Enhanced mechanical properties and wear resistance of Al–Cu–Mg alloy fabricated by press-and-sinter processing

Abstract

Two of the main challenges facing the current application of press-and-sinter aluminium alloys are poor densification and insufficient properties such as mechanical properties and wear resistance. In this work, the press-and-sinter Al–4.4Cu–1.6Mg alloy (wt.%) and 5SiC/Al–4.4Cu–1.6Mg (wt.%) composite are successfully developed through optimizing the sintering temperature and heat treatment. The results show that the relative density, hardness and mechanical properties of as-sintered alloy and composite increase with the increase of the sintering temperature. Under the same sintering temperature, the relative density of the composite is comparable to that of the single alloy. When the sintering temperature is 596 °C, both the alloy and composite achieve full densification with a relative density of 99.7% and 99.3%, respectively. Consequently, their hardness and mechanical properties also achieve the optimal value. For the peak-aged state, the alloy exhibits a favourable combination of yield strength, ultimate tensile strength of 325.3 MPa and 394.0 MPa and fracture elongation of 3.8%. The addition of the SiC particles further enhances the yield strength and ultimate tensile strength of the peak-aged alloy to 375.0 MPa and 441.0 MPa, respectively. The wear rate of the peak-aged composite is comparable to that of the 45 steel and is less than half that of the peak-aged alloy due to the enhanced hardness and the stable wear protection provided by the SiC particles.

Keywords

press-and-sinter processing Al–Cu–Mg alloy SiC particles microstructure mechanical properties wear resistance

Introduction

Press-and-sinter aluminium alloys play an important role in aerospace and automobile manufacturing due to their low density, high specific strength, low production cost and high production efficiency.^1,2 For example, the AC-2014 Al alloy camshaft bearing caps prepared by press-and-sinter processing in 1990s are still in use today, with shipment volumes exceeding 300 million.³ However, achieving full densification in the press-and-sinter aluminium alloys remains a significant challenge, primarily due to the dense thermodynamically stable oxide layer covering the surface of the aluminium powder. Such oxide layer can hinder atoms diffusion, powder wetting and liquid spreading, leading to the lower sintered density and the suboptimal mechanical properties.⁴ This severely limits the wider application of the press-and-sinter aluminium alloys.

Therefore, numerous efforts have been undertaken to alleviate or even eradicate the detrimental effects of the oxide layer, mainly including using sintering aids, optimising sintering parameters. Kehl et al.⁵ investigated the sintering process of Al–Cu compact using dilatometry and metallography. The results showed that the Al–Al₂Cu eutectic liquid phase can wet the oxide layer. Lumley et al.⁶examined the effect of Mg element on the sintering of aluminium and its alloys. The results demonstrated that trace Mg could greatly facilitate sintering by breaking up the oxide layer through the formation of MgAl₂O₄ phase. Kondoh et al.⁷ reported that the relative density of Al–12Si alloy increased by 9% with the addition of 1% Mg, and its mechanical properties were also improved. Besides, the eutectic liquid phases formed by the Cu and Mg elements can effectively fill the particle gaps, leading to an increased sintered density.^8–10 Sercombe et al.^11,12 found that the presence of Sn element prolonged the retention of the eutectic liquid phase at the particle boundaries. Benefited from the aforementioned results, various press-and-sinter aluminium alloys (primarily 2xxx, 6xxx and 7xxx series) have been successfully developed and their densification behaviours and mechanical properties have been also investigated. Martı´n et al.¹³ systematically studied the densification behaviour of four commercial aluminium alloy premixed powders (Alumix 123, 321, 13 and 431) under several sintering conditions. Among these, Alumix 123 and 13 corresponded to the 2xxx series, Alumix 321 to the 6xxx series and Alumix 431 to 7xxx series. The results suggested that the optimal relative density for Alumix 123 and 13 reached 95%, while for Alumix 321 and 431, the optimal relative density reached 97%. Du et al.¹⁴ reported a higher sinter density of 98.46% for Al–4.3Cu–1.2Mg alloy achieved by optimising the sintering parameters, but its strength was low, with a ultimate tensile strength (UTS) of only 237.9 MPa. A similar relative density of 98.4% was also observed in a Al–4.4Cu–1.5Mg alloy even with the addition of 0.2 wt.% Sn.¹⁵ Showaiter et al.¹⁶ prepared an Al–1.0Mg–0.6Si–0.25Cu alloy with a relative density of approximately 99% by adding 0.1 wt.% Sn, and its yield strength (YS), UTS and elongation (EL) after T6 heat treatment were 198 MPa, 275 MPa and 10%, respectively. In spite of the great advance in the densification, it should be noted that the press-and-sinter aluminium alloys with superior mechanical properties and wear resistance are still rarely reported.

It is well established in materials science that the introduction of alloying elements such as Cu and Mg into the Al alloys promotes the formation of nanoscale precipitates during heat treatment, thereby enhancing their mechanical strength.^17–20 Additionally, incorporating the SiC particles into these alloys has shown promise in further improving the comprehensive performance by introducing the load-bearing and thermal mismatch effects.^21,22 Herein, the press-and-sinter Al–4.4Cu–1.6Mg (wt.%) alloy and 5SiC/Al–4.4Cu–1.6Mg (wt.%) composite were successfully developed at varying sintering temperature, with the densification behaviour and properties of the as-sintered alloy and composite being investigated. The as-sintered alloy and composite prepared at the optimal sintering temperatures were then selected for heat treatment to further enhance its performance. The microstructure, mechanical properties and tribological behaviours after heat treatment were then systematically analysed. The underlying strengthening mechanisms were also discussed. This work is expected to provide reference for the preparation of high-performance press-and-sinter aluminium alloys.

Experimental methods

Sample preparation

The Al–4.4Cu–1.6Mg alloy powders were prepared by blending elemental powder Al, master alloy powder Al–Cu and Al–Mg, and elemental powder Sn in a mass ratio of 87.5:8.8:3.2:0.5. The addition of Sn element was employed to promote the sintering quality.²³ In addition, 5 wt.% SiC particles were added to the alloy powders to obtain the 5SiC/Al–4.4Cu–1.6Mg composite powders. The detailed information of the raw powders was given in Table 1. Firstly, the alloy and composite powders were individually mixed for 30 minutes in a double cone mixer (ZX-0.02) with an additional 1 wt.% SUW5000B lubricant. Secondly, both blends were compacted into 72 mm × 38 mm × 15 mm green bodies under an axial pressure of 400 MPa. For the subsequent sintering process, differential scanning calorimetry (DSC) was employed to determine the sintering temperature, as shown in Figure 1(a). According to the DSC curves of the Al–4.4Cu–1.6Mg green compact, the persistent liquid phase starts to form at about 560 °C. The formation of persistent liquid promotes densification.²⁴ Besides, the green compact reaches the rapid melting stage at 610 °C, beyond which the phase change will have a large uncertainty.²⁵ Hence, the sintering temperatures is designed to be 576, 586 and 596 °C. The green compacts were sintered in a vacuum dewaxing and sintering furnace (VDSF0408) with a 30-L/min high purity nitrogen (99.99%) as the protective atmosphere in the whole process. The sintering procedure included a 30-minute isothermal delubrication stage at 330 °C followed by an isothermal sintering stage at 576, 586 and 596 °C for 60 minutes, as presented in Figure 1(b).

Table 1.

The detailed information of the raw materials in this work.

Powder	Type	Shape	Particle size (μm)	Purity (wt.%)
Al	Elemental	Spherical	20–35	>99.8
Cu	50:50 Al–Cu master alloy	Spherical	20	>99.5
Mg	50:50 Al–Mg master alloy	Polygon	20	>99.5
Sn	Elemental	Spherical	38	>99.9
SiC	Ceramic	Angular	5–10	>99.8

Heat treatment

The as-sintered alloy and composite with the optimal microstructure and tensile properties were heat treated to further enhance their performance. The selected alloy and composite were solution treated at 510 °C for 30 minutes and then immediately quenched in water. The aging process of solution-treated alloy and composite was conducted at 190 °C for a time range of 0 to 24 hours.

Microstructure characterisation and properties tests

The density of as-sintered alloy and composite was measured by the Archimedes drainage method using an electronic balance (AG285). The microstructures and wear surface morphology were characterised using a scanning electron microscope (SEM, Sigma 300). Besides, the optical microscope (OM, Nikon MA-200) also was used to investigate the morphology of the specimens. The phase constitution was analysed on an X-ray diffraction (XRD, Bruker D8 ADVANCE) with a 2θ range of 20–90° and a scanning speed of 0.2°/s. The precipitation phase and the particle–matrix interface of the peak-aged alloy and composite were also characterised using a transmission electron microscope (TEM, FEI Talos F200x G2).

Dog-bone shaped tensile specimens with a gauge length of 25 mm and a diameter of 5 mm were machined from the as-sinter and peak-aged samples. The uniaxial tensile tests were conducted on an MTS-880 testing machine at room temperature with a nominal strain rate of 1 × 10⁻⁴ s⁻¹. The hardness was examined on an HVS-50Z/LCD Vickers hardness tester with a load of 200 g and a holding time of 15 seconds. At least 5 tests were conducted to ensure the data reproductivity. Dry sliding wear tests were conducted by a ball-on-disc mode on a rotary friction and wear tester (HT-1000 Model) with a load of 10 N, a sliding velocity of 0.3 m/s and a testing time of 60 minutes at the ambient temperature. Three types of the disc samples with a diameter of 30 mm and a height of 10 mm were fabricated from the 45 steel, the peak-aged alloy and composite, and tested at the lower side by sliding against the GCr15 ball at the upper side with a Vickers hardness of HV 786. The 45 steel was tested as a comparison. The weight of the specimens was measured before and after the test and the mass loss ( $Δ m$ ) was divided by the sliding distance (L) to obtain the average wear rate ( $Δ W$ ) during the test:

Δ W = \frac{Δ m}{L}

(1)

Results

Effect of sintering temperature on microstructure and mechanical properties

Figure 2 shows the variation trend of the relative density and hardness of the as-sintered alloy and composite with the sintering temperature. It can be observed that the sintering temperature significantly influences the relative density and hardness of both materials. With the increase in the sintering temperature, the relative density and hardness of both alloy and composite increase rapidly. When the sintering temperature reaches 596 °C, the alloy and composite all exhibit full densification, with a relative density of 99.7% and 99.3%, respectively. This is mainly attributed to the higher liquid phase content at elevated temperature.¹³ Correspondingly, the as-sintered alloy and composite achieve their maximum hardness values at 596 °C, with values of HV 97.1 and HV 101.0, respectively. Besides, there is minimal difference in the densification level between the alloy and composite at all sintering temperatures which is aligned with the results reported by Asgharzadeh²⁶ that the incorporating low content (< 9 vol. %) of the SiC particles has a marginal impact on the relative density.

Figure 3 gives the OM images of the as-sintered alloy and composite at different sintering temperature. As shown in Figure 3, the microstructure of the as-sintered alloy mainly consists of the Al matrix, Al₂Cu phase and pores, while the composite includes an additional SiC phase. Besides, the densification behaviour of both the alloy and composite exhibits significant differences at varying sintering temperatures. For the alloy sintered at 576 °C (Figure 3(a)), some large irregular pores are visible, indicating inadequate sintering. As the temperature increases, the rapid densification process is triggered. When the sintering temperature reaches 596 °C (Figure 3(c)), only a minimal amount of spherical microvoids remains. This is mostly caused by the higher liquid phase content during the sintering process.¹⁵ From Figure 3(d)–(f), the densification behaviour of the composite is similar to that of the alloy, indicating that the SiC particles have little effect on the densification behaviour of the Al alloy. The observations from the OM images are consistent with the relative density results depicted in Figure 2. In addition, from Figure 3(f), it can be observed that the SiC particles are uniformly distributed in the Al matrix and well bonded to the matrix at a sintering temperature of 596 °C.

Figure 4(a) displays the tensile engineering stress–strain curves of the as-sintered alloy and composite at room temperature. To more intuitively show the effect of the sintering temperature on the mechanical properties, the corresponding UTS, YS and EL are plotted against the sintering temperature, as depicted in Figure 4(b)–(d). The variation in the tensile properties closely follows the trend of the relative density and hardness. It can be found that the alloy and the composite fabricated at 576 °C exhibit the poorest mechanical properties due to the excessive pores. With the increase in the sintering temperature from 576 °C to 596 °C, both the alloy and composite simultaneous achieve their optimal properties. The YS, UTS and EL of the alloy are 167.6 MPa, 290.7 MPa and 6.7%, respectively. For the composite, the YS and UTS are increased to 189.0 MPa and 318.3 MPa, respectively, but the EL decreases to 5.6%.

The fracture morphologies of the as-sintered alloy and composite at different sintering temperatures are presented in Figure 5. In Figure 5(a) and (d), the obvious pores and original particles boundaries are observed at the fracture surface of the alloy and composite sintered at 576 °C, which can serve as crack initiation site and stress concentration points, and lead to the premature failure. Compared to those samples sintered at 576 °C, fewer original particles boundaries and pores can be found at the fracture surface of the alloy and composite sintered at 586 °C, as shown in Figure 5(b) and (e). This indicates that there is a relatively high interfacial bonding strength between the sintered powders. Different from the morphology of the alloy and composite sintered at 576 °C and 586 °C, a large number shallow dimples are observed at the fracture surface of 596 °C, without apparent original particle boundaries and pores, as depicted in Figure 5(c) and (f). Based on the above microstructure and mechanical properties results, the alloy and composite sintered at 596 °C have the optimal mechanical properties, which are selected for further heat treatments and analysis.

Effect of heat treatment on microstructure and properties

Figure 6 exhibits the hardness of the solid-solutioned alloy and composite as a function of the aging time. It can be observed that the addition of the SiC particles leads to a noticeable improvement in the peak-aged hardness of the composite, as compared to the unreinforced alloy. The peak-aged hardness of the alloy and composite are HV 144.2 and HV 153.4, respectively. The increased hardness is attributed to the rigidity of the SiC particle, which impedes the deformation of the Al matrix. Furthermore, it is noticed that the time to reach the peak-aged state is extended from 6 hours for the alloy to 8 hours for the composite, indicating that the aging kinetics of the alloy is decelerated by the SiC particle.²⁷

The microstructures of the peak-aged alloy and composite are analysed by OM and SEM, with the results presented in Figure 7. From Figure 7(a) and (d), it can be seen that no coarse Al₂Cu precipitates are found in the peak-aged alloy and composite, which have been dissolved into the matrix during the solid solution. Besides, the average grain size of the peak-aged alloy is 17.3 ± 1.0 μm and is refined to 14.1 ± 0.9 μm for the composite. The microstructures of the peak-aged alloy under SEM in Figure 7(b) reveal that long rod-shaped precipitates are formed at the grain boundaries and are mainly rich in the Mg and O elements as determined by the EDS mappings in Figure 7(c), which are likely of the MgAl₂O₄ phase. The O element mainly comes from the oxidation of the Al matrix during exposure to the air. The inserted XRD results in Figure 7(b) all exhibit two sets of diffraction peaks corresponding to the Al matrix and the Mg₂Sn precipitate while extra peaks of the SiC phase are presented in the composite. Other phases are not detected because of their very small sizes and/or low content.²⁸ Figure 7(e) displays an individual SiC particle in the peak-aged composite. No obvious defects and porosity were observed at the SiC/Al interface, suggesting the good bonding between them. The corresponding EDS analysis in Figure 7(f) indicates the presence of precipitates rich in the Mg and O elements at the SiC/Al interface.

Figure 8 presents the high-angle annular dark-field (HAADF) images and the corresponding EDS mappings of the intra-granular precipitates in the peak-aged alloy and composite samples. It can be observed that the acicular S precipitates, ranging in length from tens to hundreds of nanometers, are uniformly distributed within the peak-aged alloy and composite. According to the previous studies,^29,30 when the Cu/Mg mass ratio ranges from 1.5 to 5, the predominant precipitate in the Al matrix is mainly the S phase. The EDS results further confirm that the acicular phase in the peak-aged alloy and composite is indeed this S phase. It is expected that these nanoscale precipitates phases will significantly enhance the strength.

Figure 9(a) is an overall view of the SiC/Al interface in the peak-aged composite. Further observation at high magnification presents a clean and defect-free SiC/Al interface in Figure 9(b) along with high-density dislocations (marked with the white arrow) around. The corresponding EDS mapping in Figure 9(c) also identifies the enriched Mg, Cu and O elements at the SiC/Al interface, similar with Figure 7. The distribution of the Cu and Mg elements is partially superposed, and that of the Mg and O elements is completely overlapped and forms a near-continuous interface layer, suggesting the presence of the second phases. The FFT analyses into the square area in Figure 9(d) and (e) confirm that the former is of the major precipitate Al₂CuMg in the 2024 aluminium alloys,³¹ while the latter is the MgAl₂O₄ phase (Figure 9(h)). From Figure 9(d) and (e), the SiC particle is tightly bonded with the Al₂CuMg and MgAl₂O₄ phases. The presence of the MgAl₂O₄ interface layer has two favourable effects on enhancing the performance of the composite. Firstly, it can serve as a diffusion barrier, effectively impeding the formation of the Al₄C₃ phase.^32,33 Secondly, it improves the interface bonding strength between the SiC particle and the Al matrix,^34,35 thereby guaranteeing the better mutual load transfer effect. The FFT analysis of the square area in Figure 9(f) only presents two sets of spots corresponding to the SiC and Al matrix, that is, no Al₄C₃ phase is formed. According to Refs.,^36,37 the SiC phase becomes thermodynamically unstable beyond certain temperatures and reacts with the matrix. Therefore, the experimental results demonstrates that intentionally selecting a lowered sintering temperature as mentioned before is indeed effective in preventing forming this deleterious phase.³⁸ Overall, a clear bonding without apparent interfacial reaction and defects is formed between the SiC particle and the Al matrix in the composite through the press-and-sinter method in this work.

The room temperature tensile stress–strain curves for the peak-aged alloy and composite samples are presented in Figure 10(a), with their YS, UTS and EL in Table 2. The peak-aged alloy exhibits a favourable combination of YS, UTS of 325.3 MPa and 394 MPa and EL of 3.8%. The addition of 5 wt.% SiC particles increases the YS and UTS of the alloy to 375.0 MPa and 441.0 MPa, respectively. The present results are compared with those of the available aluminium alloys prepared utilizing press-and-sinter method,^39–47 as depicted in Figure 10(b). As far as we know that the 7xxx series alloy demonstrates the highest strength among the press-and-sinter aluminium alloys. For example, the UTS of Alumix 431D is 448.0 MPa, but its ductility is low.⁴⁷ In comparison, the properties of the peak-aged alloy in this work are in an outstanding position and outperform most of the mechanical properties in the relevant press-and-sinter aluminium alloy systems. The incorporation of the low-cost SiC particles into the Al–Cu–Mg alloy not only achieves additional enhancement of its strength but also offers significant cost advantages, providing a feasible strategy for improving press-and-sinter aluminium alloy for industrial production.

Table 2.

The mechanical properties of the peak-aged alloy and composite.

Materials	YS (MPa)	UTS (MPa)	EL (%)
Al–Cu–Mg alloy	325.3 ± 3.2	394.0 ± 4.4	3.8 ± 0.3
SiC/Al–Cu–Mg composite	375.0 ± 4.6	441.0 ± 13.7	2.2 ± 0.3

The fracture surfaces of the peak-aged alloy and composite are observed and displayed in Figure 11. Nearly no porosities are presented in both materials, consistent with the high relative density. As can be seen in Figure 11(a), there are large-sized cleavage planes and some shallow dimples (indicated by the yellow arrows) in the alloy, indicating the mixed ductile and brittle fracture nature. For the composite, the fracture mode changes to a more brittle pattern with the dominated cleavage fracture of the SiC particles and the matrix. Interestingly, many fractured but not debonded SiC particles can be observed in Figure 11(b). According to Refs.,^48,49 particle debonding will occur before particle fracture only if the bonding strength of the interface is lower than that of the particle. This implies the high SiC/Al interface quality, and resultantly the load can be effectively transferred between the matrix and the ceramic particles through the interface, enhancing the overall properties of the composite.³⁷ From Figure 11(c), some fine dimples surrounding the fractured SiC particles can be observed, further proving the high bonding strength.⁴⁹

The friction and wear test results of the peak-aged alloy and composite are shown in Figure 12. The coefficient of friction (COF) in Figure 12(a) increases rapidly at the beginning and then gradually stabilises when the time exceeds 20 minutes. Figure 12(b) depicts the average COF and wear rate of the peak-aged materials, where the former is the averaged value during the stable stage of 20 to 60 minutes and the latter corresponds to the whole process. The average COF of the peak-aged alloy and composite is 0.29 and 0.40, respectively. The higher average COF of composite is mainly due to the increased frictional force caused by the interaction of the SiC particles with the coupled ball.⁵⁰ Even with the slighter higher COF, the wear rate of the composite is only 3.24 mg/km, which is comparable to that of the 45 steel and 53.3% lower than that of the alloy. The result indicates that the addition of the SiC particle through the traditional press-and-sinter method effectively improves the wear resistance of the peak-aged alloy.

Figure 13(a) and (c) gives the wear tracks on the worn surface of the peak-aged alloy and composite. The width of the wear track in the composite is 1031.3 μm, which is much smaller than that in the alloy (1387.8 μm). A broader and deeper wear track represents a lower wear resistance,⁴⁹ consistent with the results in Figure 12(b). The worn surface of the alloy and composite at higher magnification is shown in Figure 13(b) and (d). For the alloy, big craters and delamination accompanied by the large plastic deformation are observed, confirming that the dominant wear mechanism is adhesive wear. In the case of the composite, parallel and shallow grooves along with small craters are found, and therefore the main wear mechanism is changed to mainly abrasive wear and partly adhesive wear. The worn hard particles like the SiC phase acts as the abrasive in this process.⁵¹

Figure 14 is the wear debris morphology of the peak-aged alloy and composite, which greatly differs from each other in the shape and size, that is, displaying big disc-shaped fragments in the alloy but much smaller ones along with some many fine particles in the composite in Figure 14(a) and (d), which is related to the changed wear mechanism. The morphology of one typical wear debris from the two materials is presented in Figure 14(b) and (e). It can be found that the surface of the wear debris in the alloy is smooth but rough in the composite. Furthermore, the EDS analysis into the wear debris in Figure 14(c) and (f) indicates that oxidation wear has occurred in both the alloy and composite. The higher O element content in the debris of the composite can be primarily ascribed to the increased frictional heat resulting from the enhanced friction force, which induces more severe oxidation.⁵²

Discussion

Strengthening mechanisms

After peak aging, the YS of the alloy and composite is significantly enhanced, and can be related to strengthening mechanisms induced by the grain boundary, precipitate, load transfer and thermal mismatch. The latter two mechanisms represent the strength contributions brought about by the introduction of SiC particles in composite. Herein, the involved strengthening mechanisms are analysed in detail for the peak-aged alloy and composite.

The increase in the YS resulting from the grain boundary strengthening can be estimated using the Hall–Petch formula³⁴:

σ_{GB} = k d^{- 1 / 2}

(2)

where k = 0.1 MNm^−3/2 is the Hall–Petch constant for the Al matrix, d is the average grain size. Therefore, the value of

σ_{GB}

is deduced to be 24.0 MPa and 26.6 MPa for the peak-aged alloy and composite, respectively.

The precipitate strengthening provided by the acicular phase (S phase) is calculated according to the following equation⁵³:

\begin{aligned} σ_{ps} = M \frac{0.13 G b}{2 \sqrt{r h}} \\ [f_{v}^{1 / 2} + 0.75 {(\frac{r}{h})}^{1 / 2} f_{v} + 0.14 (\frac{r}{h}) f_{v}^{3 / 2}] \ln (\frac{0.158 r}{r_{0}}) \end{aligned}

(3)

where M = 3 is the Taylor factor for the FCC polycrystalline matrix, G = 25.4 GPa is the shear modulus of the Al matrix, b = 0.286 nm is the Burgers vector of the matrix, r₀ is the inner radius, its value is generally considered to be equivalent to the Burgers vector, r, h and f_v are the radius, half-thickness and volume fraction of the S phase, respectively. According to the microstructure in Figure 8, the value of r, h and f_v are evaluated to be 121.8 nm, 4.3 nm and 13.9% for the alloy and 89.3 nm, 4.7 nm and 15.4% for the peak-aged composite. The value of the precipitate strengthening is 292.7 MPa and 287.8 MPa for the peak-aged alloy and composite, respectively.

The SiC particles can bear the load transferred from the matrix, which increases the YS. The strength increment can be calculated²¹:

σ_{L - T} = σ_{m} [\frac{V_{r} (S + 2)}{2} + V_{m}] = 0.5 V_{r} σ_{m}

(4)

where σ_m = 325.3 MPa is the YS of the peak-aged alloy, V_r = 0.0434 and V_m = 0.9566 are the volume fraction of the SiC particles and the Al matrix, respectively, S = 1.45 is the mean aspect ratio of the SiC particles. Therefore, the contribution of the load transfer is 10.2 MPa.

The increased dislocation density arising from the different coefficient of thermal expansion (CTE) between the SiC particles and the matrix can enhance the YS²¹:

σ_{CTE} = β G b \sqrt{\frac{12 \sqrt{2} Δ α Δ T V_{r}}{d_{p} b (1 - V_{r})}}

(5)

where β = 1.25 is the strengthening coefficient, G and b have the same definitions as above, Δα = 19.3 × 10⁻⁶/K is the CTE difference between the SiC particle and the Al matrix, ΔT = 570 K is the difference in the sintering and room temperature, d_p = 9 μm is the average diameter of the SiC particles. Therefore, the value of the dislocations strengthening is 17.0 MPa.

Based on the above calculations, it is evident that the precipitation strengthening makes the major contributions to the enhanced YS of the peak-aged alloy and composite. This is primarily attributed to the nano-sized precipitates effectively hindering the movement of dislocation and thus improving the mechanical properties of the alloy and composite. In addition, the further enhanced strength of the peak-aged composite is related to the extra strengthening effects of the load transfer and thermal mismatch introduced by the incorporation of SiC particles.

Effect of SiC on wear resistance

Figure 15 is the subsurface of the peak-aged alloy and composite after the wear test. From Figure 15(a), a large peeling area with its length reaching 100 μm can be observed in the alloy, in agreement with the severe plastic deformation observed in Figure 13(a) and (b). In addition, a mechanical mixed layer (MML) with a thickness of several micrometers is formed, as shown in Figure 15(d). Generally, a dense MML can effectively reduce the COF and thus the wear rate by preventing the direct contact between the ball and disc.⁵⁴ However, defects characterised by the many cracks are observed in the MML of the alloy, indicating its poor protection to the worn surface.⁵⁵ The cracks in the MML propagate and bridge with each other, resulting in the easy delamination and peeling. As for the composite, the SiC particles are protruded without pulling out on the worn surface in Figure 15(b) and (e). This result again demonstrates the good bonding between the Al matrix and the SiC particles and also the protection effect of the micro-sized SiC particles on the matrix by bearing the load and resisting the wear. The MML cannot be formed in the composite mainly due to the high rigidity of the SiC particles.

According to Figure 15, the improved wear resistance can be related to the added SiC particles. An illustration to depict the varied wear mechanism at different stages is shown in Figure 16 for a better understanding. As displayed in Figure 16(a), for the softer Al alloy, the plastic deformation easily occurs under the applied load and the affected materials are sheared flat in a large piece as in Figure 14(a), especially under the additional friction heat.⁵⁶ With the progress of the test, the MML is produced due to repeatedly rolling and stretching of the wear debris.⁵⁷ However, the MML with cracks fails to protecting for the worn surface due to its rapid delamination and peeling off under the cyclic load. Under this condition, the alloy exhibits a poor resistance with the low strength and hardness. For the composite in Figure 16(b), both the SiC particles and the Al matrix are in sliding contact with the counterpart at the beginning. Due to the much higher hardness of the reinforcements, the Al matrix on the top surface is gradually removed and the SiC particles are protruded.^58,59 The worn SiC particles trapped inside changes the wear pattern to the abrasive mode, resulting in the formation of the finer flake-like debris in Figure 14(d). At this stage, the protruding SiC particles greatly shield the matrix from the rapid wear while the load and stress born by the particles is effectively transferred and relieved through the strong interface to the matrix.^60–62 It is worth noting that the strong bonding between the SiC particles and the Al matrix prevents the detachment of the SiC particles, ensuring that the Al matrix can stable support the SiC particles to resist the wear. Furthermore, the added reinforcements also enhance the hardness and the deformation resistance of the composites as compared to the alloy.⁶³ These effects are synergized together to enhance the wear resistance of the peak-aged composite.

Conclusions

In this article, fully dense Al–Cu–Mg alloy and SiC/Al–Cu–Mg composite were successfully developed using traditional press-and-sinter processing, and an optimised heat treatment process was employed to further enhance the properties of both the alloy and composite. The microstructure of as-sintered and peak-aged samples was observed through the utilisation of XRD, OM, SEM and TEM. The tensile properties and dry sliding wear performance of the peak-aged alloy and composite were investigated and the corresponding mechanisms were analysed. The main results obtained in this study are summarised as follow:

Elevating the sintering temperature from 576 °C to 596 °C causes a significant improvement in the relative density, tensile strength and elongation of the Al–Cu–Mg alloy and SiC/Al–Cu–Mg composite. The alloy and composite exhibit the highest relative density after sintering at 596 °C, with value of 99.7% and 99.3%, respectively. Correspondingly, their mechanical properties also achieve optimal value.

The peak-aged alloy exhibits a favourable combination of YS, UTS of 325.3 MPa and 394.0 MPa and EL of 3.8% at room temperature. The addition of 5 wt. % SiC particles further enhances the YS and UTS of the alloy to 375.0 MPa and 441.0 MPa, respectively. The precipitation strengthening is primarily responsible for the improvement in strength.

The wear rate of the peak-aged composite is comparable to that of the 45 steel but is less than half that of the peak-aged alloy. The higher wear resistance of the composite is attributed to the enhanced hardness and the stable wear protection provided by SiC particles.

Footnotes

Acknowledgements

The authors are grateful for the financial support of the National Key Laboratory of Electromagnetic Energy under contract NO. 2022ZC035. The authors thank Dongguan Gehong Aluminum Powder Metallurgy Co., Ltd, Shenzhen Kadam Technology Co., Ltd and Xi'an Fulai Electrical Alloy Co., Ltd for raw material supply and samples preparation. The authors also thank Prof. Kemin Li of Henan University of Technology for the invaluable discussions as regards the wear results. The authors appreciated the support from Instrument Analysis Center of XJTU with respect to the SEM test.

Authors’ contributions

Pengpeng Huang did experiments, investigation, writing-original draft and data curation. Letian Wang did investigation and methodology. Zhongyue Yang did investigation and data curation. Yake Wu did conceptualisation and writing-review and editing. Zhenzhuo Ling did investigation and resources. Qianhang Cui did investigation and data curation. Hualei Wang did experiments and investigation. Feng Jiang did conceptualisation, supervision, writing-review and editing and funding acquisition.

Data availability

The data that support the findings of this study is available from the corresponding authors upon reasonable request.

Declaration of conflicting interests

The authors declared no potential conflicts of interest with respect to the research, authorship and/or publication of this article.

Funding

The authors disclosed receipt of the following financial support for the research, authorship and/or publication of this article: This work was supported by the National Key Laboratory of Electromagnetic Energy (grant number 2022ZC035).

ORCID iD

Feng Jiang

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