Microstructure Evolution and Mechanical Behaviors of SiCp/CNTs Hybrid Reinforced Al-Si-Cu-Mg Composite by Semi-solid Stir Casting

Abstract

In this study, Al-Si-Cu-Mg composites reinforced with Silicon carbide (SiC) particles and/or Carbon nanotubes (CNTs) were fabricated using semi-solid stirring casting technology (SSC). The aim of the study was to investigate the influence of SiC and/or CNTs on the mechanical properties, microstructural evolution, and deformation mechanisms of the Al-Si-Cu-Mg alloy. The findings indicated that the presence of SiC/CNTs in the composite had a significant effect on refining the α-Al phase and altering its grain orientation. Additionally, the results obtained from EBSD and TEM revealed that the microstructure of the Al-Si-Cu-Mg-SiC/CNTs hybrid composite (HAMC) underwent dynamic recrystallization (DRX) and static recrystallization (SRX), resulting in a fine equiaxed recrystallized structure. This process also led to the formation of high-angle grain boundaries (HAGBs) through dislocation rearrangement. The SiC particles were evenly distributed within the matrix, ensuring good interface bonding between SiC and α-Al phases. Moreover, the CNTs reacted with the matrix, resulting in the in-situ formation of the Al₄C₃ phase at the interface. The addition of SiC/CNTs significantly increased the tensile strength of HAMC at 25°C and 250°C, from 319MPa and 178MPa to 438MPa and 331MPa. Examination of the fracture surface of the hybrid composite unveiled that void formation occurred primarily at the regions surrounding the matrix-particle interface.

Keywords:
Semi-solid stirring casting; SiC; CNTs; Microstructure evolution; Mechanical properties

1. Introduction

Aluminum matrix composites (AMCs) possess several advantages, including low density, high elastic modulus, excellent shear performance, low thermal expansion coefficient, and good thermal conductivity. As a result, they find wide-ranging applications in various industries such as aerospace, automotive transportation, and electronic packaging¹-⁴_. For particulate reinforced AMCs, ceramic particles, such as SiC⁴,⁵, CNTs⁶-⁸,Al₂O₃⁹ and Ti₂B¹⁰-¹², have attracted much attention due to their high hardness and similarity density compared with the aluminum alloy matrix. However, AMCs prepared using conventional methods often exhibit inferior metallurgical and mechanical properties. Semi-solid stirring casting technology is an innovative technique that has been utilized to produce AMCs¹³-¹⁶. This technique takes advantage of the high viscosity of the semi-solid melt and employs enhanced shear stirring to achieve a uniform distribution of the reinforcing phase within the AMCs. The main advantage of stir casting is its ability to manufacture near net shape products at a low cost.

Several research works on the fabrication of various AMMCs were reported in the literature¹⁷-²⁶. Tang et al.²⁷ evaluated the influence of multiple pass numbers on the grain refinement and mechanical properties of SiC and stainless steel-reinforced 6061 Al-based hybrid composite materials. The findings indicated that, in comparison to the unreinforced 6061 alloy, the reinforced alloy exhibited a uniform microstructure, even particle distribution, and enhanced mechanical properties. Ma et al.²⁸ achieved successful preparation of an Al-CNTs nanocomposite using the powder metallurgy method. Subsequently, the composite was subjected to uniformization through FSP (Friction Stir Processing) channel, resulting in a significant improvement in mechanical properties. Amra et al.²⁹ prepared Al5083-SiC/CeO₂ hybrid composites and investigated the effects of SiC and CeO₂ particle volume fractions on microstructure, mechanical properties and corrosion performance. The results revealed that the Al5083-SiC(25%)-CeO₂(75%) composite displayed the most favorable combination of mechanical properties and corrosion resistance. Eskandari et al.³⁰ prepared Al8026-TiB₂/Al₂O₃ hybrid composites and investigated the effects of the reinforcement on microstructure, hardness, tensile, and wear properties. The results showed that compared to single particle-reinforced composites and alloys, the hybrid composites exhibited significantly improved mechanical and wear properties. Lu et al.³¹ utilized the squeeze casting method to fabricate composites reinforced with SiC and Cr particles in 7075 alloy. The incorporation of Cr particles resulted in improved mechanical and thermophysical properties of the materials. Additionally, the presence of nano-precipitates in the age-hardened alloy played a significant role in influencing the properties of the composites. However, there is a scarcity of reports on utilizing stirring casting technology to produce mono and hybrid composites using CNTs and/or microsized SiC particles.

In this study, SiC particles and CNTs were used as reinforcements in an Al-Si-Cu-Mg composite fabricated by SSC. SiC particles are widely employed as reinforcements in Al matrix composites due to its high hardness and elastic modulus³²-³⁶. CNTs are another potential reinforcing material known for their exceptional strength of up to 30 GPa and modulus of up to 1 TPa. With these unique properties, CNTs serve as an ideal reinforcement in Al matrix composites, contributing to the enhancement of their mechanical properties. This study systematically investigated the microstructure evolution, mechanical properties, and fracture behaviors of HAMCs.

2. Experimental

2.1. Preparation of CNT/Al-Cu-Mg-Si composites and process parameters

In this study, HAMCs were manufactured through the Ball Milling (BM) and SSC routes, as illustrated in Figure 1. The reinforcement material consisted of α-SiC particles with an average diameter of 10±4 μm. The SiCp particles were subjected to oxidation at 1200°C for 10 hours. Raw Al powders, approximately 10 μm in diameter, were used as the base material. CNTs with a purity of around 98% were fabricated through chemical vapor deposition, possessing an outer diameter ranging from 20~60 nm and a length of 0.5~2 μm. The Al powder and CNTs were mixed in alcohol at a mass ratio of 97:3, followed by ultrasonic dispersion in an ultrasonic cleaner. Subsequently, the mixture was milled in an attritor at a rotation rate of 300 rpm with a ball-to-powder ratio of 15:1 for 5 hours.

Figure 1
Schematic of fabrication process of HAMC-(SiCp/CNTs).

A series of Al-Cu-Mg-Si composites were produced using commercially pure Mg, Al-40Cu, and Al-40Si in a vacuum semi-solid strong shearing stirring device. The raw materials were rapidly heated to 700°C under vacuum, then apply mechanical stirring to the melt at a stirring rate of 500 rad/min and a stirring time of 30 minutes. Afterward, the stirring was halted, and the melt was cast into an ingot mold. To ensure better dispersion of the reinforcement within the composite, the prepared composite ingots were extruded into bars at 480°C with an extrusion ratio of 16:1. The extruded bars underwent solution treatment at 545°C for 3 hours, followed by quenching in water and subsequent aging at 170°C for 10 hours.

The chemical compositions of the composites were determined using Inductively Coupled Plasma-Atomic Emission Spectroscopy (ICP-AES) and are presented in Table 1.

Table 1
Main chemical compositions of the Al-Cu-Mg-Si aluminum alloy (wt. %).

Elements	Cu	Mg	Si	Fe	Al
Al-Cu-Mg-Si alloy	1.230	0.726	0.721	0.082	Bal.

2.2. Characterization

The locations of the microstructure observation and tensile sample are depicted in Figure 2. Following the preparation of the metallographic sample (10mm×10mm×10mm), it underwent coarse grinding, fine grinding, and polishing. Subsequently, the sample was etched for 10 seconds using Killer reagent. The samples were mechanically polished and anodized, and then observed by Optical microscopy (OM, Zeiss Axiovert .AL). The microstructures were also examined using scanning electron microscopy (SEM, Tescan MIRA3) and transmission electron microscopy (TEM, JEOL-JEM 2011Plus). TEM specimens were cut by electrical discharge machining, ground to a thickness of 60 µm, then punched to disks with a diameter of 3 mm and then dimpled to a minimum thickness of 20 mm, finally ion-beam thinned by a Gatan Model 691 ion milling system. X-ray powder diffraction (XRD, Bruker AXS D8ADVANCE X) patterns with Cu K_α radiation at a wavelength of 1.5406 Å was employed to determine the phase components of the composites. The grain size, boundary and texture orientation were investigated using electron back-scattered diffraction (EBSD, Nordly Nano, Oxford Instruments) on a MIRA3 SEM (TESCAN).

Figure 2
Schematic diagram of sampling of metallographic samples and tensile samples.

The mechanical properties of the samples were measured using a tensile machine (AGS-X) at an initial strain rate of 1 mm/s. Five specimens were measured for each condition, and then the values were averaged. The Brinell hardness was measured using an MHBS-3000AET hardness tester.

3. Result

3.1. Microstructure evolution of the prepared composites

Figure 3 presents the microstructures of the composites along the extrusion direction after heat treatment. In Figure 3a, the Al-Cu-Mg-Si alloy exhibited a typical recovery structure, with α-Al grains showing a significant variation in size and elongated along the extrusion direction. The aspect ratio was approximately 4:1, and the grain boundaries were not clearly discernible. Upon the introduction of CNTs, distinct black phases became visible in the microstructure, exhibiting a streamlined distribution. Figure 3c displays the microstructure of AMC-SiCp, characterized by a limited number of equiaxed grains and the presence of zigzag-shaped grain boundaries. In HAMC-(SiCp/CNTs), the α-Al grains exhibited a predominantly equiaxed morphology (Figure 3d), with distinct grain boundaries.

Figure 3
OM image along the extrusion directio. (a) Al-Cu-Mg-Si alloy; (b) AMC-CNTs; (c) AMC-SiCp and (d) HAMC.

Figure 4 illustrates the SEM images and EDS analysis of composites. It can be seen that SiCp were primarily concentrated at the grain boundaries, as shown in Figure 4b, c. Additionally, the regions with a higher distribution of reinforced phases featured smaller equiaxed grain sizes (Figure 4c). The EDS results show that there are obvious oxygen enrichment regions on the surface of SiCp.

Figure 4
Microstructure of Al-Cu-Mg-Si composites. (a) Al-Cu-Mg-Si alloy; (b) AMC-SiCp; (c) HAMC and (d) the EDS mapping of the SiCp/CNTs in (c).

Figure 5a displays the XRD patterns of Al-Si-Cu-Mg, AMC, and HAMC. The Al-Si-Cu-Mg alloy is represented by the green curve, AMC-CNTs by the blue curve, AMC-SiCp by the red curve, and HAMC by the black curve. A noticeable presence of SiC and C peaks can be observed in the XRD pattern of HAMC compared to that of the Al-Si-Cu-Mg alloy. Table 2 presents the variation in intensity of identified peaks corresponding to the Al phase in the composite. It is evident that the diffraction intensity of the α-Al phase on different crystal planes undergoes significant changes with an increasing amount of SiCp/CNTs. Specifically, the peak intensities of the (111) and (200) crystal planes experience substantial reductions, while the peak intensities of the (220) and (311) crystal planes gradually increase. This indicates that the addition of SiCp/CNTs alters the grain orientation of the α-Al phase, promotes α-Al growth on the (220) and (311) crystal planes, and hinders α-Al growth on the (111) and (200) crystal planes.

Figure 5
XRD patterns of Al-Si-Cu-Mg, AMC and HAMC.

Table 2
XRD data showing the intensity variation of the identified peaks for Al phase in the composite.

XRD Parameters	Standard	Al-Cu-Mg-Si	AMC/CNTs	AMC/SiCp	HAMC-(SiCp/CNTs)
2θ (deg.) (111)	38.474	38.486	38.502	38.522	38.558
Intensity (CPS)	100	1,039,711	296,694	4351	4471
FWHM (deg.)		0.141	0.167	0.147	0.161
2θ (deg.) (200)	44.722	44.710	44.736	44.760	44.799
Intensity (CPS)	45.5	80,285	28,001	4090	5442
FWHM(deg.)		0.171	0.176	0.182	0.185
2θ (deg.) (220)	65.099	65.048	65.166	65.141	65.178
Intensity (CPS)	23.3	153	395	3100	2508
FWHM (deg.)		-	-	0.192	0.249
2θ (deg.) (220)	78.234	-	-	78.280	78.299
Intensity (CPS)	32.0	-	-	1269	1131
FWHM (deg.)		-	-	0.202	0.333

Figure 6 depicts the axial crystal orientation distribution and grain boundary orientation difference statistics of the Al-Cu-Mg-Si alloy and HAMC after extrusion treatment. Figure 6a shows that under hot extrusion, the grains of Al-Cu-Mg-Si alloy are banded along the extrusion direction, and there are a large number of small Angle grain boundaries. After the addition of SiC and CNTs, although the grain distribution of HAMC was also roughly banded along the extrusion direction, the grain distribution was obviously disturbed by the reinforcement phase (Figure 6b). Compared with Al-Cu-Mg-Si alloy, a large number of fine equiaxed grains appeared, and the orientation distribution was more complex. Figure 6c~f shows the grain boundary distribution of Al-Cu-Mg-Si alloy and HAMC-(SiCp/CNTs) composite samples. In the figure, the black line denotes the high-angle boundary with an orientation difference angle greater than 15◦ from that of the adjacent grains, the green line denotes the low-angle boundary with an orientation difference angle of less than 15◦ from that of the adjacent grains. It can be inferred from Figure 7a and b that the addition of the SiCp/CNTs leads to a decrease in the proportion of small-angle grain boundaries from 77.23% to 52.19%, accompanied by an increase in the average grain orientation difference from 10.5° to 19.3°. Consequently, the inclusion of SiCp and CNTs influences the orientation change of the crystal during the alloy extrusion under stress, resulting in a tendency towards complexity.

Figure 6
EBSD of hot-extruded Al-Si-Cu-Mg alloy and HAMC: (a) IPF for Al-Si-Cu-Mg alloy; (b) IPF for HAMC; (c) Grain boundaries for Al-Si-Cu-Mg alloy; (d) Grain boundaries for HAMC; (e) the distribution of grain average misorientation (GAM) angles for Al-Si-Cu-Mg alloy; (f) the distribution of grain average misorientation (GAM) angles for HAMC.

Figure 7
EBSD of Al-Si-Cu-Mg alloy and HAMC in T6 heat treated state : (a) IPF for Al-Si-Cu-Mg alloy; (b) IPF for HAMC; (c) Grain boundaries for Al-Si-Cu-Mg alloy; (d) Grain boundaries for HAMC; (e) the distribution of grain average misorientation (GAM) angles for Al-Si-Cu-Mg alloy; (f) the distribution of grain average misorientation (GAM) angles for HAMC.

Figure 7 illustrates the axial crystal orientation distribution and grain boundary orientation difference statistics of the Al-Cu-Mg-Si alloy and HAMC after heat treatment. In Figure 7a, c, and e, it can be observed that, compared to before heat treatment, the grains of the Al-Cu-Mg-Si alloy still exhibit banding, but with larger sizes, and the average grain orientation difference is reduced to 6.9%. During the heat treatment, recrystallization occurs, leading to the merging and transformation of some grain boundaries into small-angle orientations. In comparison to HAMC before heat treatment, the zonal grain distribution disappears, and the overall microstructure displays fine equiaxed grains, resulting in a more complex grain orientation distribution.

Figure 8 presents the distribution diagrams and statistics of axial crystal recrystallization of the Al-Si-Cu-Mg alloy and HAMC before and after heat treatment. The addition of SiCp and CNTs in the Al-Cu-Mg-Si alloy increased the proportion of dynamically recrystallized grains and substructures. After heat treatment, the deformation structures in both the Al-Cu-Mg-Si alloy and HAMC are completely eliminated. Comparatively, the proportion of recrystallized grains in HAMC, in the heat treatment state, is significantly higher than that of the Al-Cu-Mg-Si alloy.

Figure 8
Distribution diagrams and statistics of axial crystal recrystallization of Al-Si-Cu-Mg alloy and HAMC: (a) hot-extruded Al-Si-Cu-Mg alloy; (b) hot-extruded HAMC; (c) T6 heat treated Al-Si-Cu-Mg alloy; (d) T6 heat treated HAMC.

In studies on the plastic deformation of particle-reinforced aluminum matrix composites, it has been observed that dense dislocation walls are formed due to the accumulation of dislocation defects. These dense dislocation walls act as strong obstacles to the movement of dislocations, resulting in grain rotation and an increase in grain misorientation³⁷. The difference in thermal expansion coefficients between SiCp, CNTs, and the aluminum matrix leads to the distribution of stress in the aluminum matrix near the SiCp and CNTs during solidification and hot extrusion. This increased distortion energy in the alloy results in the formation of vacancies and dislocations. Hence, under the same heat treatment conditions, the recrystallized grains of HAMC tend to nucleate preferentially near the reinforcement phase. Additionally, CNTs distributed in the grain boundary region exhibit a pinning effect on the grain boundaries³⁷, impeding the transformation of larger distorted structures into substructural tissues. In summary, SiCp and CNTs not only promote the recovery recrystallization of the alloy, but also significantly inhibit the growth and coarsening of recrystallized grains.

According to the bright-field TEM image shown in Figure 9a, the SiCp morphology included particles of 10μm in thickness with irregular polygonal plate-like shapes,and the distribution of CNTs was observed on the surface of SiCp. The high-resolution image in Figure 9b shows a typical interface between the CNTs and SiCp, which confirms the presence of a well-bonded interface.

Figure 9
TEM images of SiCp and CNTs in the Al matrix of the HAMC : (a) morphology of SiCp in the Al; (b) interface of CNTs/SiCp.

Figure 10 presents TEM images showcasing the distribution, morphology, and interaction of dislocations, reinforcement phases, and precipitation phases within the Al matrix of the HAMC. In Figure 10a, it can be observed that the precipitate phases, after T6 heat treatment, are uniformly distributed in the α-Al matrix in the form of rod and granule-shaped precipitates. The α-Al matrix surrounding the CNTs displays equal inclination fringes, indicating the presence of micro-aberrations induced by the CNTs. The high-resolution image in Figure 10b demonstrates the interface between the CNTs and Al. It can be seen that an Al₄C₃ phase is formed at the interface due to the preferential reaction between the amorphous carbon on the surface of the CNTs and the Al matrix. Previous studies have shown that the formation temperature of the Al₄C₃ phase ranges between 450°C and 950°C, and the preparation temperature of the composites in this study falls within this range, providing the thermodynamic conditions for the formation of Al₄C₃. Liu et al.²² found that although the wetting angle of the Al-C system was 135°~140° at 1190°C, the presence of Al₄C₃ at the interface decreased it to 55°. The formation of Al₄C₃ at this interface layer is beneficial for enhancing the bonding strength between α-Al and CNTs. In Figure 10c, the morphologies of θʹ and θ phases in the matrix are displayed after T6 heat treatment. High-resolution transmission electron microscopy (HRTEM) images of the precipitate phases indicate a clean and well-bonded interface between the precipitates and the α-Al matrix, as shown in Figure 10e and f. Additionally, significant dislocation tangles are observed in the HAMC-(SiCp/CNTs) composites, as shown in Figure 10d.

Figure 10
TEM images of distribution and morphology of dislocation, reinforcement phase and precipitation phase in the Al matrix of the HAMC. (a) distribution of precipitation phase;(b) interface of CNTs/Al; (c) morphologies of θʹ and θ phases; (d) morphologies of dislocation; (e,f) HRTEM images of θʹ and θ phases.

3.2 Mechanical properties of the prepared composites

Figure 11 illustrates the aging hardening curves of an Al-Si-Cu-Mg alloy, AMC , and HAMC treated at 170°C. The addition of reinforcements significantly accelerates the aging response of the Al-Si-Cu-Mg alloy and increases its hardness values both in the unaged and peak-aged states. The hardness begins to increase rapidly from 1 hour, indicating a large curvature in the aging hardening curve, indicating that the four materials are in an under-aging state. As the aging treatment continues, the Al-Si-Cu-Mg alloy reaches its peak hardness value at 12 hours but starts to exhibit a significant decrease with further aging treatment. AMC-CNTs achieves its peak hardness at 11 hours of aging treatment, with a slight decrease during continued aging and then stabilizing. AMC-SiCp reaches its peak hardness at 9 hours, with a slight decrease during continued aging as well. HAMC-(SiCp/CNTs) reaches its peak hardness at 7 hours after aging treatment. During solution quenching, the mismatch in thermal expansion coefficients between SiCp, CNTs, and Al generates residual stress at the interface between the matrix and the reinforcement, leading to the formation of numerous dislocations. The presence of residual stress and high-density dislocations can accelerate the precipitation of the precipitated phases within the matrix.

Figure 11
Aging hardening curves of Al-Si-Cu-Mg alloy、AMC and HAMC at 170°C.

Figure 12 presents the engineering stress-strain curves for both the Al-Si-Cu-Mg alloy and composite samples at 25°C and 250°C. In Figure 12a, b, the tensile strength, yield strength, and elongation of the Al-Si-Cu-Mg alloy and composite samples at 25°C are shown. It is evident that all the composite samples exhibit higher tensile strength but with a subsequent reduction in elongation compared to the Al-Si-Cu-Mg alloy. Among the composite samples, HAMC demonstrates the highest tensile strength of 438 MPa. The ultimate tensile strengths for Al-Si-Cu-Mg, AMC-CNTs, and AMC-SiCp are recorded as 319 MPa, 386 MPa, and 416 MPa, respectively. Figure 12c, d illustrate the tensile strength, yield strength, and elongation of the Al-Si-Cu-Mg alloy and composite samples at 250°C. As the temperature rises, the elastic modulus of the Al-Si-Cu-Mg alloy decreases, elongation increases, and tensile strength decreases to 176 MPa. The tensile properties of AMC-SiCp and HAMC composites at high temperatures remain consistent with those observed at room temperature. In the early stage of the tensile process, elastic deformation occurs, and no evident plastic deformation stage is observed, indicating brittle fracture characteristics.

Figure 12
Tensile properties of Al-Si-Cu-Mg alloy、AMC and HAMC after T6 treatment at 25°C and 250°C: (a) Stress-strain curve of composites at 25°C; (b) Yield strength, tensile strength and elongation of composites at 25°C; (c) Stress-strain curve of composites at 250°C; (d) Yield strength, tensile strength and elongation of composites at 250°C.

The results above indicate that the tensile strength of the Al-Si-Cu-Mg alloy is significantly enhanced at both room temperature and high temperature after the addition of SiCp and CNTs. It is noteworthy that compared to the Al-Si-Cu-Mg alloy, the tensile strength of AMC-CNTs and HAMC exhibits a smaller decrease at 250°C. This can be attributed to several factors. Firstly, CNTs and SiCp themselves possess excellent high-temperature resistance. Secondly, the main reason for the decline in strength of the aluminum alloy at high temperature is the softening effect caused by recovery and recrystallization of the aluminum matrix. After heat treatment, the extruded Al-Si-Cu-Mg alloy retains a significant number of substructures (Figure 8c) that have not fully undergone recrystallization, thereby preserving a significant amount of deformation energy. At elevated temperatures, these substructures continue to grow and form recrystallized grains, which greatly reduces the material’s strength. However, with the addition of SiCp and CNTs, the substructure of the composite is significantly reduced (Figure 8d), making it less susceptible to recrystallization at high temperatures. Furthermore, SiCp and CNTs can inhibit the recrystallization process of the aluminum matrix, thereby playing a “nailing” role in enhancing the composite’s strength.

Figure 13 shows the fracture morphology of Al-Si-Cu-Mg alloy and composite samples stretched at temperatures of 25°C and 250°C. It can be seen that in the fracture morphology of Al-Si-Cu-Mg alloy at 25°C (Figure 12a), the lighter part is the aluminum matrix phase, and the darker part is the dimple. There are obvious dimple with uneven size distribution on the cross section, showing significant ductile fracture, and the fracture mode is mainly the tearing of the aluminum matrix phase. After the crack is generated, it mainly spreads inside the aluminum matrix phase and finally runs through the whole section. It can be seen from the fracture morphology of Al-Si-Cu-Mg alloy at 250°C (Figure 12e) that the dimple at the fracture become large and deep, indicating that the toughness of the alloy increases with the increase of temperature. With the addition of CNTs, the fracture dimples of AMC-CNTs became smaller and shallow with uneven distribution (Figure 13 b.f), and the elongation decreased. The fracture morphology of AMC-SiCp and HAMC composites at temperatures of 25°C and 250°C is similar, and the fracture is smooth without obvious dimple structure (Figure 12c.g), showing obvious brittle fracture. In the process of plastic deformation of alloy matrix, SiCp with high elastic modulus is difficult to cooperate with its deformation, resulting in stress concentration and crack initiation at the junction of the two phases. After crack initiation, it spreads rapidly under the action of tensile stress, and tears the aluminum matrix (circled by blue ellipses) under the action of shear stress, and quickly passes through the entire section, resulting in brittle fracture of the material.

Figure 13
SEM images of fracture morphology at 25°C and 250°C: (a,e) Al-Cu-Mg-Si alloy; (b,f) AMC-CNTs; (c,g) AMC-CNTs; (d,h) HAMC.

4. Conclusion

In this study, AMC-CNTs, AMC-SiCp, and HAMC composites were successfully prepared using a combination of stir casting and hot extrusion technology. The study focused on analyzing the microstructure, mechanical properties, and phase composition of these composites. Based on the findings, the following conclusions can be drawn:

1. The incorporation of SiCp-CNTs in the composites resulted in a more intricate crystal orientation during the extrusion process. This complexity, in turn, facilitated dynamic recrystallization and led to the formation of fine equiaxed grains. SiCp/CNTs were found to enhance the growth of α-Al grains on the (220) and (311) crystal planes, while inhibiting their growth on the (111) and (200) crystal planes.

2. SiCp/CNTs have been found to enhance the dynamic recrystallization (DRX) of the aluminum matrix during the extrusion process and the static recrystallization (SRX) during heat treatment. This promotion of DRX and SRX leads to a more intricate crystal orientation. Additionally, the inclusion of SiCp/CNTs results in significant grain refinement, resulting in the formation of High-Angle Grain Boundaries (HAGBs) through dislocation rearrangement.

3. In the composites, SiC and CNTs were primarily observed to be distributed along the α-Al grain boundaries. This distribution played a significant role in impeding the growth of recrystallized grains. The SiC particles exhibited a homogeneous dispersion with strong interfacial bonding, while the CNTs partially reacted with the Al matrix, forming an in-situ Al₄C₃ intermetallic compound.

4. The addition of SiCp-CNTs significantly accelerated the aging response of Al-Si-Cu-Mg. Compare with Al-Cu-Mg-Si alloy, the tensile strength of HAMC at 25°C and 250°C, from 319MPa and 178MPa to 438MPa and 331MPa.

5. Acknowledgments

This study was financially supported by the Huxiang Youth Talent Project (2022RC1056).

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Publication Dates

• Publication in this collection
  14 June 2024
• Date of issue
  2024

History

• Received
  25 Dec 2023
• Reviewed
  28 Mar 2024
• Accepted
  14 Apr 2024