Research Article | Open Access
Solidification Behavior and Microstructure of Al-7Si Alloys with Individual and Combined Additions of Sr and Yb
The effects of individual and combined additions of Sr and Yb on the solidification behavior and microstructure of Al-7Si alloys were investigated using thermal analysis, optical microscopy, and scanning electron microscopy. The results showed that the nucleation temperature, minimum temperature, and growth temperature of eutectic solidification were lower than those of the unmodified alloys, and eutectic recalescence at 2°C in the Yb- and (Sr + Yb)-modified alloys was observed. The decrease in the nucleation temperature and an increase in recalescence obviously refined the eutectic Si morphology of the Sr-, Yb-, and (Sr + Yb)-modified alloys from a coarse plate-like structure to a fine fibrous structure. Moreover, the three-dimensional microstructure of eutectic Si in the (Sr + Yb)-modified alloy showed the finer fibrous structure, which was mainly attributed to formation of the Al2Si2Yb phase. The formation enthalpy of Al2Si2Yb was lower than that of Al-(Si)-Sr in the (Sr + Yb)-modified alloy, thus suppressing the formation of the Al2Si2Sr intermetallic compound. The more effective Sr together with Yb further improved the modification effect of eutectic silicon.
Al-Si casting alloys have been used for a variety of purposes in the automotive and aircraft industries, due to their low density, superior castability, corrosion resistance, high strength to weight ratio, ease of recycling, high efficiency, and high wear resistance . It is well known that the morphology of eutectic Si greatly affects the mechanical properties of cast Al-Si alloys [2, 3]. Coarse primary Si and long needle-like eutectic Si affect the machinability of Al-Si alloys. These Si shapes can severely impact the matrix alloy, producing a significant concentration of stress on the Al-Si alloy, which, in turn, causes a decline in the mechanical properties of the Al-Si alloy, particularly in terms of its ductility. The literature on this subject indicates that the mechanical properties of Al-Si alloys can be improved by refining the morphology of the eutectic Si.
Modifying the Si phase in Al-Si alloys can be achieved by adding rare earth elements or other elements, such as Sr and Na [4, 5]. Recent studies have discussed the effects of rare earth elements, including La , Nd , Eu , Ce , Er , Yb , Y , Sc , and Sm  on the morphology of eutectic Si and the modification mechanisms in cast Al-Si alloys. Experimental results indicate that the morphology of eutectic Si changes from a coarse plate-like structure to a fine fibrous morphology after adding such rare earth elements, significantly improving the mechanical properties of the Al-Si alloys. The effects of the combined modifications on eutectic Si and on the mechanisms of modification have been reported [15, 16]. Liao and Sun  investigated the effect of Sr and B on the refinement of eutectic Si in near eutectic Al-11.6% Si alloys. The results showed that adding 0.03% Sr and 0.028% B changes the dendritic structure from a columnar structure to an equiaxed structure; further, the eutectic silicon fibers had a fine texture. Kumari et al.  reported that Sr, Be, Ca, and Mn modify the eutectic Si phase from an acicular form to a fibrous structure when added to Al-7 Si-0.3 Mg-0.8 Fe alloys. Furthermore, the platelet β-phase was refined. These modifications resulted in a considerable improvement in impact strength. Li and Yan  found that adding La + Yb effectively improves the mechanical properties of Al-10%Si-3%Cu alloys, significantly improving the ultimate tensile strength, hardness, and elongation and increasing the elasticity modulus. Adding mixed rare earth materials (0.6 wt.% La + Yb) improves the fluidity of ADC12 alloys by up to 1,321 mm. However, investigations on the effects of modifications and of the mechanism for adding a combination of elements are very limited in Al-Si casting alloys.
Sr and Yb elements change the morphology of eutectic Si from a coarse plate-like form to a fine fibrous structure when used to modify Al-Si foundry alloys [20, 21]. Li et al.  directly observed the adsorption of Sr atoms along the <112>Si growth direction of Si and/or at the intersection of multiple Si twins. Conversely, Yb displays a different mechanism in terms of adsorption of Sr atoms. No significant Yb-rich cluster is observed at the intersection of Si twins. Therefore, in the present study, we investigated the solidification behavior and effects of modification effects of the Al-7 wt.%Si alloy with individual and combined additions of Sr and Yb.
The material used in the study was Al-7Si alloy, which was prepared using pure aluminum (99.7%) and Al-24 wt.% Si. The graphite crucible was preheated to 400°C. Pure aluminum ingots and Al-24 wt.% Si master alloys were completely melted in the graphite crucible in a resistance furnace, with the temperature was adjusted to approximately 730°C for 10 min. After applying a thermal shield and degassing (high-purity nitrogen gas emitted through an immersed graphite rod was dipped into the melt for 10 min to ensure low hydrogen content), the material underwent drossing. Subsequently, the Al-10 wt.% Sr and Al-10 wt.% Yb modifiers were added to the graphite crucible. The system was allowed to cool at room temperature for 10 min. A thermal analysis was performed using preheated K-type thermocouples inserted into the center of the melts, with the tip placed 10 mm from the bottom of the crucible to record the thermal kinetics of the melts during solidification. The cooling curves for the melts were recorded during the cooling process. The actual composition and designation of the alloys are shown in Table 1 as analyzed by inductively coupled plasma-atomic emission spectroscopy (ICP–AES).
Cooling and derivative curves were plotted for each set of thermal data. These curves were used to measure the three temperatures of the eutectic reaction. Figure 1(a) shows an example of the measured cooling curve and the first derivative curve. Figure 1(b) is a magnification of Figure 1(a) to highlight the eutectic reaction. The nucleation temperature (TN) was defined as the temperature corresponding to the first noticeable change on the derivative of the cooling curve, where nucleation behavior noticeably changed heat flow; the minimum temperature (TMin) was the minimum temperature prior to recalescence; and the growth temperature (TG) was the maximum reaction temperature reached after recalescence, as shown in Figure 1. We then drew a straight line from the right coordinate origin, and TG and TMin are marked in Figure 1(b). The difference between TG and TMin (i.e., ∆TR = TG − TMin) was defined as recalescence of the eutectic arrest (∆TR). Note that ∆TR is sometimes denoted as undercooling.
The samples used for the microstructural observations were also used for the metallographic observations. These samples were made using traditional mechanical devices, and they were etched with a solution of 0.5 vol.% HF. Optical microscopy was conducted using a Nikon M300 microscope (Tokyo, Japan). After observing the microstructure, these samples were deeply etched with a solution (15 vol.% HCl) to reveal the three-dimensional morphology of the eutectic Si phase. A scanning electron microscope (SEM: conducted with VEGA3 TESCAN) was used to observe the microstructure of the eutectic Si after adding different rare earth elements.
The formation enthalpy of the Al2Si2Yb phase was obtained via the first-principles method. The Perdew–Berke–Ernerhof genre of the generalized gradient approximation was introduced to portray the exchange correlation potential energy. The cutoff energy for the plane wave was set to 470 eV, and a 5 × 5 × 3 grid was utilized to sample the Brillouin zone with the Monkhorst–Pack scheme. The Broyden–Fletcher–Goldfarb Shanno minimization scheme was utilized for geometric optimization, and the elastic constants were determined by the stress-strain method .
3.1. Solidification Behavior of the Alloys Modified by Different Rare Earth Element Combinations
Figure 2 shows a cluster of the cooling curves with the corresponding optical microstructures taken from Al-7 wt.% Si alloys after adding different rare earth elements. The TN, TMin, and TG values of the unmodified alloy were 571, 568, and 568°C, respectively (Table 2). When the alloys were modified with different rare earth elements, the TN, TMin, and TG values of eutectic solidification decreased. The Sr-modified alloy developed a stable temperature at a certain time compared with the matrix alloys. A 2°C eutectic recalescence in the Yb-modified alloy was observed, and the TMin and TG were 556 and 558°C, respectively. Moreover, the (Sr + Yb)-modified alloy revealed a 2°C increase in eutectic TN compared with the Yb-only modification detected by thermal analysis and shown in Figure 2. The increase in TN means that noticeable nucleation occurred at a higher temperature. This was also evident from the microstructures (Figures 3(c) and 3(d)). The spacing of the fine eutectic Si fibers in the honeycomb structure of the (Sr + Yb)-modified alloy was smaller than that of the Yb-modified alloy.
The corresponding morphologies of the unmodified alloy and the Sr-, Yb-, and (Sr + Yb)-modified alloys are illustrated in Figure 2. The eutectic Si phase in the unmodified alloy was coarse and flake-like. The modified alloys were refinable by a single or complex modification and revealed a fine lath-shaped or fibrous eutectic Si. However, the degree of refinement in the eutectic Si was not clearly visible with an optical microscope. Therefore, the deep etching method was used to reveal the effects of modifying the eutectic Si in the alloys, and the corresponding 3D SEM microstructures are shown in Figure 3. A typical coarse plate-like eutectic Si structure was discovered in the unmodified alloy, as shown in Figure 3(a). The eutectic Si structure in the Yb-modified and Sr-modified, and (Sr + Yb)-modified alloys changed in morphology in which the coarse plate-like eutectic Si was changed to a honeycomb structure composed of fine fibrous eutectic Si, as shown in Figures 3(b)–3(d). The spacing of the fine eutectic Si fibers in the honeycomb structure was related to the modifiers and was approximately 400 nm when the Sr + Yb combined modifier was added to the alloy, as shown in the upper-right corner of Figure 3(d).
3.2. Solidification Behavior of the (Sr + Yb)-Modified Alloy with Different Holding Time
The (Sr + Yb)-modified alloy exhibited the optimal modification effect, as shown in Figures 2 and 3. The melts were placed at 730°C for 180 min after adding Sr and Yb to further study the possible interactions between Sr and Yb. The cooling curves of the (Sr + Yb)-modified alloy with different holding time are shown in Figure 4. The characteristic temperatures during solidification and recalescence are shown in Figure 5 and Table 3.
It is clear from Figures 4 and 5 that the characteristic temperature decreased as holding time increased at the beginning of the experiment. The characteristic temperature was the lowest when holding time was 15 min. The TN, TMin, and TG values were 543, 541, and 541°C, respectively. After being held for 15 min, the characteristic temperature increased as holding time increased. In the end, all characteristic temperatures stabilized with increased holding time. Micrographs of the (Sr + Yb)-modified alloy at different periods of time are shown in Figure 6. The microstructure of eutectic Si was refined when holding time was 5 min, but there was still some coarse plate-like eutectic Si, as shown in Figure 6(a). Furthermore, the fine fibrous eutectic Si structure was observed after heating for 15 min, as shown in Figure 6(b). However, the fine fibrous eutectic structure worsened when the holding time was 60 min, as shown in Figure 6(c). The modification effect was completely lost after 120 min; instead, a laminar structure was observed, as shown in Figure 6(d).
In the present study, a kind of Sr-rich intermetallic compound with a particle size of 10 µm was observed in the Sr-modified alloy, as shown in Figure 7. The main components of the Sr-rich phase were Al, Si, and Sr. Some intermetallic compounds were observed in the (Sr + Yb)-modified alloy, and the results are shown in Figure 8. The mapped scanning images showed that the bright and grey intermetallic compounds were the Al-Si-Yb and Al-Si-Fe phases, respectively. However, the Sr element and Sr-rich intermetallic compounds in the combined modified alloy were not observed. Particles similar to Al-Si-Sr and the Al-Si-Yb intermetallic compound in Al-Si alloys have been identified as the Al2Si2Sr and Al2Si2Yb phase by transmission electron microscopy [21, 22].
The formation enthalpies of the Al-(Si)-RE intermetallic compounds were referenced to further analyze the probability of forming Al-(Si)-RE (RE = Sr, Yb), and the results are listed in Table 4. Due to the lack of thermodynamic parameters for the Al2Si2Yb intermetallic compound, the formation enthalpy of the Al2Si2Yb intermetallic compound was characterized by first-principle calculation, which is widely used when computing the structural and thermodynamic properties of intermetallic compounds, such as Al2Sr and Al2Si2Sr [24–26]. Table 4 shows that the formation enthalpy of the Al2Si2Sr phase was lower than those of the Al-Sr phases, suggesting that the Al2Si2Sr phase is easier to form from Al-Si alloys after adding Sr. Similar results were found for the Sr-modified alloy, as shown in Figure 7. The formation enthalpy of the Al-(Si)-Yb phases in the Al-Si alloys with the (Sr + Yb) combination was less than that of the Al2Si2Sr phase. The Al2Si2Yb phase had the lowest formation enthalpy value among the Al-(Si)-Yb phases (−34.33 kJ/mol), which caused the Al2Si2Yb intermetallic compound to form first in the Al-Si alloy modified with added (Sr + Yb). These results are consistent with the mapped scanning results for the (Sr + Yb)-modified alloy. However, the Sr element and Sr-rich intermetallic compounds were not observed in the (Sr + Yb)-modified alloy, as shown in Figure 8.
4.1. Thermal Analysis
As shown in Figure 2 and Table 2, the thermal temperatures of the alloys during the cooling process showed that the individual and combination of Sr and Yb lead to reductions in TN, TMin, and TG by 10–13°C; a 2°C recalescence was exhibited in the cooling curves of the Yb-modified alloys. Shabestari et al.  reported that the decrease in TN is related to the absorption of impure elements at the interface of Si and liquid aluminum. These impure elements would block growth, leading to a reduced TN of the alloy at an equivalent growth rate. Djurdjevic et al.  discovered that the decline in eutectic TG in the modified Al-Si alloys is related to a modification of the eutectic Si morphology. The greater the differences between the unmodified and modified Al-Si eutectic TGs, the greater the level of eutectic Si modification. Moreover, recalescence was also identified as another important factor for modifying eutectic Si morphology. An increase in recalescence under cooling is usually observed in a full modification alloy with a fine fibrous eutectic structure. Knuutinen et al.  reported that recalescence is an indication of nucleation effects, while a reduced TG indicates an effect on growth of the Al-Si eutectic in the solidification front. In this study, the Sr-, Yb-, and (Sr + Yb)-modified alloys with full modifications have lower TN and TG, as shown in Figure 2. Moreover, the cooling curves of Yb-modified alloy exhibited 2°C recalescence. The lower nucleation temperature and increase in recalescence obviously affected the changes in eutectic Si morphology. The 3D eutectic Si microstructures of the Yb-modified alloy had a finer fibrous eutectic Si, as shown in Figure 3.
When holding time was 15 min, the characteristic temperature was the lowest, as shown in Figure 5. Similar results were observed in the hypoeutectic Al-Si alloys after adding Na and Sr + Na . The reason may be adequate diffusion of modifiers into the melts when the holding time was increased from 5 to 15 min, causing the characteristic temperature to decrease as holding time increased. Then, the characteristic temperature increased as holding time increased. After being held for 90 min, recalescence disappeared during the cooling process. The differences between the unmodified and modified Al-Si eutectic growth temperatures were reduced, causing the eutectic Si modification of the alloys to worsen, as shown in Figure 6(d).
4.2. Modification Mechanism
The Sr and Yb elements significantly modified the eutectic Si in the Al-Si alloys, as shown in Figure 2. However, the mechanisms for the modifying the eutectic Si differ [30, 31]. When the Al-Si alloy was modified with Sr, the Sr elements in the alloy adsorbed on the growth steps of the eutectic Si and hindered it, causing the eutectic Si to grow into flakes via a step growth mechanism . According to the impurity induced twinning mechanism, twinning is most easily induced when the atomic radius of the modifier elements relative to that of the Si is 1.646 . The modified atom is easily adsorbed on the surface of the Si atom, effectively inhibiting growth of the Si phase . The ratio of the radius of the Yb element to the radius of the Si element was 1.65, suggesting that the Yb element will have a significant modifying effect on Al-Si alloys.
When Sr content in the alloy reaches a particular level, the Sr-rich phase in the Al-Sr master alloy, such as Al4Sr, will gather within the eutectic Si phase to form an Sr-rich intermetallic compound, as shown in Figure 7. Dahle et al.  claimed that the decrease in nucleation frequency is caused by the formation of the Sr containing intermetallics (likely Al2Si2Sr), which has the toxic effects on the nuclei and grain refinement, causing the modifying effect of eutectic Si to decrease. In the (Sr + Yb)-modified alloy, the AlxSr, AlxYb compounds in the master alloys and the Si elements preferentially reacted to form the Al2Si2Yb intermetallic compound, as the formation enthalpy of Al2Si2Yb is lower than that of Al-(Si)-Sr, as shown in Table 4. In addition, Al2Si2Sr was not observed in the (Sr + Yb)-modified alloy due to the low actual content of Sr. Therefore, the more effective Sr element was dispersed uniformly in the alloys, leading to a finer fibrous eutectic structure in the (Sr + Yb)-modified alloy, as shown in Figure 3(d).
The following conclusions were made from the present investigation:(1)Adding Sr, Yb, and Sr + Yb reduced the nucleation temperature, minimum temperature, and growth temperature of the Al-7 Si alloys. The cooling curves of the Yb-modified alloy exhibited 2°C recalescence.(2)The morphology of eutectic Si in the Sr-, Yb-, and (Sr + Yb)-modified alloys changed from a plate-like form to a fine fibrous structure. The 3D microstructure of eutectic Si in the (Sr + Yb)-modified alloy was the finest fibrous structure.(3)The formation enthalpy of Al2Si2Yb was lower than that of Al-(Si)-Sr in the (Sr + Yb)-modified alloy, thus suppressing the formation of the Al2Si2Sr intermetallic compound. The more effective Sr element uniformly dispersed in the alloys, leading to the fine fibrous eutectic structure in the (Sr + Yb)-modified alloy.
The data used to support the findings of this study are available from the corresponding author upon request.
Conflicts of Interest
The authors declare that there are no conflicts of interest regarding the publication of this paper.
The authors would like to acknowledge the financial support by the National Natural Science Foundation of China (51405216), Training Program Foundation for Young Scientists of Jiangxi Province (20153BCB23023), and the Natural Science Foundation of Jiangxi Province (20171BAB206005).
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