Microstructure and Mechanical Properties of the SiCp/A357 Composites Materials

Abstract

In this study, SiCp/A357 composites containing 10% SiCp by mass were successfully fabricated using the semi-solid stir casting technique. The investigation focused on the microstructure, interfacial characteristics, and mechanical properties of the SiCp/A357 composites. The experimental findings indicated uniform distribution of SiCp within the matrix, with minimal interfacial reaction at the matrix and SiCp interface. Additionally, the interfacial bonding between SiCp and the matrix was found to be highly robust. The composites achieved an ultimate tensile strength of 233.4 MPa and an elongation of 7.67% prior to heat treatment. Following T6 heat treatment, the tensile strength showed significant improvement, reaching 431.9MPa, while the elongation decreased to 5.7%. The orientation relationship between the matrix and reinforcement was investigated, and first-principles calculations were employed to analyse the interfacial bonding between the reinforcement and the matrix. The simulations confirmed effective bonding, which consistent with the experimental observations.

Keywords:
Semi-solid stir casting; Aluminium matrix composites; Mechanical properties; Microstructure; Silicon carbide particles

1. Introduction

In recent years, aluminium matrix composites (AMMC) have garnered significant attention from researchers due to their exceptional properties, including a high specific modulus, high specific strength, low thermal expansion coefficient, enhanced mechanical properties, and superior wear resistance¹,². These materials hold significant potential for applications across various sectors such as automotive, electronic packaging, aerospace, and national defence. They are also viable replacements for certain traditional materials³-⁵ such as automotive brake discs, pistons, and aircraft wings and fuselages⁶,⁷.

At present, particle reinforced aluminum matrix composites can be fabricated through various methods, including melt infiltration, powder metallurgy⁸, stir casting, spray forming⁹, and in-situ synthesis¹⁰, among others. These methods with distinct characteristics. Such as powder metallurgy, this method is that it facilitates the preparation of metal matrix composites with a high content of reinforcing phases. It also enables better control over the homogeneity of particle distribution within the matrix. Additionally, composites produced by this method exhibit more stable property specifications¹¹. The spray forming can reduce interfacial reactions between reinforcing particles and the matrix, promote grain refinement and uniform distribution of the reinforcing phase, and minimize compositional segregation¹². The in-situ reaction synthesis technique effectively reduces reinforcement agglomeration and segregation while producing a clean reinforcement-matrix interface with strong bonding. Furthermore, the in-situ formed reinforcing particles exhibit smaller sizes and more uniform matrix distribution, resulting in enhanced mechanical properties¹³.

Although these techniques are capable of producing metal matrix composites with uniform particle distribution and excellent properties, they require more complex processes and equipment compared to stir casting. Moreover, their technical limitations and high costs in producing large parts preclude their use in the production of aluminum composites on an industrial scale. In contrast, stir casting is widely adopted for large-scale manufacturing due to its simplicity, cost-effectiveness, and production scalability¹⁴.

Stir-casting leverages the vortex generated by stirring to incorporate reinforcement particles into the alloy melt, ensuring a uniform distribution thereby enhancing the mechanical properties of the matrix alloy. However, the stir-casting process is often hindered by challenges such as SiCp agglomeration, poor interfacial wettability, and interfacial reactions between the reinforcement and matrix. These issues affect the uniform distribution of SiCp within the alloy matrix, ultimately lowering the material's mechanical properties¹⁵. Therefore, how to solve these difficulties is the focus of research. As of now, there are two main methods to solve these problems. One is to pretreat SiCp, including SiC surface oxidation treatment¹⁶ and SiCp surface modification¹⁷. The second is to optimize the preparation process parameters, such as melt temperature, ultrasonic vibration, and SiCp size¹⁸.

Mazahery and Shabani¹⁹ prepared SiC nanoparticles reinforced A356 composites by stir casting and the experimental results showed that the yield strength, ultimate tensile strength and modulus of elasticity of the composites increased with the addition of the particles, but the ductility decreased. The highest yield strength and tensile strength of 145 MPa and 280 MPa were obtained when the content of SiC nanoparticles was 3.5%. Lakshmikanthan et al.²⁰ fabricated double sized SiCp reinforced A357 composites using the stir-casting technique, which significantly enhanced their mechanical properties. The tensile and yield strengths of the composites were increased by 45.30% and 76.33%, respectively, compared to the cast A357 alloy. Zhu et al.²¹ prepared SiCnp/Al 6082 aluminum matrix composites using a combination of squeeze casting and stir casting. The results showed that the SiCnp particles were uniformly distributed, significantly refining the composite microstructure. Compared with the 6082 aluminum alloy without SiCnp addition, the grain refinement reached 30.55%. Moreover, the tensile strength, yield strength and elongation increased by 8.92%, 12.16% and 12.60%, respectively. Zhang et al.²² developed homogeneous Al/SiCp composites using air-isolated stir-casting method, which effectively avoided the formation of the brittle Al₄C₃ phase. The resulting composites exhibited excellent comprehensive performance. Venkatesh et al.²³ produced aluminium matrix composites reinforced with nano-SiCp using a stir-casting technique aided by ultrasound. The nano-SiCp was evenly distributed throughout the matrix and hardness and tensile strength of the composites reached 163 BHN and 431 MPa, respectively, when the SiC content was 1.5 wt%.

In this study, a semi-solid stir casting method was employed to address the issues with the stir-casting process in producing the SiCp/A357 composites. This method achieved uniform SiCp distribution while avoiding interfacial reactions, thereby significantly enhancing the mechanical properties of the composites. Furthermore, the microstructure and mechanical properties were systematically characterized, and complementary first-principles calculations were used to investigate interfacial characteristics. Through combined simulation and experimental approaches, the mechanism by which the reinforcement strengthens the matrix mechanical properties was revealed.

2. Experimental Materials and Methods

2.1. Composites preparation

In this experiment, a commercial A357 alloy (Tianjin Lang Feng Da Metal Products Co., Ltd., China) was used to prepare SiCp/A357 composites. The SiCp particles(Hebei Chuancheng Metal Materials Co., Ltd., China) with an average size of 15μm were used as the reinforcement. The chemical composition of the A357 alloy, analysed using ICP-OES (Thermo Fisher iCAP PRO) is presented in Table 1.

Table 1
A357 aluminium alloy composition/wt.%.

Alloy	Elements
	Si	Mg	Ti	Cu	Fe	Al
A357	7.00	0.58	0.11	0.02	0.04	Bal.

The raw materials were weighed and proportioned according to different mass fractions. The matrix alloy was melted in a resistance furnace at 750°C and held for 15 min. Hexachloroethane was added to refine and degas the melt, followed by stirring and slag removal. The alloy melt was then cooled to 620°C, after which SiCp particles were added.

The molten composites were stirred continuously at 350 rpm for 25 minutes. Finally, the composite melt was poured into a preheated metal mould to form composite ingots. The process flow is illustrated in Figure 1.

Figure 1
Process flow.

2.2. Characterisation and performance testing of composite materials

Experimental specimens were prepared from composite ingots and subsequently subjected to a heat treatment in a muffle furnace. This treatment entailed heating the specimens to 540°C and maintaining this temperature for 9 hours, during which time the specimens underwent a solution treatment. Following this, the specimens were quenched in warm water at 60°C and then artificially aged at 180°C for 6 hours, after which the specimens were cooled down.

The microstructural analysis of the composites was conducted using an optical microscope (OM, Nikon ECLIPSE MA200) and a scanning electron microscope (SEM, SU3900) for both the as-cast and heat-treated specimens. Additionally, high-resolution transmission electron microscope (TEM, Tecnai G2 TF30) was employed to investigate the interfacial relationships between the matrix and the reinforcement. Room temperature tensile experiments were conducted on an electronic universal testing machine (C45.105, MTS, USA) at a tensile rate of 1 mm/min. The geometry and dimensions of the tensile specimens are shown in Figure 2.

Figure 2
Schematic diagram of tensile specimen (unit: mm).

To quantitatively evaluate the uniformity of SiCp distribution, 100× metallographic images were divided into 3×3 rectangular regions of equal size. The area fraction of SiCp in each region was measured, and the coefficient of variation (denoted as ξ) was calculated as the ratio of the standard deviation to the mean area fraction. A smaller ξ value indicates better distribution uniformity within the analyzed area. For quantitative evaluation of SiCp agglomeration, Image-Pro Plus software was used to analyze 500× metallographic images from multiple regions. The software calculated both SiCp area fraction and agglomerate area fraction per image. The agglomeration parameter β was defined as their mean ratio. A smaller β value indicates lower agglomeration, while a larger β reflects more severe particle clustering.

2.3. Materials studio calculation

The Materials Studio software was used to calculate the properties of the interfacial structure between the reinforcement and the matrix, focusing on the stability of the interface and the bond strength. These calculations were based on Density Functional Theory (DFT) using the CASTEP software package²⁴,²⁵. The electron exchange and correlation energies were calculated using the Perdew-Burke-Ernzerhof (PBE) in the generalized gradient approximation (GGA)²⁶,²⁷. The plane-wave cut-off energy was set to 380 eV, and the Brillouin zone was summed using the K-point grid proposed by Monkhorst-Pack's K-point grid. In particular, the k-points for the bulk phases of Al, SiC, and Al/SiC interfaces were 6×6×6, 5×5×2, and 10×10×1, respectively. The following convergence criteria were applied in the computational process: the energy change required to converge was set to a value of less than 1×10⁻⁵ eV/atom; a maximum stress limit of 0.05 GPa, a maximum force of 0.03 eV/Å, and a maximum displacement of 0.001 Å.

Prior to optimising the interface, the bulk phase properties of Al and SiC were calculated. The calculated values of lattice constants for the bulk phase of FCC-Al were a = b = c = 4.0476Å, which are in good agreement with the experimental data of 4.05Ų⁸. For SiC bulk phase, the calculated values of lattice constants were a = b = 3.083Å and c = 10.092Å, which are also in good agreement with the experimental values, i.e., a = b = 3.073Å and c = 10.053Ų⁹. The aforementioned calculations demonstrate that the parameters employed in our computations ensure sufficient accuracy for reliable interfacial property analysis.

3. Results and Discussions

3.1. Microstructure morphology

SiCp/A357 aluminium matrix composites were obtained using a semi-solid state stir casting technique combined with mechanical stirring. During the stirring process, the high viscosity of the semi-solid melt exerts a grinding effect, enhancing both particle dispersion and interfacial wettability. This results in a more uniform distribution of reinforcement particles within the alloy melt, which in turn leads to an improvement in the mechanical properties. Figure 3 illustrates the microstructure of SiCp/A357 composites before and after heat treatment. The metallographs of the composites before heat treatment are presented in (a)-(c), while (d)-(f) show the metallographs after heat treatment. The results demonstrate that SiCp is uniformly distributed in the Al matrix in both cases. However, at higher magnification, it can be observed that there are still SiCp agglomerates present in some regions of the Al matrix. These agglomerates contribute to stress concentration within the matrix, potentially promoting crack formation in the affected areas. This can lead to a reduction in the composite's load-bearing capacity, ultimately impairing its mechanical properties³⁰.

Figure 3
Microstructure and morphology of the composites: (a, b, c) non-heat treat; (d, e, f).

The distribution and agglomeration of SiCp were quantified using image pro plus software. The results are shown in the Table 2. As can be seen from the table, the SiCp distribution is relatively uniform, but agglomeration still exists.

Table 2
Quantitative analysis of the distribution of reinforcements.

Composites	Distribution parameter of reinforcements
	ξ	β
SiCp/A357	0.077	10.2%
SiCp/A357 heat treatment	0.089	14.9%

Figure 4 presents the line scan analysis of the composite interface, with the scan direction indicated in Figure 4a, extending from the SiCp surface to the matrix surface. Figure 4b and 4c illustrate the EDS line scan curves showing atomic concentration changes for Si and Al, where the curves overlap in the transition layer region at the bonding interface between SiCp and the matrix. This transition layer is a consequence of atomic diffusion across the interface, and its thickness can be taken as an indicator of the strength of the interfacial bonding to a certain extent. The combined analysis of SEM images and EDS line scan results demonstrates that the interface between the reinforcement particles and the matrix is well bonded, exhibiting no discernible interfacial reaction layer. This robust bonding facilitates effective load transfer from the matrix to the reinforcement particles, enhancing the load-bearing capacity of the composites and, consequently, improving its mechanical properties. Furthermore, The EDS line scan reveals that the interface is primarily composed of Al and Si elements, with no evidence of harmful interfacial products such as Al₄C₃.

Figure 4
Line scan of the composite interface.

To further examine the interfacial bonding between the Al matrix and SiCp, as well as the presence of any interfacial reactions, the samples were analysed by HRTEM. Figure 5 presents the HRTEM image of the composite interface. As can be seen from Figure 5a, a relatively clean interface is visible between SiCp and Al matrix, with no evidence of Al₄C₃ formation, indicating the absence of interfacial reaction. This observation aligns with the results of the EDS line scan.

Figure 5
HRTEM image of the composites interface.

Figure 5b shows the diffraction spot pattern at the interface, where the red colour corresponds to SiCp and yellow to the Al matrix. At the interface, the crystal plane spacing of Al(111) is measured at 0.236 nm, and that of SiC(0001) is 0.2796 nm. The mismatch between the two crystal planes is calculated to be 15.6%. This indicates that there exists a semi commensurate relationship between the two crystal planes, i.e., there exists a site-directional relationship between Al(111) and SiC(0001).

3.2. Mechanical properties

The comprehensive performance of SiCp/A357 composites can be reflected by their mechanical properties. Figure 6a illustrates the stress-strain curves for the composites. Samples 1 and 2 correspond to specimens before heat treatment, while Samples 3 and 4 represent those after heat treatment. From the figure, it can be observed that before heat treatment, the samples yielded early under increasing load and when the strain exceeds 1%, strain growth rate slows down as the stress increases. In contrast, the heat-treated samples exhibited no yielding at the initial stage under similar loading conditions. The uniform dispersion of SiCp in the matrix alloy enhances tensile strength by hindering dislocation motion. Consequently, this constraint on dislocation movement under external forces enhances the tensile strength of the composites. A more intuitive visualisation of the tensile strength and elongation of the composites can be seen in Figure 6b. Without heat treatment, the composite exhibited an ultimate tensile strength of 233.4 MPa and an elongation of 7.67%. Following heat treatment, the ultimate tensile strength increased to 431.9 MPa, while the elongation decreased to 5.7%.

Figure 6
Composite stress-strain curves showing tensile strength and elongation: (a) stress-strain curve; (b) tensile strength and elongation.

The addition of SiCp enhances tensile strength while retaining the material’s inherent malleability, improving its suitability for engineering applications. However, heat treatment increases tensile strength at the expense of reduced elongation.

Figure 7 shows the scanning electron microscope (SEM) images of the fracture morphology of the composite. The images in panels (a) and (b) display the fracture morphology of the composites ​​prior to​​ heat treatment, while those in panels (c) and (d) depict the morphology ​​after​​ heat treatment. As seen in Figure 7a and 7c, a considerable number of small, shallow tough nests are present in the composites ​​both before and after​​ heat treatment. Upon closer observation (Figure 7b and 7d), brittle fracture features (e.g., ​​cleavage planes) are also evident, with flat sections and the absence of obvious cavities. However, the brittle fracture characteristics are more pronounced in comparison to the specimens prior to heat treatment. While the heat treatment led to a notable increase in tensile strength, it also caused a reduction in elongation, highlighting the trade-off between strength and ductility.

Figure 7
Fracture morphology of material tensile specimen.

Based on the above analysis, the fracture behaviour of SiCp/A357 composites is characterized by ​​a combination of brittle and quasi-cleavage​​ modes. This results from the fragmentation of eutectic Si particles or SiCp. In the composites, dislocations tend to accumulate near SiCp due to their pinning effect, generating localized stress concentrations that lead to the cracking of SiCp. Notably, the strong interfacial bonding between the A357 matrix and SiCp causes the SiCp to fracture during tensile loading rather than debond. This highlights the critical role of interfacial bonding under tensile loading: when the bond strength exceeds that of the SiCp themselves, the load is effectively transferred to the particles, subjecting them to sufficiently high stresses until fracture occurs. Conversely, weak interfacial bonding hinders stress transfer, resulting in particle is pulled out and the formation of interfacial voids.

3.3. First principles calculations

3.3.1. Surface properties

The measurement of surface energy is an important indicator for evaluating surface stability. In the case of a two-phase interface, the preference is for low-surface-energy surfaces as this facilitates the formation of a more stable interface structure and a stronger interface bond. To ensure the reliability and accuracy of the interface structure, seven atomic layers of Al(111) and 13 atomic layers of SiC(0001) were constructed, as determined by HRTEM. The bonding strength at the interface is significantly influenced by the diverse atomic composition of the SiC surface. Accordingly, the surface energies of the Si and C-terminated surfaces were calculated separately to assess the stability of the SiC surface. The formula used for calculating surface energy is as follows³¹:

${\delta }_{surf}=\frac{E_{slab}^A-\frac{N_{slab}^A}{N_{bulk}^A}{E}_{bulk}^A}{2*A}$

where ${\delta }_{surf}$ denotes the surface energy of substance A, $E_{slab}^A$denotes the energy of the surface slab of A, $E_{bulk}^A$is the bulk energy of A, $N_{slab}^A$ denotes the number of atoms in the slab, $N_{bulk}^A$ denotes the number of atoms in the crystal structure of substance A, and A denotes the cross-sectional area of the surface.

Table 3 illustrates the surface energies of SiC with varying terminations. The surface energies of SiC truncated by Si atoms and by C atoms are 4.4532 J/m² and 7.8606 J/m², respectively. These values are considerably higher than the surface energy of Al (0.8078 J/m²), indicating that the SiC surface is less stable compared to the Al surface. Consequently, SiC tends to adsorb other atoms to reduce its surface energy and enhance stability. Additionally, the surface energy of the C-truncated SiC surface is significantly higher than that of the Si-truncated SiC surface. Consequently, the C-truncated SiC surface is more prone to forming a bond with the (111) surface of Al than the Si-truncated SiC surface.

Table 3
Surface energies of different terminal atoms of SiC.

Terminal Atom	$E_{slab}^A\left(eV\right)$	$E_{bulk}^A\left(eV\right)$	$N_{slab}^A$ (个)	$N_{bulk}^A$ (个)	${\delta }_{surf}\left(J\text{/}{m}^{\text{2}}\right)$
C	-1681.5119	-1039.5481	13	8	7.8606
Si	-1685.0139	-1039.5481	13	8	4.4532

3.3.2. Al(111) and SiC(0001) interface

The Al(111)/SiC(0001) interface model was constructed using the stacking method, where seven layers of Al(111) were positioned atop 13 layers of SiC(0001), with a 15 Å vacuum layer in-between the two to prevent their interaction. Given that the C-atom truncated SiC surface is more likely to bond with the low-index surface of Al, the C atoms were selected as the terminals. These can be classified into three types according to the stacking of interfacial atoms: TOP, FCC and HCP. Consequently, there are three interfacial models in total, as illustrated in Figure 8. The 'TOP' sequence indicates that the first layer of C atoms is directly above the first layer of Al atoms; the 'HCP' sequence indicates that the first layer of C atoms is above the second layer of Al atoms; and the 'FCC' sequence indicates that the first layer of C atoms is located above the third layer of Al atoms.

Figure 8
Schematic diagram of the Al(111)/SiC(0001) C-terminal interface model, with blue balls representing Al atoms, pink balls representing Si atoms, and yellow balls representing C atoms.

Based on the constructed interfacial structure model, the adhesion work of the interface can be calculated. This value reflects the stability and bonding strength of the material interface, providing insights into the nature of the interfacial bonding. The formula is as follows³²:

$W_{ad}=\frac{E_{slab}^{Al}+E_{slab}^{SiC}-E_{slab}^{Al/SiC}}{A_{int}}$

where $E_{slab}^{Al/SiC}$ denotes the total energy of the interface, $E_{slab}^{Al}$denotes the energy of the Al(111) surface, $E_{slab}^{SiC}$ denotes the energy of the SiC(0001) surface, and $A_{int}$ is the cross-sectional area of the interface. The results of the calculations are presented in Table 4. It is worth assuming that greater the adhesion work, the stronger the interfacial bonding. A comparison of the adhesion work of the three structures reveals that the interfacial adhesion work is the largest when the C atom of SiC(0001) is directly atop the Al atom. This suggests that the TOP stacking arrangement provides the strongest interfacial bonding, enhancing the stability and mechanical integrity of the Al/SiC composite interface.

Table 4
Adhesion work at the interface of Al(111) and SiC(0001).

Termination	Stacking sequences	$W_{ad}\left({\text{J/m}}^{\text{2}}\right)$	Interface distance d0 (Å)
C-terminated	TOP	3.88	1.98
	HCP	2.89	1.73
	FCC	2.35	1.84

3.3.3. Electronic properties

The electronic characteristics of the interface play a crucial role in understanding its bonding strength. To gain further insight into the nature of the Al(111)/SiC(0001) interface, we also investigated the interfacial electronic structure, focusing on the bonding properties. Based on the previous findings regarding interfacial adhesion work, we conducted a detailed investigation of the SiC(0001)/Al(111) interface using the C-TOP interfacial model. This analysis involves calculating key electronic structures, such as the density of states (DOS) and differential charge density, to elucidate the bonding states and electronic properties.

Figure 9 shows the differential charge density of the SiC(0001)/Al(111) C-TOP interface model. The chromatic variation is indicative of the variation in charge density across the interfaces. The blue regions indicate areas where C atoms gain electrons, while the red regions show Al and Si atoms losing electrons. This pattern highlights the pronounced electron migration at the interface, with electrons transferring predominantly from the Al side to the C atoms, indicative of robust ionic bonding characteristics. Nevertheless, the majority of the charge transfer occurs within a few layers of atoms in the vicinity of the interface. On the Al side, the first layer of Al atoms undergoes a substantial charge transfer, whereas the second layer exhibits no discernible change in charge. In contrast, the charge transfer on the SiC side is predominantly concentrated at the interface between the two layers of C and Si atoms. In the interface model, the uniform charge distribution on the Al side suggests limited charge transfer between Al atoms, indicating no strong bonding tendency among them. In contrast, the SiC side shows notable electron transfer phenomenon, with a considerable amount of charge concentrated around C atoms and a substantial charge transfer of Si atoms. This enables formation of highly robust covalent bonds between Si and C.

Figure 9
Differential charge density map of Al(111)/SiC(0001) C-TOP interface model.

On the contrary, the Al-C atoms demonstrate minimal charge transfer to the second layer, with the majority of the charge transfer concentrated on C and Si atoms in both layers of the interface. At the interface, the charge transfer between Al-C atoms is more pronounced, indicating that the covalent bond formed between Al-C atoms is considerably stronger than the Al-Si bond. It is worth concluding that the SiC(0001)/Al(111) interface formed at the C-terminated side is highly stable. Moreover, the formation of Al-C covalent bonds strengthens the interfacial bonding, thereby enhancing the overall performance of the composite.

To gain deeper insight into the bonding properties and charge distribution at the Al(111)/SiC(0001) interface, the density of states for the most stable C-TOP interface model was calculated. The resulting data are presented in Figure 10, with Figure 10a showing the total density of states of the interfacial structure and the fractional-wave density of states, while Figure 10b displays the layered density of states of the interfacial model. The red dashed line in Figure 10b represents the Fermi energy level. Figure 10a reveals that near the Fermi energy level, the total density of states is predominantly influenced by Al-p, s and C-p orbitals. The predominant type of orbital hybridization is s-p hybridization. In the valence band region of the structure (below the Fermi energy level), the densities of states are primarily contributed by the C-p, s, and Si-p, s orbitals, whereas it is predominantly contributed by the Al-p and Si-p, s orbitals in the conduction band region (above the Fermi energy level).

Figure 10
Density of states plots for Al(111)/SiC(0001) C-TOP interfacial model.

Figure 10b illustrates that the Si and C atoms situated at the interface exhibit a more pronounced pseudogap, characterised by two distinct peaks on either side of the Fermi energy level. The presence of a non-zero density of states between these peaks confirms that the gap between them is a pseudogap. The pseudogap serves as a clear indicator of the system’s covalent bonding tendency, with a wider gap signifying stronger covalent character. Notably, the C-p orbitals exhibit a larger pseudogap than the Si-p orbitals, indicating the potential for a robust covalent bond, specifically the Al-C bond, to be formed at the interface.

The bonding strength between SiCp and the Al matrix was evaluated by calculating the adhesion work of the interfacial structure. The results indicate strong interfacial bonding between the two, which aligns with previous SEM and TEM observations showing a clean, reaction-free interface. Subsequently, the electronic properties of the interfacial structure were analyzed. By examining the bonding state and electron distribution at the interface, it was further demonstrated that SiCp and the Al matrix form strong covalent Al-C bonds, enhancing interfacial bonding strength and stabilizing the interface. This improvement contributes to the alloy’s overall properties. These findings are consistent with mechanical property tests and fracture morphology analysis. Under external loading, the strong reinforcement-matrix interfacial bonding enables effective load transfer to the SiCp particles. The particles ultimately fracture when the applied stress exceeds their capacity, significantly improving the material’s performance.

4. Conclusion & Future Scope

Through experimental and first-principle computational analysis of SiCp-reinforced A357 aluminum matrix composites, we demonstrate that: (1) The semi-solid stir casting method effectively improves SiCp distribution in the matrix; (2) Reinforcing particles form strong interfacial bonds with the matrix (consistent with simulations) without interfacial reactions; (3) Heat-treated composites exhibit an ultimate tensile strength of 431.9 MPa and elongation of 5.7%.

Semi-solid stir casting provides a cost-effective, industrial-scale solution for SiCp/Al composites, simplifying manufacturing while enhancing material performance. Building on these findings, future research should investigate:​​ (1) alloying or rare-earth doping to improve SiCp-Al wettability and synergistically enhance properties;​​ (2) gradient SiCp addition to optimize distribution uniformity while studying volume fraction effects on A357 alloy microstructure and properties.

The developed SiCp/A357 composites show significant potential in automotive, aerospace, and rail applications. Their high specific strength enables weight reduction in vehicle/aircraft structural components, improving fuel efficiency and payload capacity. Enhanced wear resistance also benefits high-speed train gearboxes by extending service life and reducing maintenance frequency. Additionally, in this study, due to the limited number of samples, there may be some errors in the quantitative analysis of SiCp distribution.

5. Acknowledgements

This work was supported by Yunnan Province Major Science and Technology Special Plan (202402AB080008), Yunnan Province Major Science and Technology Special Plan (202302AB080007), Yunnan Province Science and Technology Talent and Platform Plan (202105AD160020), The Yunnan Talent Support Plan (yfgrc202406).

• Data Availability
  Data will be made available on request.

6. References

• ¹ Wang Y, Monetta T. Systematic study of preparation technology, microstructure characteristics and mechanical behaviors for SiC particle-reinforced metal matrix composites. J Mater Res Technol. 2023;25:7470-97. http://doi.org/10.1016/j.jmrt.2023.07.145
  » http://doi.org/10.1016/j.jmrt.2023.07.145
• ² Liu Y, Jia L, Wang W, Jin Z, Zhang H. Reinforcing Al matrix composites by novel intermetallic/SiC interface and transition structure. J Mater Res Technol. 2023;26:164-75. http://doi.org/10.1016/j.jmrt.2023.07.166
  » http://doi.org/10.1016/j.jmrt.2023.07.166
• ³ Yang Z, Fan J, Liu Y, Nie J, Yang Z, Kang Y. Strengthening and weakening effects of particles on strength and ductility of SiC particle reinforced Al-Cu-Mg alloys matrix composites. Materials. 2021;14(5):1219. http://doi.org/10.3390/ma14051219
  » http://doi.org/10.3390/ma14051219
• ⁴ Samal P, Vundavilli PR, Meher A, Mahapatra MM. Recent progress in aluminum metal matrix composites: a review on processing, mechanical and wear properties. J Manuf Process. 2020;59:131-52. http://doi.org/10.1016/j.jmapro.2020.09.010
  » http://doi.org/10.1016/j.jmapro.2020.09.010
• ⁵ Chen F, Wang K. Fabrication and Sc-refinement of SiCp/A356 composites by vacuum stirring casting method. J Electr Comput Eng. 2022;2022:1-11. http://doi.org/10.1155/2022/5460706
  » http://doi.org/10.1155/2022/5460706
• ⁶ Saxena A, Singh N, Kumar D, Gupta P. Effect of ceramic reinforcement on the properties of metal matrix nanocomposites. Mater Today Proc. 2017;4(4):5561-70. http://doi.org/10.1016/j.matpr.2017.06.013
  » http://doi.org/10.1016/j.matpr.2017.06.013
• ⁷ Valibeygloo N, Azari Khosroshahi R, Taherzadeh Mousavian R. Microstructural and mechanical properties of Al-4.5wt% Cu reinforced with alumina nanoparticles by stir casting method. Int J Miner Metall Mater. 2013;20(10):978-85. http://doi.org/10.1007/s12613-013-0824-2
  » http://doi.org/10.1007/s12613-013-0824-2
• ⁸ Sadhu KK, Mandal N, Sahoo RR. SiC/graphene reinforced aluminum metal matrix composites prepared by powder metallurgy: a review. J Manuf Process. 2023;91:10-43. http://doi.org/10.1016/j.jmapro.2023.02.026
  » http://doi.org/10.1016/j.jmapro.2023.02.026
• ⁹ Patil IS, A A, Rao SS, Herbert MA. Experimental investigation of tensile fractography and wear properties of Al-12.5Si alloy reinforced with ZrO₂ using spray deposition method. Mater Today Commun. 2022;30:103217. http://doi.org/10.1016/j.mtcomm.2022.103217
  » http://doi.org/10.1016/j.mtcomm.2022.103217
• ¹⁰ Girish CV, Singh H, Prasad RM. In-situ synthesized polymer-derived SiC reinforced aluminum matrix composites. J Alloys Compd. 2024;976:173270. http://doi.org/10.1016/j.jallcom.2023.173270
  » http://doi.org/10.1016/j.jallcom.2023.173270
• ¹¹ Xing Y, Li NY, Li CJ, Gao P, Guan HD, Yang CMY, et al. Effects of size and oxidation treatment for SiC particles on the microstructures and mechanical properties of SiCp/Al composites prepared by powder metallurgy. Mater Sci Eng A. 2022;851:143664. http://doi.org/10.1016/j.msea.2022.143664
  » http://doi.org/10.1016/j.msea.2022.143664
• ¹² Fan CH, He WJ, Hu ZY, Wu Q, Ni YM. Interface structure and properties of spray-forming (SiCp+β-LiAlSiO4)/6092Al matrix composites. Trans Nonferrous Met Soc China. 2023;33(4):1029-37. http://doi.org/10.1016/S1003-6326(23)66163-1
  » http://doi.org/10.1016/S1003-6326(23)66163-1
• ¹³ Ramesh CS, Pramod S, Keshavamurthy R. A study on microstructure and mechanical properties of Al 6061–TiB2 in-situ composites. Mater Sci Eng A. 2011;528(12):4125-32. http://doi.org/10.1016/j.msea.2011.02.024
  » http://doi.org/10.1016/j.msea.2011.02.024
• ¹⁴ Soltani S, Azari Khosroshahi R, Taherzadeh Mousavian R, Jiang Z-Y, Fadavi Boostani A, Brabazon D. Stir casting process for manufacture of Al–SiC composites. Rare Met. 2017;36(7):581-90. http://doi.org/10.1007/s12598-015-0565-7
  » http://doi.org/10.1007/s12598-015-0565-7
• ¹⁵ Sivashanmugam N, Harikrishna KL. Effect of SiC addition in fabrication of Al/SiC metal matrix composite by stir casting. Mater Today Proc. 2024. In press. http://doi.org/10.1016/j.matpr.2024.03.006
  » http://doi.org/10.1016/j.matpr.2024.03.006
• ¹⁶ Ureña A, Martínez EE, Rodrigo P, Gil L. Oxidation treatments for SiC particles used as reinforcement in aluminium matrix composites. Compos Sci Technol. 2004;64(12):1843-54. http://doi.org/10.1016/j.compscitech.2004.01.010
  » http://doi.org/10.1016/j.compscitech.2004.01.010
• ¹⁷ Afkham Y, Fattahalhoseini SM, Khosroshahi RA, Avani C, Mehrooz N, Mousavian RT. Incorporation of silicon carbide and alumina particles into the melt of A356 via electroless metallic coating followed by stir casting. Silicon. 2018;10(5):2353-9. http://doi.org/10.1007/s12633-018-9771-x
  » http://doi.org/10.1007/s12633-018-9771-x
• ¹⁸ Jiang W, Zhu J, Li G, Guan F, Yu Y, Fan Z. Enhanced mechanical properties of 6082 aluminum alloy via SiC addition combined with squeeze casting. J Mater Sci Technol. 2021;88:119-31. http://doi.org/10.1016/j.jmst.2021.01.077
  » http://doi.org/10.1016/j.jmst.2021.01.077
• ¹⁹ Mazahery A, Shabani MO. Characterization of cast A356 alloy reinforced with nano SiC composites. Trans Nonferrous Met Soc China. 2012;22(2):275-80. http://doi.org/10.1016/S1003-6326(11)61171-0
  » http://doi.org/10.1016/S1003-6326(11)61171-0
• ²⁰ Lakshmikanthan A, Bontha S, Krishna M, Koppad PG, Ramprabhu T. Microstructure, mechanical and wear properties of the A357 composites reinforced with dual sized SiC particles. J Alloys Compd. 2019;786:570-80. http://doi.org/10.1016/j.jallcom.2019.01.382
  » http://doi.org/10.1016/j.jallcom.2019.01.382
• ²¹ Zhu J, Jiang W, Li G, Guan F, Yu Y, Fan Z. Microstructure and mechanical properties of SiCnp/Al6082 aluminum matrix composites prepared by squeeze casting combined with stir casting. J Mater Process Technol. 2020;283:116699. http://doi.org/10.1016/j.jmatprotec.2020.116699
  » http://doi.org/10.1016/j.jmatprotec.2020.116699
• ²² Zhang WY, Du YH, Zhang P, Wang YJ. Air-isolated stir casting of homogeneous Al-SiC composite with no air entrapment and Al₄C₃ J Mater Process Technol. 2019;271:226-36. http://doi.org/10.1016/j.jmatprotec.2019.04.001
  » http://doi.org/10.1016/j.jmatprotec.2019.04.001
• ²³ Venkatesh VSS, Rao GP, Patnaik L, Gupta N, Kumar S, Saxena KK, et al. Processing and evaluation of nano SiC reinforced aluminium composite synthesized through ultrasonically assisted stir casting process. J Mater Res Technol. 2023;24:7394-408. http://doi.org/10.1016/j.jmrt.2023.05.030
  » http://doi.org/10.1016/j.jmrt.2023.05.030
• ²⁴ Li Z, Lv Q, Wei H, Zhou M, Zhang F, Shan Q. Effects of RE (Nd, Ce, Y, La) on TiC(100)/Fe(110) interface based on first-principles calculations. Vacuum. 2022;205:111449. http://doi.org/10.1016/j.vacuum.2022.111449
  » http://doi.org/10.1016/j.vacuum.2022.111449
• ²⁵ Wang Q, Shi J. First-principles analysis of SiC/Al composites interface. International Journal of Fluid Mechanics & Thermal Sciences. 2019;5(4):91. http://doi.org/10.11648/j.ijfmts.20190504.11
  » http://doi.org/10.11648/j.ijfmts.20190504.11
• ²⁶ Lu J, Guan K, Rao P, Zeng Q, Liu J, Feng Z. First-principles calculation of interfacial stability, energy, electronic properties, ideal tensile strength and fracture toughness of SiC/BN interface. Appl Phys, A Mater Sci Process. 2021;127(2):125. http://doi.org/10.1007/s00339-021-04289-z
  » http://doi.org/10.1007/s00339-021-04289-z
• ²⁷ Wang C, Chen W, Jia Y, Xie J. Calculating study on properties of Al (111)/6H-SiC (0001) interfaces. Metals. 2020;10(9):1197. http://doi.org/10.3390/met10091197
  » http://doi.org/10.3390/met10091197
• ²⁸ Liu LM, Wang SQ, Ye HQ. Adhesion and bonding of the Al/TiC interface. Surf Sci. 2004;550(1-3):46-56. http://doi.org/10.1016/j.susc.2003.12.031
  » http://doi.org/10.1016/j.susc.2003.12.031
• ²⁹ Xu XY, Wang HY, Zha M, Wang C, Yang ZZ, Jiang QC. Effects of Ti, Si, Mg and Cu additions on interfacial properties and electronic structure of Al(111)/4H-SiC(0001) interface: a first-principles study. Appl Surf Sci. 2018;437:103-9. http://doi.org/10.1016/j.apsusc.2017.12.103
  » http://doi.org/10.1016/j.apsusc.2017.12.103
• ³⁰ Wang Z, Song M, Sun C, He Y. Effects of particle size and distribution on the mechanical properties of SiC reinforced Al–Cu alloy composites. Mater Sci Eng A. 2011;528(3):1131-7. http://doi.org/10.1016/j.msea.2010.11.028
  » http://doi.org/10.1016/j.msea.2010.11.028
• ³¹ Wu Q, Xie J, Wang A, Ma D, Wang C. First-principle calculations on the structure of 6H-SiC/Al interface. Mater Res Express. 2019;6(6):065015. http://doi.org/10.1088/2053-1591/ab0be1
  » http://doi.org/10.1088/2053-1591/ab0be1
• ³² Zhang F, Li Q, Glazoff MV, Ott RT. First-principles study of interfaces in Al/SiC metal-matrix composite system. Comput Mater Sci. 2023;229:112444. http://doi.org/10.1016/j.commatsci.2023.112444
  » http://doi.org/10.1016/j.commatsci.2023.112444

• Associate Editor:
  Aloisio Klein

• Editor-in-Chief:
  Luiz Antonio Pessan

Data availability

Data will be made available on request.

Publication Dates

• Publication in this collection
  04 July 2025
• Date of issue
  2025

History

• Received
  18 Jan 2025
• Reviewed
  16 Apr 2025
• Accepted
  29 May 2025