed little, but the grain spheroidization degree was further deepened. The increase in isothermal temperature has a limited impact on grain size but accelerates the spheroidization process of grains. When the alloy is extruded to produce plastic strain, the gradual rise of dislocations and defects in the lattice leads to the formation of non-equilibrium grain boundaries and the accumulation of a substantial amount of strain energy. Although a significant portion of the strain energy is released in the form of heat, part of the strain energy still exists inside the deformed alloy. During the subsequent semi-solid isothermal treatment, the remaining strain energy acts as the driving force of recrystallization, recombining the extruded grains into higher spheroidized and finer equiaxed grains [18].

Figure 5 indicates the relationship between the AGS and SF of α-Al grains and isothermal temperature and time during semi-solid isothermal treatment. The findings indicate a noticeable increase in both AGS and SF with a prolonged isothermal time at a constant isothermal temperature. At different isothermal times, the transformation laws of AGS and SF with an isothermal temperature are different. When the isothermal time is 10 min, there is an initial increase followed by a decrease in AGS and SF with a rising isothermal temperature. When the isothermal time is extended to 20 min, the AGS increases gradually with a rising isothermal temperature, while the SF experiences an initial decrease followed by an increase. When the isothermal time reaches 30 min, the AGS reaches its peak at 610 °C and subsequently decreases gradually, whereas the SF is positively correlated with isothermal temperature.

The grain evolution mechanism in semi-solid isothermal treatment is relatively complex, and currently, researchers generally believe that there are two main grain evolution mechanisms in isothermal treatment, namely coarsening and spheroidization [19,20]. The coarsening process of solid grains is mainly represented by coalescence and Ostwald ripening mechanisms [15,21]. During the isothermal treatment process, the eutectic phase with a low melting point at the grain boundary melts first to form a liquid phase. When the isothermal temperature and time are low, the liquid phase is less, and there is contact between solid grains. At this time, the adjacent grains merge through the coalescence mechanism, and the grain growth rate is faster. The study by S. Takajo et al. [22] shows that the probability of grain coalescence is related to the number of adjacent solid grains. Therefore, the coalescence mechanism is the main coarsing mechanism under high-solid-phase conditions. As the liquid-phase composition at the grain boundary increases, the grains are completely surrounded by the liquid phase, the Oswald ripening mechanism plays the main role.

The microscopic images in Figure 4 can further verify this conclusion. As shown in Figure 4, the coalesce mechanism is marked by the red arrow, and the Ostwald ripening mechanism is marked by the blue box. It can be found that the coalescence mechanism is common at high solid fractions (Figure 4b,c,e,f), while at low solid fractions, most grains are separated by the liquid phase, and the Ostwald ripening mechanism comes into play (Figure 4i,k,l). Chang et al. [10] found that the spheroidization of grains during the isothermal process was affected by the curvature effect; due to the effect of curvature, the atoms at the high-curvature grain boundaries spread to the low-curvature grain boundaries, which promotes the spheroidization of grains. The rate of spheroidization is primarily determined by the atom diffusion rate, with the diffusion rate of liquid atoms significantly surpassing that of solid atoms [23]. Therefore, the more liquid-phase composition at the grain boundary, the faster the grain spheroidization rate.

As shown in Figure 6, there are some fine particles in the microstructure of the sample after isothermal treatment, and these fine particles are generally distributed in three or even multiple-grain contact regions. The composition of these particles was analyzed by an energy dispersive spectrometer.

Table 2. EDS analysis results of isothermal samples in different regions when the isothermal temperature is 620 °C and the isothermal time is 20 min.

Element	A (at.%)	B (at.%)	C (at.%)	D (at.%)	α-Al Solid Grain (at.%)
Al	95.18	94.76	90.12	94.37	92.72
Zn	3.41	4.01	4.13	3.85	4.81
Mg	0.08	1.10	2.82	0.10	1.20
Cu	1.26	0.08	2.80	1.30	1.10
Si	0.05	0.02	0.09	0.01	0.06
Fe	0.02	0.03	0.04	0.37	0.11

shows the EDS analysis results at different locations of the sample. The results show that the element composition of these small particles is similar to that of α-Al solid grains. This phenomenon is attributed to the grain refinement mechanism during isothermal treatment. Throughout this process, the liquid phase diffuses along the grain boundary and gradually accumulates in three or more grain contact regions. The increasing liquid-phase composition in these regions enhances permeability, enabling it to penetrate the grains along irregular depressions. It combines with the internal droplets of the grains, destroys the original grain boundaries, and then splits the coarse solid grains into relatively small solid grains. This mechanism also elucidates the reason for the decrease in AGS in certain isothermal samples at higher isothermal times and temperatures. This phenomenon was also reported by JIANG et al. [14].

In order to understand the hardness variation trend of the samples at different isothermal temperatures, the Vickers hardness (HV0.2) of the samples after isothermal treatment was tested by a microhardness tester. The experiment frequency was 50 Hz, the test force was set to 0.2 kgf, and the test force holding time was set to 10 s. At least six different regions were selected for testing during the experiment, and the test point was ensured to be at the center of the α-Al solid grain so as to reduce the measurement error during the experiment as much as possible. The HV0.2 change curve of the 7075 aluminum alloy sample after the 20 min isothermal treatment at different temperatures is shown in Figure 7. The findings reveal a gradual reduction in the alloy’s HV0.2 as the isothermal temperature increases. Considering the differences of AGS in the microstructure at different isothermal temperatures (Figure 5a), we chose the Hall–Petch relationship (Equation (3)) to represent the relationship between HV0.2 and AGS [24]. It can be seen from the formula that the hardness is negatively correlated with the AGS, indicating that the smaller the AGS is, the higher the hardness will be. According to Figure 5a, with the 20 min isothermal time, the AGS gradually increases with rising isothermal temperature, which explains why the sample’s hardness decreases with rising isothermal temperature. Additionally, we can also observe from Figure 7 that the declining trend of hardness significantly intensifies when the isothermal temperature escalates from 620 °C to 630 °C. This phenomenon is primarily caused by the heightened composition of the liquid phase. According to Figure 4, with the 20 min isothermal time, the composition of the liquid phase progressively increases with the increase in isothermal temperature, and the content of the as-cast structure formed after rapid cooling is higher. In the as-cast structure, there are usually a certain number of defects, such as pores and inclusions, which will reduce the hardness of the alloy.

$${ H}_V=H_0+k_H{D}^{-1/2}$$

(3)

where $H_V$ is the Vickers hardness, $H_0$ is the intrinsic hardness, $k_H$ is the Hall–Petch coefficient, and D is the AGS.

According to the changes in the AGS and SF relative to the isothermal temperature and time, the optimal process parameters for producing semi-solid billets of extruded 7075 aluminum alloy via the EASSIT process are determined to be 620 °C for the isothermal temperature and 20 min for the isothermal time. The SF is 0.80, and the AGS is 93 μm. The disparity in microstructure between the axial and radial orientations of the initial sample is evident from Figure 3a,b. Therefore, to determine the uniform distribution of the microstructure of the sample after isothermal treatment along different directions, the axial and longitudinal microstructures following isothermal treatment were compared, as detailed in the findings presented in Figure 8. These results highlight the absence of substantial differences in the microstructure between the longitudinal and transverse directions after isothermal treatment, which meets the requirements of the subsequent semi-solid processing.

3.3. Phase Evolution and Alloying Element Distribution During Semi-Solid Isothermal Treatment

During isothermal treatment, the presence of the liquid phase results in the inevitable occurrence of element segregation within the microstructure. The effect of the segregation of most elements on the properties of the alloy is unfavorable, so it is necessary to study the distribution of elements during the semi-solid isothermal treatment. Figure 9 indicates the EDS analysis results and the SEM image of a sample with an isothermal temperature of 620 °C and an isothermal time of 20 min. The results reveal distinct segregation of Cu and Si at the grain boundaries and within intracrystalline droplets (Figure 9c,d). Conversely, as the primary elements of the 7075 aluminum alloy, Zn and Mg exhibit an even distribution within the microstructure (Figure 9f,g). This is not exactly the same as the results of BINSEH et al. [11], who found Mg loss at the grain boundaries of semi-solid 7075 aluminum alloy. During isothermal treatment, the solute element is difficult to diffuse due to the restriction of crystal position. At the grain boundary, the solute element is easier to diffuse and migrate, so the solute element will tend to accumulate around the grain boundary. The segregation of Cu and Si culminates in the formation of Al-Cu and Al-Si eutectic phases at the grain boundary, whose melting points are lower than that of the α-Al grain [25]. With the increase in temperature, the Al-Cu and Al-Si eutectic phases first melt to form an intercrystalline liquid phase and permeate between α-Al solid grains. In addition, the formation of intracrystalline droplets is also due to the segregation of Cu and Si elements in the initially melted solid grains. Zhang et al. [26] believe that the formation of intracrystalline droplets is mainly caused by the segregation of alloying elements and the decrease in solid–liquid interface energy caused by the agglomeration of solid grains with complex shapes during the isothermal process. Zn and Mg have similar crystal structures and atomic sizes to Al and are easier to form solid solutions with Al, so they are easily soluble in α-Al solid grains [16], and there is no obvious segregation of Zn and Mg elements during isothermal treatment. Except for the above elements, segregation of Fe and Mn was not found during the isothermal treatment (Figure 9e,h).

A predecessor’s research has highlighted the significant influence of the secondary phase within the alloy [27,28] on its properties, necessitating an investigation into phase evolution during semi-solid isothermal treatment. The X-ray diffraction (XRD) diagram in Figure 10 illustrates 7075 aluminum alloy samples held at varying isothermal temperatures for 20 min. From Figure 10a, it is evident that the intensity of the α-Al peak conspicuously declines with an increasing isothermal temperature, reflective of the augmented liquid-phase composition. In addition, the intensity of diffraction peaks corresponding to the Al₇Cu₂Fe, Al₆(Cu, Fe), Al₂CuMg, and MgZn₂ phases gradually diminishes as the isothermal temperature rises, as depicted in Figure 10b. Notably, the melting temperature of the MgZn₂ phase in the 7xxx series aluminum alloy ranges from 475 to 480 °C [29]. In semi-solid isothermal treatment, the MgZn₂ phase undergoes initial melting, followed by the gradual dissolution of the α-Al and Al₂CuMg phases with an increased heating time and isothermal temperature. Upon reaching 630 °C, a small amount of Al₇Cu₂Fe and Al₆(Cu, Fe) is still present in the alloy’s microstructure, owing to their high dissolution temperatures [30]. Research findings by B. Binesh et al. [11] suggest that Al₆(Cu, Fe) particles can influence the liquid film at the grain boundary during the process of grain coarsening.

3.4. Coarsening Kinetics in Semi-Solid Isothermal Treatment

The coarsening kinetics during semi-solid isothermal treatment can be described by the Lifshitz–Slyozov–Wagner (LSW) equation [9,31]:

$$D_t^n-D_0^n=Kt$$

(4)

where $D_t$ and $D_0$ are the final and initial grain sizes, t is the isothermal time, K is the coarsening rate constant, and n is the coarsening index.

The coarsening index of the coarsening behavior caused by different reasons is different. At present, it is generally believed that n = 2 indicates grain coarsening controlled by an interface reaction, while n = 3 suggests grain coarsening controlled by volume diffusion. Additionally, coarsening indices of n = 4 and n = 5 denote grain coarsening controlled by boundary diffusion and dislocation, respectively [9,32,33]. According to the Ostwald ripening mechanism, coarsening behavior during semi-solid isothermal treatment is mainly controlled by volume diffusion [12,34,35], so n is 3. The variation in the cubic average size of primary solid grains with isothermal times under different isothermal temperatures is demonstrated in Figure 11. Through linear fitting of the curve, the coarsening rate constant K and regression coefficient R² at various temperatures are obtained, as displayed in

Table 3. Coarsening rate constant and regression coefficient of linear fitting curves at different isothermal temperatures.

Temperature (°C)	K (μm3 × s−1)	R2
600	767.17	0.9981
610	803.35	0.9989
620	629.73	0.9489
630	645.98	0.8800

. The results indicate that when the isothermal temperature is below 620 °C, the regression coefficient exceeds 0.94, demonstrating a high degree of curve fitting. At an isothermal temperature of 630 °C, the regression coefficient is 0.88, reflecting a good curve fitting. Therefore, the LSW equation of n = 3 can be used to describe the microstructure coarsing kinetics of extruded 7075 aluminum alloy in a semi-solid state.

An observation from Figure 11 reveals an increase in the coarsening rate constant as the isothermal temperature ascends from 600 °C to 610 °C. Ostwald ripening is a process in which small particles dissolve and large particles grow by ingesting the mass of small particles [36]. In this process, the rate of atom diffusion plays an important role. As the isothermal temperature increases, the intergranular liquid film thickens, resulting in an expansion of alloying element diffusion paths and thereby escalating the atom diffusion rate within the liquid phase. This contributes to an increase in the coarsening rate constant.

In addition, the precipitated particles also affect the coarsening rate constant during the isothermal process. It can be seen from Figure 10b that Al₆(Cu, Fe) particles are not dissolved at a lower isothermal temperature. B. Binesh et al. [11] found that Al₆(Cu, Fe) particles precipitated at grain boundaries in the semi-solid microstructure of 7075 aluminum alloy. The inhibition effect of these particles on the migration of the liquid film at the grain boundary leads to a decrease in the coarsing rate constant. When the temperature increases from 600 °C to 610 °C, the intergranular liquid film becomes thicker, and the blocking effect of some small size Al₆(Cu, Fe) particles on grain boundary migration is weakened, which leads to an increase in the coarsing rate constant.

However, an evident decline in the coarsening rate constant is observed when the isothermal temperature reaches 620 °C and 630 °C, reflecting a more complex coarsening behavior of semi-solid grains. Although the Ostwald maturation mechanism is still the main role, the liquid phase under the condition of high-liquid-phase composition, driven by the interface energy, is more likely to penetrate into the grain along the depression of irregular grains and gradually connect with the liquid droplets inside the grains to form new grain boundaries, thus achieving the effect of grain refinement [14]. This grain refinement mechanism will hinder the grain coarsening process. In addition, an increase in the liquid phase will also lead to the melting of the bumps on the surface of the solid grain, thereby reducing the coarsening rate constant [10]. This also explains why the regression coefficient of the fitted curve gradually decreases with an increase in isothermal temperature. According to Section 3.2, the optimal isothermal temperature of the EASSIT process is 620 °C. According to Table 3, the coarsening rate constant at this time is 629.73 μm³/s.

4. Conclusions

In this study, the semi-solid billets of 7075 aluminum alloy were prepared using the EASSIT process. The impact of isothermal time and isothermal temperature on the microstructure of 7075 aluminum alloy during the EASSIT process was thoroughly examined. Based on the findings of this study, several conclusions can be derived:

(1) The EASSIT process can influence 7075 aluminum alloy to obtain an equiaxed structure. As the isothermal time increases, there is a notable increase in the AGS and SF. However, the relationship between the AGS and SF with isothermal temperature is intricate due to the influence of the grain refinement mechanism. The optimal operation parameters of the EASSIT process were determined: the isothermal temperature was 620 °C, and the isothermal time was 20 min.

(2) The HV0.2 decreases with an increasing isothermal temperature, primarily attributed to the increase in the AGS and liquid-phase composition.

(3) During isothermal treatment, there is obvious segregation of Cu and Si at the grain boundaries and intracrystalline droplets. The formation of intracrystalline droplets is caused by the segregation of Cu and Si elements in the initially melted solid grains. The diffraction peaks of Al₇Cu₂Fe, Al₆(Cu, Fe), Al₂CuMg, and MgZn₂ gradually decrease with an increase in isothermal temperature.

(4) Due to the influence of the grain refinement mechanism and melting mechanism, the coarsening behavior of grains at high isothermal temperatures is more complicated, and the coarsening rate constant increases first and then decreases with an increase in isothermal temperature. The coarsening rate constant at 620 °C is 629.73 μm³/s.