Characterization of (B₄C+Ti) Hybrid Reinforced Mg and AZ91D Composites

Abstract

Ceramic hybrid particulate reinforced magnesium matrix composites have attracted much attention in recent years, owing to their light weight, high specific strength and stiffness, excellent wear and damping properties, and have very wide popularization and application prospects in the automotive and aerospace fields. A simple and novel technique, in situ reactive infiltration technique, was utilized to prepare (B₄C+Ti) hybrid reinforced Mg and AZ91D composites. The microstructure, produced phases, wear, compression and bending performance were characterized and analyzed in depth. The results showed that extending the holding time for preparing the composites and magnesium alloy as the matrix were both beneficial for the generation of interpenetrating networks within the composites, which means more sufficient and complete in situ reactions between B₄C and Ti particles. More sufficient and complete reactions facilitate stronger interfacial bonding, leading to an improvement of the wear, compression and bending performance of the composites. The wear mechanisms for both kinds of composites were abrasive and delamination wear. (B₄C+Ti)/AZ91D composites exhibit better compression performance and this can be attributed to the microstructure of the composites and work-hardening and softening during compression.

1. Introduction

Hybrid metal matrix composites (HMMC) refer to composites that consist of more than one kind of reinforcement with different sizes and shapes, which will lead to higher strength, higher resistance to wear and other excellent properties. In addition, some of the drawbacks of one type of reinforcement can be compensated by additional reinforcement [1,2]. Mohammad et al. [3] studied the tribological behavior of mono and hybrid ceramic-reinforced AZ31 matrix composites, and the results showed that hybrid composites demonstrated higher wear resistance than the mono ones. Li et al. [4] proposed an effective method to achieve excellent performance of Mg matrix composites by incorporating carbon nanotubes (CNTs) and in-situ micro-Mg₂Si_m and nano-Mg₂Si_n into magnesium matrix, and the results showed that the hardness, compressive ultimate strength, compressive yield strength and elongation of the HMMC increased by 15.6%, 53.2%, 50.4% and 2.9%, respectively, compared to the AZ91D matrix. HMMCs have been a hot issue in the high strength applications field in the past decades, owing to their excellent performance [1].

Boron carbide (B₄C) has high hardness, low density, high melting point, good chemical stability, excellent wear resistance and many other advantages [5]. In our previous study [6], Ti effectively improved the wettability of the metal melt with ceramic preform during fabricating composites by metal melt infiltration. In addition, the reaction product of the B₄C-Ti system was TiC and TiB₂, and they have merits of high melting point, light weight, high hardness and modulus, superior wear resistance, and excellent chemical stability, and rather well wettability exists between TiC and Mg [7]. Therefore, B₄C and Ti are appropriate, promising reinforcing candidates for magnesium or magnesium alloy.

The conventional methods to prepare HMMCs include mechanical alloying [8], powder metallurgy [9], stir casting [10], and infiltration routes with or without pressure [11]. With the fast development of science and technology, many novel processing methods for HMMCs have been developed, such as the precursor infiltration and pyrolysis method [12], chemical vapor deposition [13], electrical discharge machining [14] and in situ reactive infiltration [7]. Reinforcement in HMMCs prepared via in situ reactive infiltration can be in situ synthesized and the melt matrix can infiltrate the preform simultaneously. By using this method, fine and thermal stable reinforcement and a clean interface between the reinforcement and matrix can be uniformly produced within the matrix, which ensure the improvement of the mechanical performance of the composites [15].

Recently, ceramic hybrid particulate reinforced magnesium matrix composites have received much attention due to their light weight, high specific strength and stiffness, excellent wear and damping properties and other isotropic characteristics. They are the most potential candidates for structural and functional applications in aerospace, automotive and electronic devices. Many researchers have conducted studies about the mechanical properties about the hybrid particulate reinforced magnesium matrix composites. Shen et al. [16] studied the tensile property of the micro+submicro+nano SiC-reinforced magnesium matrix composites prepared by stir casting, and found that the tensile strength were significantly improved, owing to the effective load transfer from the matrix to the reinforcement material and the decrease in the grain size. Sankaranarayanan et al. [17] conducted a study about the micro-sized titanium and nano-sized boron carbide particle reinforced magnesium matrix composites prepared by the liquid metallurgy route, based on the disintegrated melt deposition technique, and the results showed that the best combination of strength and ductility was observed in the hybrid composite, which can be attributed to the presence of nano particles, the uniform distribution of the reinforcement and the strong interfacial bonding between the matrix and reinforcement. Sahoo et al. [18] investigated the age hardening behavior and mechanical properties of the as-cast in-situ TiC-TiB₂-reinforced AZ91 composites, and found that the ultimate tensile strength of the composites was increased by 66.49% after age hardening, compared with the base monolithic alloy. However, reports about the mechanical properties of hybrid magnesium matrix composites prepared by in situ reactive infiltration are rare.

In this study, in situ reactive infiltration was utilized to prepare (B₄C+Ti) hybrid particulate reinforced magnesium matrix composites, and the microstructure, produced phases and mechanical properties, including the compressive, three-point bending and wear properties, of (B₄C+Ti)/Mg and (B₄C+Ti)/AZ91D composites will be discussed in detail. Influence of the holding time and matrix on the microstructure, produced phases and mechanical properties will be mainly investigated.

2. Materials and Methods

The starting raw materials used were B₄C particles, with an average particle size of 28 μm and a purity of 94.66%, purchased from Mudanjiang Precision Abrasive Boron carbide Co., Ltd., (Mudanjiang, China), and Ti particles, with an average size of 44 μm and a purity of 99.5%, provided by Beijing Xingrongyuan Technology Co., Ltd., (Beijing, China). The pure magnesium and AZ91D ingot used were purchased from Yueyang Yuhua New Metallurgical Materials Co., Ltd., (Yueyang, China), and pure magnesium has a purity of 99.95%, and the chemical composition of the AZ91D magnesium alloy is Mg-9.5%Al-0.5%Zn-0.21%Mn (wt%).

Before the preparation, an STA449F3 simultaneous thermal analyzer was used to carry out the differential thermal analysis of the mixed powder, and according to the DTA-TG curve of the mixed powder in Figure 1, the peak value of the reaction system was 1213 K, which indicates a relatively complete reaction between the starting materials. Considering that pure Mg or Mg alloy are chemically active and flammable at higher temperatures, 1173 K was chosen as the holding temperature.

Based on the results of DTA, the detailed processing procedures for (B₄C+Ti)/Mg and (B₄C+Ti)/AZ91D composites by the in situ reactive infiltration method can be summarized as follows. Firstly, Ti and B₄C powders (in a molar ratio of 3:1, as shown in their equilibrium reaction equation 3Ti + B₄C = TiC + 2TiB₂) were mixed and mechanical blended after addition of a certain amount of binder (polyvinyl alcohol solution). Then, the mixed powders were compacted into a green preform of rectangular solid with a size of 60 mm × 25 mm × 20 mm and 50% relative density. Together with a pure Mg or magnesium alloy ingot on it, the compacted preform was then put into a graphite mold. The in situ reactive infiltration process was finally performed in an electric furnace under the presence of a flowing argon atmosphere (99.999% purity). The infiltration system was finally heated up to 1173 K for 30 min, 90 min, and 180 min, respectively, to study the effect of the holding time on the mechanical properties at a rate of 5 K min⁻¹. After that, the samples were cooled with the furnace down to the room temperature.

A pin-on-disc friction and wear tester (MM-200) was used to test the dry sliding wear performance. The specimens had a size of ø5.95 mm ×12 mm and the steel counter disc was 45# steel, with a hardness of HRC 56. It should be noted that the contacting surfaces of the specimens and the counter discs should be ground to 800-grit abrasive paper before the test. The sliding speed and time were 250 r/min and 5 min, respectively, and the applied load was set as 40 N. Wear rates were calculated by the ratio of volume loss to wear distance and the volume losses were obtained by the ratio of mass loss to the density of the specimens. To confirm the mass loss precisely, the specimens and the counter discs should be washed by acetone in an ultrasonic cleaning machine for 10 min and then weighted by using an electronic balance with a precision of 0.01 mg before and after the test.

The compression and bending behaviors of (B₄C+Ti)/Mg and (B₄C+Ti)/AZ91D composites were tested by using an electronic universal testing machine (CMT5105). The specimens for compressive and bending behavior had a size of ø4 mm× 6 mm and 40 mm× 4 mm × 4 mm, respectively, and the upper and lower surfaces of the specimens had to be polished by abrasive paper prior to the test in order to reduce the friction resistance between the sample and the indenter during the performance test. The span for the bending performance test was 30 mm, and the indenter velocity was set as 1 mm/min.

A Mettler-Toledo analytical balance (XS105 DualRange) and a full-automatic Rockwell type hardness tester were used to measure the density according to Archimedes’s principle and the hardness of the as-fabricated composites. Scanning electron microscopy (SEM, FEI Quanta 600, Massachusetts, America) was used to characterize the microstructure, worn morphology and the bending fracture surfaces of (B₄C+Ti)/Mg and (B₄C+Ti)/AZ91D composites. X-ray diffraction (XRD, X’Pert Pro, PANalytical, Netherlands) was utilized to determine the produced phases within the composites.

3. Results

3.1. Microstructure and Produced Phases

Figure 2 shows the SEM images and EDS results of the (B₄C+Ti)/Mg and (B₄C+Ti)/AZ91D composites prepared for different holding time. From Figure 2a–d, it can be observed that when the composites were prepared for 30 min or 90 min, black B₄C and gray Ti particulates with irregular shapes were distributed uniformly within the matrix, and no segregation or obvious voids or pores were observed. In addition, white phases were produced around the B₄C and Ti particulates, and it was confirmed to be titanium carbide or titanium boride by EDS, which means that there is an in situ reaction between the starting powders. However, according to Wang [19], it is not easy to distinguish titanium boride from titanium carbide by EDS due to the existence of Ti in both of them and the lightness of boron and carbon elements. In (B₄C+Ti)/Mg composites, when the holding time was prolonged from 30 min (Figure 2a) to 90 min (Figure 2c), lots of finer particles appeared, and it was confirmed to be TiC by EDS. This can be attributed to particle fragmentation of Ti and the reaction between Ti and dissociative C diffused away from B₄C with the time extension. As the time continued to extend to 180 min, the microstructure of (B₄C+Ti)/Mg composite changed to be interpenetrating network and titanium boride and titanium carbide were produced within the matrix (Figure 2e). In the (B₄C+Ti)/AZ91D composites (Figure 2b,d,f), it can be observed that when the holding time was extended from 30 min to 90 min, the microstructure tended to change to be interpenetrating network (Figure 2d). When the holding time was 180 min (Figure 2f), the microstructure was still interpenetrating network and even no starting particles were observed. From the comparison of the microstructure of both kinds of composites prepared under different holding time, it can be deduced that the in situ reaction between the starting powders tended to be more sufficient as the holding time increased. In addition, taking the AZ91D alloy as the matrix instead of pure Mg, the in situ reaction between B₄C and Ti tended to be more complete and sufficient.

XRD spectra, identified by the experimental PXRD Database of the as-fabricated (B₄C+Ti)/Mg and (B₄C+Ti)/AZ91D composites, is displayed in Figure 3. As shown in Figure 3a, when the holding time was 30 min, the main produced phases within the (B₄C+Ti)/Mg composite was B₁₃C₂, Ti, Mg, TiB and TiC. It has been pointed out that the original reactant B₄C can be transformed into B₁₃C₂, mainly because with the decrease in carbon concentration in B₄C, element B replaces element C on the icosahedral chain, making the C-B-C chain become the C-B-B chain [20]. The peak of TiB and TiC verified the preliminary in situ reaction between B₄C and Ti and the peak of Mg illustrated the successful infiltration of Mg melt into the ceramic preform. When the holding time was prolonged to 90 min, the main phases were B₁₃C₂, Ti, Mg, TiB, TiC and Ti₃B₄. As the time extended to 180 min, the main phases were B₁₃C₂, Mg, TiB, TiC, Ti₃B₄ and TiB₂. By comparing the peaks of the produced phases, it can be observed that the peaks of B₁₃C₂ and Ti gradually decreased, while that of TiB, TiC and Ti₃B₄ increased as time extended because the reaction between B₄C and Ti was strengthened. As the holding time prolonged to 180 min, TiB₂ was also detected, while B₁₃C₂ and Ti were almost invisible, which proves a more sufficient and complete in situ reaction between B₄C and Ti as the holding time extends. In addition, the peak of Mg was increased with the extension of holding time, and this is mainly because sufficient time allows the infiltration of the molten metal matrix into the ceramic reinforcement to be more sufficient.

It could be observed from Figure 3b that when the holding time was 30 min, the main phases of the (B₄C+Ti)/AZ91D composite was B₁₃C₂, Ti, Mg, TiB and TiC. As the holding time was extended to 90 min, TiB₂ was also detected, in addition to those phases observed as 30 min. When the holding time was 180 min, the main phases were Ti, Mg, TiB, TiC and TiB₂, and rare peak of B₁₃C₂ was detected. It can be inferred by the observation above that the AZ91D alloy facilitates the reaction between B₄C and Ti, making the reaction more sufficient and complete compared to pure Mg, as observed in the SEM images of the corresponding composites. The main reason is that the wettability between the Mg alloy and the ceramic reinforcement is better than that of pure Mg, making the infiltration of the molten metal matrix into the ceramic reinforcement easier. The infiltration of the metal matrix among the reinforcement promotes the diffusion of B and C element to diffuse away from B₄C, and then promotes the reaction between B or C element and Ti.

Table 1. Density and hardness of (B₄C+Ti) hybrid-reinforced Mg and AZ91D composites.

Holding Time/min	(B4C+Ti)/Mg Composites		(B4C+Ti)/AZ91D Composites
	ρ/(g/cm3)	HRB	ρ/(g/cm3)	HRB
0	2.77 ± 0.001	40.16 ± 0.01	2.81 ± 0.001	45.32 ± 0.01
30	2.53 ± 0.001	55.55 ± 0.01	2.64 ± 0.001	59.14 ± 0.01
90	2.49 ± 0.001	56.74 ± 0.01	2.50 ± 0.001	69.43 ± 0.01
180	2.48 ± 0.001	72.95 ± 0.01	2.68 ± 0.001	97.90 ± 0.01

shows the physical performance of the as-fabricated (B₄C+Ti)/Mg and (B₄C+Ti)/AZ91D composites. When the samples were heated to 1173 K without any holding time, the density and hardness of the (B₄C+Ti)/Mg composite was about 2.77 g/cm³ and 40.16 HRB, while that of the (B₄C+Ti)/AZ91D composite was 2.81 g/cm³ and 45.32 HRB. When the samples were heated with a specific holding time (30 min, 90 min or 180 min), the density tended to decrease, while that of hardness increased, and it could be observed that the density of those composites ranged from 2.47 g/cm³ to 2.69 g/cm³, and hardness from 50 HRB to 98 HRB. The density was even lower than most commercial alloys, so light-weight hybrid reinforced magnesium matrix composites were successfully prepared by in situ reactive infiltration technique. In addition, by comparing the density and hardness of both kinds of composites, it could be observed that by extending the holding time, the density of the composites was on a downward trend, while the hardness was on the rise. Through the above analysis, extending the holding time makes the reaction between the starting materials stronger, thus leading to stronger interfacial strength. Furthermore, it is obvious that (B₄C+Ti)/AZ91D composites tended to be denser and harder than (B₄C+Ti)/Mg composites fabricated under the same processing conditions, and the reason is that taking AZ91D as the matrix also leads to a more complete and sufficient reaction between the starting powder. In addition, stronger interfacial strength means that their performance of local resistance to hard material pressing into its surface is better, i.e., higher hardness. As mentioned above, for both kinds of composites, extending the holding time and taking the AZ91D alloy as the matrix contribute to a higher hardness.

3.2. Wear Performance

Figure 4 shows the relationship plot of the wear rate and holding time for both kinds of composites. It is obvious that the wear rate of the (B₄C+Ti)/Mg composites tends to be higher than the (B₄C+Ti)/AZ91D composites fabricated under the same conditions, which means that (B₄C+Ti)/AZ91D composites exhibit better wear resistance. The main reason is that AZ91D magnesium has better mechanical performance than pure magnesium, and after adding or producing the same kinds of reinforcement, the corresponding composites demonstrate better performance. Furthermore, with the extension of holding time, both kinds of composites have a lower wear rate, which means better wear resistance. As mentioned above, as the holding time is extended, the in situ reaction between B₄C and Ti is more sufficient; thus, the interface bonding is more solid and exhibits better wear resistance.

SEM micrographs that show the worn surfaces of the composites corresponding to Figure 4 are displayed in Figure 5. When the holding time was 30 min (Figure 5a), the worn surface of the (B₄C+Ti)/Mg composite is rough, and there are lots of pits, grooves and craters running parallel to the sliding direction, which is the characteristic of abrasive wear and adhesive wear [3]. These pits were formed when B₄C or Ti particles were pulled out of the matrix during wear and the formation procedure can be explained as follows: early wear and tear of the soft matrix phase make the rigid particles gradually protrude. After running for a period of time, the wear-resisting, hard particles tend to bear the main loads. The starting raw materials B₄C and Ti demonstrate a certain degree of in situ reaction, and the interface between the produced phases and the matrix has a certain binding force. Under the aegis of the soft matrix, hard particles do not easily wear off, thus leading to particle exposure and making the worn surface become rough. If the friction continues to increase, the rigid particles will be pulled out or will fall off from the soft matrix, leaving pits on the worn surface. Then, plowing or grooves will be caused by the friction of hard particles between the samples and counter discs.

Furthermore, it could be observed from Figure 5a that there were obvious craters within the worn surfaces. As adhesive wear occurs, a series of small crack appeared on the worn surface along the direction perpendicular to the wear and tear. Those cracks interact with each other or cause the peeling of layered or large materials at the intersection, thus forming a crater on the worn surface, which means that delamination wear occurs [21]. Figure 5b shows that the worn surface of the (B₄C+Ti)/AZ91D composite prepared under the same condition is more smooth, and the craters tend to be shallower and smaller, and there are also significantly fewer pits, which all indicate that the wear resistance of the (B₄C+Ti)/AZ91D composite is better than that of the (B₄C+Ti)/Mg composite.

By comparing the microstructure of the (B₄C+Ti)/Mg composites (Figure 5a,c,e) with (B₄C+Ti)/AZ91D composites (Figure 5b,d,f), it could be observed that with prolonged holding time, the craters on the worn surfaces of both kinds of composites tended to be shallower and smaller, and the worn surfaces more smooth, which illustrates that the wear resistance of both kinds of composites gradually improves. The main reason can be attributed to more sufficient in situ reactions between the starting materials, leading to stronger interface bonding. To sum up, the main wear mechanisms for both kinds of composites are abrasive and delamination wear, and the wear resistance of the (B₄C+Ti)/AZ91D composite is better than that of the (B₄C+Ti)/Mg composite prepared under the same condition, and the wear resistance of both kinds of composites increased with the extension of holding time.

3.3. Compression and Bending Performance

Figure 6 shows the relationship between the compressive and bending strength of the (B₄C+Ti)/Mg and (B₄C+Ti)/AZ91D composites and the holding time. It could be observed that the (B₄C+Ti)/AZ91D composites exhibit higher compressive strength and bending strength than that of the (B₄C+Ti)/Mg composites. When the holding time was 180 min, the compressive and bending strength of the (B₄C+Ti)/Mg composite was 370 MPa and 263 MPa, respectively, while that of the (B₄C+Ti)/AZ91D composite was 550 MPa and 480 MPa, respectively, nearly 1.5 and 1.8 times higher than the (B₄C+Ti)/Mg composites. It could also be observed from Figure 6 that the compressive and bending strength of the (B₄C+Ti)/Mg and (B₄C+Ti)/AZ91D composites increased with the extension of holding time, which means that the compressive and bending properties gradually improved. This is mainly because magnesium melt can infiltrate into the preform more sufficiently, and the in situ reaction between the starting materials was also more complete and sufficient; therefore, the interface between the reinforcing phase and the matrix became stronger.

The bending fracture surfaces of (B₄C+Ti)/Mg and (B₄C+Ti)/AZ91D composites are shown in Figure 7. When the holding time was 30 min, the bending fracture surface of the (B₄C+Ti)/Mg composite was smooth and the composite showed an obvious brittle fracture on a macroscopic scale. It could be observed from Figure 7a that the outline of the reinforcing particles B₄C and Ti with irregular shapes was clear. Brittle fractures of the reinforcing particles also exist and the main reason is that the particles tended to bear the main loads. Furthermore, there were little dimples in the fracture surface, and parts of the reinforcing phases were pulled out of the matrix. The interface bonding between the reinforcing phases and matrix was not strong enough, so some particulates were pulled out under high stress.

Figure 7b shows that the outlines of the reinforcing particles in (B₄C+Ti)/AZ91D were not clear and parts of the reinforcing particles presented transcrystalline fractures, and also had no obvious cracks. There were some voids in the fracture surface, but the metal tearing edge in the fracture surface significantly increased, which has obvious characteristics of a ductile fracture. Therefore, the bending performance of the (B₄C+Ti)/AZ91D composites was better than that of the (B₄C+Ti)/Mg composites. The cause of metal tearing edge can be explained as follows: when the composite was subjected to a bending stress, the AZ91D magnesium alloy matrix underwent plastic deformation firstly, and the dislocation moved within the matrix. When the dislocation moved to the vicinity of the reinforcing particles, the dislocation was blocked by the reinforcing particles and piled up here. As the stress continued to increase, the cracks nucleated and grew at the poorly bonded interface. If the interfacial strength of the dislocation pile-up is high, cracks will preferentially spread within the magnesium matrix, and the magnesium matrix will be torn, thus forming tearing edges.

From Figure 7c,e, it can be observed that the size of reinforcing particles within the (B₄C+Ti)/Mg composite gradually decreased, but the amount of metal tearing edges tended to increase and even shallower dimples appeared. No brittle fractures of the reinforcing particles occurred, and the cracks at the interface were greatly reduced. It can be explained by the fact that as the holding time prolongs, the in situ reaction between the starting materials became more sufficient and the interface bonding was stronger, so the cracks preferentially originated within the matrix. It can be deduced that the extension of holding time is beneficial for the improvement of the bending performance. By comparing Figure 7b,d,f, similar results can be obtained that the size of reinforcing particles within the composites tends to decrease, and no cracks within the fracture surface exist, and the amount of dimples gradually increased. Furthermore, the metal tearing edges and dimples within the fracture surface of the (B₄C+Ti)/AZ91D composites were more, leading to much better bending performance than (B₄C+Ti)/Mg composites prepared under the same conditions.

Figure 8 shows the true stress–strain curve of the (B₄C+Ti)/Mg and (B₄C+Ti)/AZ91D composites. It could be observed that the compression strength of the (B₄C+Ti)/AZ91D composite was higher than that of the (B₄C+Ti)/Mg composite, and the fracture strain value was relatively higher. Figure 9 shows the SEM micrographs of the compression fracture surfaces of (B₄C+Ti)/Mg and (B₄C+Ti)/AZ91D composites corresponding to Figure 8. As shown in Figure 9, from a macro perspective, the fractures of the compression samples of both kinds of composites were all produced along the direction of 45° with the compression axis and had cutting-off characteristics, showing obvious brittle fracture characteristics. As observed from the microcosm, there were cracks perpendicular to the fracture surface in both composites, and the cracks ran through the interior of the fracture surface. However, the crack in the fracture of the (B₄C+Ti)/AZ91D composite was much finer, which means its compression property is better. This can be explained by the structure of both kinds of composites and the work hardening and softening caused by the thermal effect during compression deformation. On one hand, the in situ reaction of the starting materials in the (B₄C+Ti)/AZ91D composite tended to be more sufficient, and the reinforcement tended to be network interpenetration, thus leading to stronger interfacial bonding. Under the action of compressive shear stress, cracks were mainly initiated and propagated at the interface between the matrix and reinforcing phases. Under equivalent deformation, the deformation in the (B₄C+Ti)/AZ91D composite can be dispersed in more grains, thus leading to smaller difference in the strain degree in the grain and near grain boundary and more uniform deformation. Relatively speaking, the reduced stress concentration caused in the composite enables the material to withstand a larger amount of deformation before a fracture, thus obtaining a larger strain. On the other hand, the strength of AZ91D is higher than that of Mg, leading to a greater thermal effect during cold compression. The increase in flow stress caused by softening and work hardening in the (B₄C+Ti)/AZ91D composite is smaller than that in the (B₄C+Ti)/Mg composite. With the increase in deformation, the work hardening effect of the (B₄C+Ti)/Mg composites reaches the maximum, and brittle fractures occur. However, at this time, the thermal softening effect of the (B₄C+Ti)/AZ91D composite is weakened, and the work hardening process becomes more and more obvious, resulting in its final compression strength being higher than that of the (B₄C+Ti)/Mg composite.

4. Conclusions

From this work, the following conclusions can be drawn:

(1)

Light-weight (B₄C+Ti)/Mg and (B₄C+Ti)/AZ91D composites were successfully prepared by in situ reactive infiltration, and extending the holding time and magnesium matrix alloying are both beneficial for preparing composites with interpenetrating network structures. The main reason is that extending the holding time and magnesium matrix alloying both facilitate more sufficient and complete in situ reactions between the starting materials.

(2)

(B₄C+Ti)/AZ91D composites exhibit better wear performance than (B₄C+Ti)/Mg composites prepared under the same conditions, which can be attributed to the higher hardness of the (B₄C+Ti)/AZ91D composites. Ultimately, more sufficient and complete reactions between B₄C and Ti in the (B₄C+Ti)/AZ91D composites occurred. In addition, their dominant wear mechanisms are abrasive and delamination wear.

(3)

The compression and bending performance of the (B₄C+Ti)/AZ91D composites were better than that of the (B₄C+Ti)/Mg composites, and the extension of holding time leads to an improvement in both kinds of properties for the composites.