Next Article in Journal
On the Mechanical Behavior of LP-DED C103 Thin-Wall Structures
Previous Article in Journal
Development and Characterisation of a New Die-Casting Die Cooling System Based on Internal Spray Cooling
 
 
Search for Articles:
Title / Keyword
Author / Affiliation / Email
Journal
Article Type
 
 
Advanced Search
 
Section
Special Issue
Volume
Issue
Number
Page
 
Journals
Metals
Volume 14
Issue 9
10.3390/met14090957
metals-logo
Submit to this Journal Review for this Journal Propose a Special Issue
► Article Menu

Article Menu

share Share announcement Help format_quote Cite thumb_up
...
Endorse textsms
...
Comment
first_page
settings
Order Article Reprints
Font Type:
Arial Georgia Verdana
Font Size:
Aa Aa Aa
Line Spacing:
Column Width:
Background:
Article

Improving the Mechanical Properties of Al-Si Composites through the Synergistic Strengthening of TiB2 Particles and BN Nanosheets

by
Yiren Wang
1,
Jian Wang
2,
Zunyan Xu
2,*,
Baoqiang Xu
1,
Bingheng Yu
2,
Jianwu Dong
2 and
Caiju Li
2,*
1
Faculty of Metallurgical and Energy Engineering, Kunming University of Science and Technology, Kunming 650093, China
2
Faculty of Materials Science and Engineering, Kunming University of Science and Technology, Kunming 650093, China
*
Authors to whom correspondence should be addressed.
Metals 2024, 14(9), 957; https://doi.org/10.3390/met14090957 (registering DOI)
Submission received: 22 July 2024 / Revised: 12 August 2024 / Accepted: 16 August 2024 / Published: 23 August 2024
Browse Figures
Versions Notes

Abstract

:
The size and distribution of the silicon phase and intermetallic phase are important factors affecting the properties of Al11Si3Cu2NiMg alloy (M142). In this study, BNNS and micro-TiB2 were used to synergistically refine and reinforce M142 composites (M142-BNNS-TiB2). After T6 heat treatment, the comprehensive mechanical properties of M142-BN-TiB2 composites were excellent, with an ultimate tensile strength of 463 MPa and an elongation of 2.6%. In addition, the introduction of BNNS and micro-TiB2 changed the fracture mode of M142 from brittle fracture to quasi-cleavage fracture, and the introduction of BNNS and micro-TiB2 refined the Si phase and intermetallic phase, which could change the origin of the crack in the composite, thus improving the ductility of the composite.

1. Introduction

Aluminum silicon (Al-Si) alloys are widely used in the automotive industry due to their high strength, low coefficient of expansion, low density, and high thermal conductivity [1,2,3,4,5,6]. In order to obtain better performance, elements such as Cu [2,7], Mg [8], and Ni [9] are usually added to Al-Si alloy for alloying. Meanwhile, Al-Si alloys cannot be subjected to precipitation hardening through artificial aging, and by adding alloying elements such as Cu and Si, aluminum alloys series 2xxx (with Cu) and 6xxx (with Si) are formed, which can be subjected to artificial aging and thus enhance their properties due to precipitation hardening. The microstructure of Al-Si series alloys mainly includes the silicon phase and intermetallic phase [10]. The research shows that in the deformation process of the Al-Si series alloy, the cracks will first appear in the silicon phase and the intermetallic phase [11] because in the deformation process, the load will be transferred from the matrix to the silicon phase and the intermetallic phase, and the silicon phase and the intermetallic phase belong to the brittle phase, easy to produce cracks [12,13,14,15]. The number of crack phases increases with increasing strain, and the ultimate failure of the alloy is caused by the joint of a large number of cracks [16]. Therefore, the size and distribution of the silicon phase and intermetallic phase are important factors affecting the performance of Al-Si alloy.
Casting and powder metallurgy are common methods for preparing Al-Si alloy and its composites [17,18,19], among which powder metallurgy (PM) is considered a promising method for manufacturing Al-Si alloy, which can produce Al-Si alloy with fine grains while maintaining small sizes of silicon particles and metal phases, and the resulting Al-Si alloy has suitable strength and ductility [20]. In this way, Cai et al. [21] prepared Al-22Si alloy with a tensile strength of 148 MPa and elongation of 5% by hot pressing. Kai et al. [22] prepared Al10SiMgFe alloy with a tensile strength of 291 MPa and elongation of 17% by sintering, hot deformation (hot rolling), and T6 heat treatment. In recent years, the Al-Si composite prepared by powder metallurgy has attracted much attention for its high strength, low thermal expansion coefficient, and suitable wear resistance [23]. TiB2 [24,25], SiC [26,27], TiC [28], CNTs [29], B4C, h-BN, and other hard ceramic particles are often used as reinforcements for aluminum matrix composite. Among them, TiB2 is of concern because of its suitable wettability with the aluminum matrix. Xi et al. [30] prepared a TiB2/Al-12Si composite with yield strength up to 247 MPa and showed that there was a strong interfacial bond between Al-12Si and TiB2. In addition, h-BN has the advantages of low density [31], high melting point, high hardness [32], small thermal expansion coefficient, and high thermal conductivity, which makes it an ideal reinforced material [33]. Shakti Corthay et al. [34] prepared Al-2% h-BN composites with a strength of 405 MPa and maintained high ductility. However, according to the research, a large amount of cracks in Al-Si composite prepared by powder metallurgy are preferentially formed near the silicon phase and intermetallic phase, resulting in material fracture failure, and the reinforcing body cannot give full play to its reinforcement effect [23,35]. Therefore, it is a challenging task to achieve the refinement of silicon particles and intermetallic phase in Al-Si composite, and at the same time give full play to the reinforcement effect of reinforcement, so as to prepare high-performance Al-Si composite.
The size of Si and intermetallic compounds in Al-Si alloys can significantly affect their mechanical properties. In this study, Al11Si3Cu2NiMg (M142) powder was used as raw material, and TiB2 and BN nanosheets were used as reinforcers. The addition of hybrid reinforcers inhibited the cold welding of metal powders during ball milling, and the purpose of refining the Si phase and intermetallic phase was achieved. Then, high-performance M142-BNNS-TiB2 composites were prepared using powder metallurgy and a hot extrusion process. The effects of BNNs and TiB2 on the microstructure of M142-BNNS-TiB2 composites were studied systematically. In addition, the relationship between microstructure and properties of M142-BNNS-TiB2 composites was studied, and its rule was revealed.

2. Materials and Methods

Figure 1 schematically shows the preparation process of the M142-BNNS-TiB2 composites. Materials used in this work include Al11Si3Cu2NiMg (M142) powders (spherical, particle size 30–52 μm), nanoscale lamellar h-BN (BNNS, diameter 1–5 μm, thickness 5 nm), and micron-scale TiB2 (particle size 1–5 μm). First, the M142 powder, BNNS (2 wt%), and TiB2 (0.5 wt%) were mechanically ball milled for 10 h under an argon atmosphere (PM0.4-A) (ball material ratio 10:1, ball milling speed 200 r/min) to obtain M142-BNNS-TiB2 composite powder, and then the obtained M142-BNNS-TiB2 composite powder was cold-pressed into a cylinder with a diameter of 30 mm using a single-arm hydraulic press (W41-100T) (pressure: 600 Mpa, pressure maintained for 30 min). The densification process involved heating the cold-pressed composite to 750 °C for 90 min and subsequently hot-extruding at 450 °C (extrusion ratio: 25:1) to obtain the M142-BNNS-TiB2 composite rod. Finally, the M142-BNNS-TiB2 composite rod was subjected to T6 heat treatment (solid solution at 510 °C for 3 h, then water quenching, artificial aging at 180 °C for 6 h). For a comparative study, a control group was prepared in this paper, and the same content of nano BNNS (2 wt%) was added to the M142 powder. The composite was prepared using the same process as described above, and the prepared composite was named M142-BNNS. The microstructure of the samples was analyzed using field emission scanning electron microscopy (FE-SEM, Nova Nano 450) and high-resolution transmission electron microscopy (HRTEM, FEI Tecnai F30). The tensile test specimens (scale length of 5 mm, thickness of 1.7 mm) were tested at room temperature at a constant strain rate of 5 × 10−4 s−1 in a general tensile testing machine (AG Xplus 50 kN) equipped with an automatic contact tester.

3. Results and Discussion

3.1. Microscopic Morphology of Ball Milled Composites

Figure 2a–c show the surface morphologies of pure M142 powder and M142-BNNS and M142-BNNS-TiB2 composite powder after ball milling. Figure 2a shows that spherical M142 powder becomes larger irregular M142 powder due to repeated impact, fracture, and cold welding [36] in the ball milling process. Figure 2b shows M142-BNNS composite powder after ball milling. BNNS nanosheets are uniformly dispersed on the surface of the powder, no cold welding of M142 powder occurs, and part of the spherical M142 powder is flaky after ball milling. This is because the BNNS uniformly dispersed on the surface of the powder has a suitable lubrication effect, which prevents the cold welding of the M142 powder in the ball milling process. Figure 2c shows that M142-BNNS-TiB2 composite powder after ball milling also does not occur during cold welding, and part of the spherical M142 powder is flaky after ball milling. In addition, the average particle diameters of M142-BNNS and M142-BNNS-TiB2 are 33.97 μm and 19.59 μm, respectively, which shows that the addition of TiB2 can further refine the powder. This is because the added hard ceramic particles TiB2 continuously impact and wear M142 articles in the ball milling process, which makes the size of M142 powder decrease. Finally, the introduction of BNNS and micron-level TiB2 particles can realize the flaking and particle size refinement of the powder in the ball milling process.

3.2. Composition and Distribution of Intermetallic Phases in Composites

Figure 3(a1,b1,c1), respectively, show the micro-SEM images of hot extruded M142, M142-BNNS, and M142-BNNS-TiB2. It can be observed that the second phase in the matrix is distributed along the extrusion direction, which is caused by the shear stress in the hot extrusion process. Figure 3(a2,b2,c2), respectively, show the morphology and distribution of the second phases of hot extruded M142, M142-BNNS, and M142-BNNS-TiB2. Figure 3(a2) shows that the second phases mainly exist in M142 alloy, including granular eutectic Si, irregular lumpy Q-Al5Cu2Mg8Si6, and δ-Al3CuNi (confirmed by the point scanning results of EDAX in Figure 3d). In Figure 3(b2), it can be observed that the addition of BNNS leads to the irregular blocky δ-Al3CuNi being broken, resulting in the formation of many fine granular δ-Al3CuNi. In Figure 3(c2), it can be observed that the micron-level TiB2 is uniformly dispersed in the Al matrix. In addition, the addition of BNNS and TiB2 further smashes and refines δ-Al3CuNi. In order to quantitatively analyze the synergistic effect of BNNS and TiB2 on the refinement of Al3CuNi, the grain size statistics of Al3CuNi in M142, M142-BNNS, and M142-BNNS-TiB2 are conducted, as shown in Figure 3e–g. The results show that the introduction of BNNS can refine the size of Al3CuNi in M142 from 14.29 μm to 3.05 μm, and the simultaneous introduction of BNNS and TiB2 further refines the average particle size of Al3CuNi to 1.84 μm. This is because the simultaneous introduction of BNNS and micron-level TiB2 particles in the ball milling process realizes the lamellarization and particle size refinement of the powder (Figure 2c), thereby refining the Al3CuNi phase in the composite.
Figure 4 shows SEM images of M142 and composites after T6 heat treatment and statistical graphs of the average size of Al3CuNi. Figure 4a is the SEM image of M142 alloy after T6 heat treatment. the second phase in M142 alloy after T6 heat treatment is still Si, Q-Al5Cu2Mg8Si6 and δ-Al3CuNi. After T6 heat treatment, Al5Cu2Mg8Si6 exhibited significant morphological changes, changing from irregular blocky Q-Al5Cu2Mg8Si6 to skeletal, which was dispersed at the grain boundary. Figure 4b,c show that T6 heat treatment does not change the morphology of Q-Al5Cu2Mg8Si6 in M142-BNNS and M142-BNNS-TiB2. In addition, the average size of Al3CuNi in M142, M142-BNNS, and M142-BNNS-TiB2 after T6 heat treatment is statistically analyzed, as shown in Figure 4d–f. By comparing the average size of Al3CuNi in a hot extrusion state and M142, M142-BNNS, and M142-BNNS-TiB2 after T6 heat treatment (as shown in Table 1), it is found that Al3CuNi in composite material after T6 heat treatment is coarsened. This is because the temperature increases in the heat treatment process, providing a driving force for the diffusion of Cu and Ni atoms, which makes Al3CuNi coarsened [37].

3.3. Morphology and Distribution of Si Phase in Composites

The mechanical properties of Al-Si composites are known to be largely dependent on the Si phase characteristics (size, shape, and distribution). Figure 5 shows the BSE image of the hot extruded M142 and composites, the EDS spectrum of Si element, and the grain size statistical diagram of the corresponding Si phase. Through the observation of Figure 5(a2), Si in M142 mainly exists in two forms. One is a large lump of Si particles that are evenly dispersed in the aluminum matrix, with grain size between 30 and 50 μm, and the other is a small Si particle cluster area, the grain size of small Si particles is between 5 and 15 μm. It can also be seen from the Si grain statistics of M142 in Figure 5(a3) that the Si grain distribution presents a bimodal grain distribution. This is because the ball milling process passed M142 cold welding (Figure 2a), resulting in the inability to refine some Si particles. Figure 5(b2) shows that the morphology of Si particles changes after BNNS is added. The Si particles are mainly divided into two morphologies: one is small Si particles, which are evenly dispersed in the aluminum matrix, with a grain size between 10 and 15 μm; the other is a strip region of small Si particle clusters. In addition, it can be concluded from Figure 5(b3) that the average grain size of Si particles in M142-BNNS is 17.66 μm. According to Figure 5(c2), Si particles in M142-BNNS-TiB2 mainly exist in two forms, which are similar to the forms of silicon particles in M142-BNNS. The difference is that there are fewer small Si particles clustered in strip areas in M142-BNNS-TiB2, and most of the small Si particles present a streamlined distribution along the direction of hot extrusion, the average grain size of Si particles is 14.23 μm (Figure 5(c3)). The reasons for this phenomenon mainly include two aspects: (1) the addition of BNNS nanosheets and micron-level TiB2 particles in the ball milling process realizes the lamellarization and particle size refinement of the powder, which makes the Si particles in the powder damaged and refined, (2) the diffusion and precipitation of Si during sintering and hot extrusion lead to the presence of a large number of fine Si particles along the extrusion direction [22].
Figure 6 shows the BSE images of M142 and composites after T6 heat treatment, the EDS spectrum of Si element, and the statistical graph of grain size of the corresponding Si phase. According to the observation in Figure 6(a2), Si particles in M142 after T6 heat treatment still mainly exist in two forms, i.e., large massive Si particles and aggregation area of small Si particles. However, it is worth noting that compared with the unheat-treated M142 (Figure 5(a2)), it can be found that the proportion of large Si particles in M142 after T6 heat treatment increases, which can also be observed by the statistical figure of Si particle size (Figure 6(a3)). Figure 6(b2) shows that after T6 heat treatment, Si particles in M142-BNNS mainly exist in two forms: small Si particles and small Si particle cluster regions. It is worth noting that, compared with the unheat-treated M142-BNNS (Figure 5(b2)), the strip region of small Si particle clusters disappears and becomes small Si particle cluster regions, which are more evenly dispersed. Figure 6(c2) shows that after T6 heat treatment, the cluster area of small Si particles in M142-BNNS-TiB2 decreases, and a large number of small Si particles are evenly distributed in the Al matrix along the direction of hot extrusion. This is due to the fact that the T6 heat treatment promotes the further recrystallization and uniform precipitation of Si particles, resulting in a fine-grained microstructure.
To analyze the effect of T6 heat treatment on Si size, the grain size of Si in the composites in the hot extrusion state and T6 heat treatment state was compared (as shown in Table 2). It can be seen that the average grain size of Si in M142, M142-BNNS, and M142-BNNS-TiB2 composites increased by 18.40%, 5.94%, and 2.81% year-on-year after T6 heat treatment. It is noteworthy that the coarsening of Si in M142 is the most significant after T6 heat treatment because the difference in Si particle size provides the premise for Ostwald ripening [4]. The Si grains in M142 are distributed in bimodal grain size, and the difference in Si particle size is large, which makes the Si particles in M142 more prone to Ostwald ripening. The size difference of Si particles in M142-BNNS and M142-BNNS-TiB2 composites is small, and the Ostwald ripening effect is weakened.

3.4. Microstructure of M142-BNNS-TiB2 Composite

Figure 7 is the TEM microanalysis of the M142-BNNS-TiB2 composite after T6 heat treatment. Figure 7a is the low-magnification TEM image showing the distribution of elliptical and hexagon particles along the grain boundary (GBS) of M142-BNNS-TiB2 composites, which are interface products. Figure 7b,c are the electron diffraction patterns of the hexagon grains and elliptical grains at the grain boundary, respectively. Through the diffraction pattern, it can be determined that both the hexagon grains and elliptical grains are Q-Al5Cu2Mg8Si6. At the same time, it can be seen from Figure 7a that there are a large number of dislocations in the perimeter of hexagonal particle Q-Al5Cu2Mg8Si6, which proves that Q-Al5Cu2Mg8Si6 has a suitable dislocation-strengthening effect. Such submicron particles distributed along the grain boundary can store more dislocations in the strain process, improving the mechanical properties of the composites [26]. Figure 7d shows the presence of high-density granular and acicular nanoprecipitates in the M142-BNNS-TiB2 composite. Figure 7e shows the HRTEM of the particles and needle-like nanoprecipitate phase in Figure 7d. The HRTEM, fast Fourier transform (FFT), and inverse fast Fourier transform (IFFT) (Figure 7e–g) results demonstrate that these intragranular particles nanoprecipitate are Q’-Al5Cu2Mg8Si6, needle-like nanoprecipitate phase are θ’-Al2Cu. The high density of θ’ and Q’ nanoprecipitates is beneficial for enhancing the strength of composites [28]. In conclusion, the introduction of BNNS and micron-scale TiB2 allows for the refinement and uniform dispersion of Q-Al5Cu2Mg8Si6.
To further analyze the microstructure in M142-BNNS-TiB2, TEM observation was conducted on Si, BNNS and TiB2 in the composite, and the results are shown in Figure 8. Figure 8a shows the TEM image of ellipsoidal Si particles in the M142-BNNS-TiB2 composite. The ellipsoidal Si is not prone to stress concentration in the strain process, which is conducive to improving the strength and plasticity of the composites. In addition, there are a large number of dislocations around the Si particles. Figure 8b BNNS (marked by yellow arrow) is uniformly distributed in the grain boundary. There is a thermal mismatch between BNNSs and Al, and the geometric necessary dislocation caused by thermal mismatch also promotes the strengthening behavior of the composite [32]. Figure 8c shows that there are dislocations around the TiB2 particle, which proves that TiB2 has a suitable dislocation-strengthening effect. Figure 8d is the HRTEM of the marked area in Figure 8c, and Al and TiB2 have a tight interface, indicating that Al and TiB2 have a suitable interface combination, which is conducive to the load transfer of TiB2.

3.5. Mechanical Properties of Composites

Figure 9a shows a comparison of typical engineering stress–strain curves for pure M142, M142-BNNS, and M142-BNNS-TiB2 composites after hot extrusion. Specifically, the ultimate tensile strength (UTS) of the M142-BNNS-TiB2 composite was 313 ± 5 MPa, an increase of 22.2% compared with the value of the pure M142 sample. Compared with the M142-BNNS composite, the M142-BNNS-TiB2 composite also showed an improvement in UTS while maintaining a comparable elongation.
Figure 1 shows the comparison of typical engineering stress–strain curves of pure M142, M142-BNNS, and M142-BNNS-TiB2 composites treated with T6 heat treatment. As shown in Table 3, the UTS of M142 after T6 heat treatment is 362 ± 4 MPa, 41.4% higher than that of M142 without T6 heat treatment, which is due to the precipitation strengthening introduced by T6 heat treatment (Figure 7). The UTS and elongation (EI) of M142-BNNS-TiB2 after T6 heat treatment are 463 ± 4 MPa and 2.6 ± 0.2%, respectively, which are increased by 22.2% and 4% compared with pure M142 after T6 heat treatment, and by 17.2% and 52.9% compared with M142-BNNS composite after T6 heat treatment, shows a superior strength–ductility balance. Figure 4 and Figure 6 show that the addition of BNNS and TiB2 makes the intermetallic compound phase and Si phase in M142-BNNS-TiB2 fine and uniformly dispersed, and the fine and uniformly dispersed second phase is not easy to crack when subjected to external loading and can store more dislocations, which is conducive to improving the mechanical properties of M142-BNNS-TiB2. In addition, as shown in Figure 7, the addition of BNNS and TiB2 makes the composite material produce uniformly dispersed submicron Q-Al5Cu2Mg8Si6 at the grain boundary, which is conducive to improving the strength of the composite material. At the same time, after T6 heat treatment, the high-density θ’ and Q’ nano precipitation phase formed in the grain is also conducive to improving the strength of the composite.

3.6. Fracture Morphology of Composite

Figure 10(a1) shows that M142 exhibits typical brittle fracture after T6 heat treatment, with a large-size cleavage surface and a typical river-like cleavage surface. Fractured Si and Q-Al5Cu2Mg8Si6 were observed on the fracture surface of M142 in Figure 10(a2). In addition, there were cracks on the fractured eutectic Si, which penetrated through the whole eutectic Si, indicating that the eutectic Si was prone to fracture when subjected to external loads. Figure 10(b1) shows that M142-BNNS-TiB2 after T6 heat treatment exhibits typical quasi-cleavage fracture characteristics, and there are typical small-sized cleavage surfaces and tear edges. Figure 10(b2) shows that the M142-BNNS-TiB2 fracture surface is distributed with smaller sizes of eutectic Si, Q-Al5Cu2Mg8Si6, and δ-Al3CuNi. These findings confirm that the introduction of BNNS and TiB2 refines the Si phase and intermetallic phase, which can change the location of crack initiation in the composite, thereby improving the ductility of the composite.

4. Conclusions

We prepared M142 composites with synergistic refinement and reinforcement by BNNS and micron TiB2, studied the effects of BNNS and micron TiB2 on eutectic Si and intermetallic phases, and obtained the following conclusions:
  • BNNS and micron TiB2 particles were introduced simultaneously in the ball milling process, which can realize the lamellarization and particle size refinement of M142 powder;
  • The introduction of BNNS and micron TiB2 can realize the uniform refinement of the Si phase and intermetallic phase in M142-BNNS-TiB2;
  • After T6 heat treatment, the M142-BN-TiB2 composite exhibits excellent comprehensive mechanical properties, with a UTS of 463 MPa (Compared with M142 and M142-BNNS, it has increased by 101 and 68 MPa, respectively) and an elongation of 2.6% (Compared with M142 and M142-BNNS, it has increased by 0.1% and 0.9%, respectively);
  • The introduction of BNNS and micro-TiB2 changed the M142 fracture mode from brittle fracture to quasi-cleavage fracture. At the same time, the introduction of BNNS and micro-TiB2 to refine the Si phase and intermetallic phase can change the location of crack initiation in the composite, thus improving the ductility of the composite.

Author Contributions

Conceptualization, Y.W., J.W., Z.X. and C.L.; investigation, Y.W., J.W., B.X. and B.Y.; methodology, Y.W., J.W., B.X. and B.Y.; supervision, Z.X. and C.L.; writing—original draft, Y.W. and J.W.; writing—review and editing, Z.X. and J.D. All authors have read and agreed to the published version of the manuscript.

Funding

This work was supported by Yunnan Major Scientific and Technological Projects [grant no. 202202AG050011, 202202AB080004], the National Natural Science Foundation of China [grant no. 52061021], and the Yunnan Industrial Technology Innovation Talent Project.

Data Availability Statement

The original contributions presented in the study are included in the article, further inquiries can be directed to the corresponding authors.

Conflicts of Interest

The authors declare no conflicts of interest.

References

  1. Alemdag, Y.; Beder, M. Effects of zinc content on strength and wear performance of Al−12Si−3Cu based alloy. Trans. Nonferrous Met. Soc. China 2019, 29, 2463–2471. [Google Scholar] [CrossRef]
  2. Asghar, Z.; Requena, G.; Zahid, G.H.; Rafiud, D. Effect of thermally stable Cu- and Mg-rich aluminides on the high temperature strength of an AlSi12CuMgNi alloy. Mater. Charact. 2014, 88, 80–85. [Google Scholar] [CrossRef]
  3. Beder, M.; Alemdag, Y. Influence of Mg addition and T6 heat treatment on microstructure, mechanical and tribological properties of Al–12Si–3Cu based alloy. Trans. Nonferrous Met. Soc. China 2021, 31, 2208–2219. [Google Scholar] [CrossRef]
  4. Ding, C.; Hao, H.; Lu, Z.; Yu, C.; Wu, X.; Yu, P.; Ye, S. Fabrication of hypereutectic Al–Si alloy with improved mechanical and thermal properties by hot extrusion. Mater. Charact. 2023, 202, 113026. [Google Scholar] [CrossRef]
  5. Dong, J.; Jiang, J.; Wang, Y.; Huang, M.; Cui, J.; Qin, T.; Kong, L. Effect of casting process and Zr, V addition on microstructure and mechanical properties of Al–12Si-4.5Cu–2Ni alloy. J. Mater. Res. Technol. 2024, 29, 1428–1440. [Google Scholar] [CrossRef]
  6. Li, D.; Zhao, K.; Han, M.; Liu, G.; Sun, Q.; Liu, S.; Liu, X. Optimizing microstructure and enhancing mechanical properties of Al–Si–Mg–Mn–based alloy by novel C-doped TiB2 particles. J. Mater. Res. Technol. 2023, 26, 9450–9466. [Google Scholar] [CrossRef]
  7. Yang, Y.; Yu, K.; Li, Y.; Zhao, D.; Liu, X. Evolution of nickel-rich phases in Al–Si–Cu–Ni–Mg piston alloys with different Cu additions. Mater. Des. 2012, 33, 220–225. [Google Scholar] [CrossRef]
  8. Yamamoto, K.; Takahashi, M.; Kamikubo, Y.; Sugiura, Y.; Iwasawa, S.; Nakata, T.; Kamado, S. Effect of Mg content on age-hardening response, tensile properties, and microstructures of a T5-treated thixo-cast hypoeutectic Al–Si alloy. Mater. Sci. Eng. A 2020, 798, 140089. [Google Scholar] [CrossRef]
  9. Ma, S.; Li, Y.; Kan, W.; Zhang, J.; Wang, M.; Wang, L.; Wu, Y.; Wang, H.; Chen, Z. Enhancement of grain refinement and heat resistance in TiB2-reinforced Al-Cu-Mg-Fe-Ni matrix composite additive manufactured by electron beam melting. J. Alloys Compd. 2022, 924, 166395. [Google Scholar] [CrossRef]
  10. Zhang, M.; Wang, D.; Nagaumi, H.; Wang, R.; Zhang, X.; Zhou, P.; Wu, F.; Zhang, B. Achieving excellent strength-ductility in Al–Si–Cu–Mg cast alloy via effective work hardening. Mater. Sci. Eng. A 2024, 889, 145840. [Google Scholar] [CrossRef]
  11. Liu, H.; Pang, J.; Wang, M.; Li, S.; Zhang, Z. The effect of thermal exposure on the microstructure and mechanical properties of multiphase AlSi12Cu4MgNi2 alloy. Mater. Charact. 2020, 159, 110032. [Google Scholar] [CrossRef]
  12. Liu, H.Q.; Pang, J.C.; Wang, M.; Li, S.X.; Zhang, Z.F. Effect of temperature on the mechanical properties of Al–Si–Cu–Mg–Ni–Ce alloy. Mater. Sci. Eng. A 2021, 824, 141762. [Google Scholar] [CrossRef]
  13. Sui, Y.; Wang, Q.; Wang, G.; Liu, T. Effects of Sr content on the microstructure and mechanical properties of cast Al–12Si–4Cu–2Ni–0.8Mg alloys. J. Alloys Compd. 2015, 622, 572–579. [Google Scholar] [CrossRef]
  14. Tan, P.; Yang, Y.; Sui, Y.; Wang, Q.; Jiang, Y. The influence of Al–10Sr or/ and Al–5Ti–1B on microstructure and mechanical properties of Al–12Si–4Cu–2Ni–0.8 Mg alloys. J. Alloys Compd. 2019, 809, 151856. [Google Scholar] [CrossRef]
  15. Zhang, N.; Feng, Y.; Zhao, S.; Fu, Y.; Wang, L.; Guo, E. Microstructure evolution and mechanical properties of Al-12Si-xCu-yNi-1Mg alloy with different Ni/Cu ratios. Mater. Today Commun. 2024, 38, 108081. [Google Scholar] [CrossRef]
  16. Shin, S.-S.; Kim, E.-S.; Yeom, G.-Y.; Lee, J.-C. Modification effect of Sr on the microstructures and mechanical properties of Al–10.5Si–2.0Cu recycled alloy for die casting. Mater. Sci. Eng. A 2012, 532, 151–157. [Google Scholar] [CrossRef]
  17. Zasadzińska, M.; Strzpek, P.; Mamala, A.; Noga, P. Reinforcement of Aluminium-Matrix Composites with Glass Fibre by Metallurgical Synthesis. Materials 2020, 13, 5441. [Google Scholar] [CrossRef] [PubMed]
  18. Gupta, M.K. Mechanical behaviors of Al 6063/TiB2 composites fabricated by stir casting process. Mater. Today Proc. 2023, 82, 222–226. [Google Scholar] [CrossRef]
  19. Borodianskiy, K.; Selivorstov, V.; Dotsenko, Y.; Zinigrad, M. Effect of Additions of Ceramic Nanoparticles and Gas-Dynamic Treatment on Al Casting Alloys. Metals 2015, 5, 2277–2288. [Google Scholar] [CrossRef]
  20. Zhao, X.; Meng, J.; Zhang, C.; Wei, W.; Wu, F.; Zhang, G. A novel method for improving the microstructure and the properties of Al-Si-Cu alloys prepared using rapid solidification/powder metallurgy. Mater. Today Commun. 2023, 35, 105802. [Google Scholar] [CrossRef]
  21. Cai, Z.; Zhang, C.; Wang, R.; Peng, C.; Qiu, K.; Feng, Y. Preparation of Al–Si alloys by a rapid solidification and powder metallurgy route. Mater. Des. 2015, 87, 996–1002. [Google Scholar] [CrossRef]
  22. Chang, K.-C.; Hung, F.-Y.; Zhao, J.-R.; Miu, C.-F. Advancements in microstructure, texture characterization, and mechanical properties of powder metallurgy Al-Si-Mg-Fe alloy for engineering applications. J. Alloys Compd. 2024, 973, 172941. [Google Scholar] [CrossRef]
  23. Liu, Y.-X.; Wang, R.-C.; Peng, C.-Q.; Cai, Z.-Y.; Zhou, Z.-H.; Li, X.-G.; Cao, X.-Y. Microstructures and mechanical properties of in-situ TiB2/Al−xSi−0.3Mg composites. Trans. Nonferrous Met. Soc. China 2021, 31, 331–344. [Google Scholar] [CrossRef]
  24. Akbari, M.K.; Shirvanimoghaddam, K.; Hai, Z.; Zhuiykov, S.; Khayyam, H. Al-TiB2 micro/nanocomposites: Particle capture investigations, strengthening mechanisms and mathematical modelling of mechanical properties. Mater. Sci. Eng. A 2017, 682, 98–106. [Google Scholar] [CrossRef]
  25. Chen, C.; Sun, C.; Wang, W.; Qi, M.; Han, W.; Li, Y.; Liu, X.; Yang, F.; Guo, L.; Guo, Z. Microstructure and mechanical properties of in-situ TiB2/AlSi7Mg composite via powder metallurgy and hot extrusion. J. Mater. Res. Technol. 2022, 19, 1282–1292. [Google Scholar] [CrossRef]
  26. Liu, K.; Su, Y.; Wang, X.; Cai, Y.; Cao, H.; Ouyang, Q.; Zhang, D. Achieving simultaneous enhancement of strength and ductility in Al matrix composites by employing the synergetic strengthening effect of micro- and nano-SiCps. Compos. Part B Eng. 2023, 248, 110350. [Google Scholar] [CrossRef]
  27. Liu, G.; Zhou, S.; Wang, C.; Chen, S.; Sun, D.; Chen, A.; Li, Z.; Yan, C.; Shi, Y. Fabrication and characterization of SiC/Al composites prepared by laser powder bed fusion (LPBF) combined with vacuum pressure infiltration. Mater. Des. 2023, 236, 112495. [Google Scholar] [CrossRef]
  28. Gao, Y.-Y.; Qiu, F.; Zhao, Q.-L.; Jiang, Q.-C. A new approach for improving the elevated-temperature strength and ductility of Al–Cu–Mg–Si alloys with minor amounts of dual-phased submicron/nanosized TiB2–TiC particles. Mater. Sci. Eng. A 2019, 764, 138266. [Google Scholar] [CrossRef]
  29. Borodianskiy, K. Effect of carbon nanotubes’ addition on mechanical properties and thermal conductivity of copper. J. Mater. Sci. 2019, 54, 13767–13774. [Google Scholar] [CrossRef]
  30. Xi, L.X.; Zhang, H.; Wang, P.; Li, H.C.; Prashanth, K.G.; Lin, K.J.; Kaban, I.; Gu, D.D. Comparative investigation of microstructure, mechanical properties and strengthening mechanisms of Al-12Si/TiB2 fabricated by selective laser melting and hot pressing. Ceram. Int. 2018, 44, 17635–17642. [Google Scholar] [CrossRef]
  31. Ma, L.; Zhang, X.; Pu, B.; Zhao, D.; He, C.; Zhao, N. Achieving the strength-ductility balance of boron nitride nanosheets/Al composite by designing the synergistic transition interface and intragranular reinforcement distribution. Compos. Part B Eng. 2022, 246, 110243. [Google Scholar] [CrossRef]
  32. Ma, L.; Zhang, X.; Duan, Y.; Zhang, H.; Ma, N.; Zhu, L.; Rong, X.; Zhao, D.; He, C.; Zhao, N. Achieving exceptional high-temperature resistant Al matrix composites via two-dimensional BN pinning grain rotation. Compos. Part B Eng. 2023, 253, 110570. [Google Scholar] [CrossRef]
  33. Wang, Y.; Zhu, Y.; Li, R.; Liu, H.; Liu, W.; He, P. High-performance aluminum-based matrix composites reinforced with high content of boron nitride nanosheets. J. Alloys Compd. 2022, 906, 164358. [Google Scholar] [CrossRef]
  34. Corthay, S.; Kutzhanov, M.K.; Matveev, A.T.; Bondarev, A.V.; Leybo, D.V.; Shtansky, D.V. Nanopowder derived Al/h-BN composites with high strength and ductility. J. Alloys Compd. 2022, 912, 165199. [Google Scholar] [CrossRef]
  35. Xiao, Y.K.; Chen, H.; Bian, Z.Y.; Sun, T.T.; Ding, H.; Yang, Q.; Wu, Y.; Lian, Q.; Chen, Z.; Wang, H.W. Enhancing strength and ductility of AlSi10Mg fabricated by selective laser melting by TiB2 nanoparticles. J. Mater. Sci. Technol. 2022, 109, 254–266. [Google Scholar] [CrossRef]
  36. Yang, C.M.Y.; Li, X.; Li, C.J.; Peng, Y.Z.; Xing, Y.; Feng, Z.X.; Tan, J.; Tao, J.M.; Li, Z.L.; Wang, Y.R.; et al. Interface and strengthening mechanisms of Al matrix composites reinforced with in-situ CNTs grown on Ti particles. Mater. Des. 2023, 229, 111923. [Google Scholar] [CrossRef]
  37. Mozammil, S.; Karloopia, J.; Verma, R.; Jha, P.K. Effect of varying TiB2 reinforcement and its ageing behaviour on tensile and hardness properties of in-situ Al-4.5%Cu-xTiB2 composite. J. Alloys Compd. 2019, 793, 454–466. [Google Scholar] [CrossRef]
Metals 14 00957 g001
Figure 1. Schematically shows the preparation process of M142-BNNS-TiB2 composite.
Figure 1. Schematically shows the preparation process of M142-BNNS-TiB2 composite.
Metals 14 00957 g001
Metals 14 00957 g002
Figure 2. SEM images of pure M142 powder and composite powder after ball milling: (a) M142; (b) M142-BNNS; (c) M142-BNNS-TiB2.
Figure 2. SEM images of pure M142 powder and composite powder after ball milling: (a) M142; (b) M142-BNNS; (c) M142-BNNS-TiB2.
Metals 14 00957 g002
Metals 14 00957 g003
Figure 3. (a1,b1,c1) are SEM images of hot extruded M142, M142-BNNS, and M142-BNNS-TiB2, respectively; (a2,b2,c2) are locally enlarged images of (a1,b1,c1); (d) is the point scanning result of EDAX for the labeled phase; (e–g) are statistical graphs of the average size of Al3CuNi in M142, M142-BNNS, and M142-BNNS-TiB2, respectively.
Figure 3. (a1,b1,c1) are SEM images of hot extruded M142, M142-BNNS, and M142-BNNS-TiB2, respectively; (a2,b2,c2) are locally enlarged images of (a1,b1,c1); (d) is the point scanning result of EDAX for the labeled phase; (e–g) are statistical graphs of the average size of Al3CuNi in M142, M142-BNNS, and M142-BNNS-TiB2, respectively.
Metals 14 00957 g003
Metals 14 00957 g004
Figure 4. SEM images of M142 and composites after T6 heat treatment: (a) M142; (b) M142-BNNS; (c) M142-BNNS-TiB2; statistical chart of average size of Al3CuNi after heat treatment of M142 and composites: (d) M142: (e) M142-BNNS; (f) M142-BNNS-TiB2.
Figure 4. SEM images of M142 and composites after T6 heat treatment: (a) M142; (b) M142-BNNS; (c) M142-BNNS-TiB2; statistical chart of average size of Al3CuNi after heat treatment of M142 and composites: (d) M142: (e) M142-BNNS; (f) M142-BNNS-TiB2.
Metals 14 00957 g004
Metals 14 00957 g005
Figure 5. BSE images of hot extruded M142 and composites, EDS spectrum of Si element and grain size statistical diagram of corresponding Si phase: (a1–a3) M142; (b1–b3) M142-BNNS; (c1–c3) M142-BNNS-TiB2.
Figure 5. BSE images of hot extruded M142 and composites, EDS spectrum of Si element and grain size statistical diagram of corresponding Si phase: (a1–a3) M142; (b1–b3) M142-BNNS; (c1–c3) M142-BNNS-TiB2.
Metals 14 00957 g005
Metals 14 00957 g006
Figure 6. BSE images of M142 and composites after T6 heat treatment, the EDS spectrum of Si element, and the statistical graph of grain size of corresponding Si phase: (a1–a3) M142; (b1–b3) M142-BNNS; (c1–c3) M142-BNNS-TiB2.
Figure 6. BSE images of M142 and composites after T6 heat treatment, the EDS spectrum of Si element, and the statistical graph of grain size of corresponding Si phase: (a1–a3) M142; (b1–b3) M142-BNNS; (c1–c3) M142-BNNS-TiB2.
Metals 14 00957 g006
Metals 14 00957 g007
Figure 7. TEM micrographs of M142-BNNS-TiB2 composite after T6 heat treatment: (a) ellipsoidal and hexagonal Q-Al5Cu2Mg8Si6 distributed at the grain boundary; (b,c) diffraction patterns of hexagonal and ellipsoidal Q-Al5Cu2Mg8Si6 at the grain boundary, respectively; (d) θ′ and Q′ nano−precipitated phases inside the grains; (e) HRTEM of θ′ and Q′ nano-precipitated phases in Figure (d); (f) FFT and IFFT images of the θ′ nano−precipitated phases in Figure (e); (g) FFT and IFFT images of the Q′ nano-precipitated phases in Figure (e).
Figure 7. TEM micrographs of M142-BNNS-TiB2 composite after T6 heat treatment: (a) ellipsoidal and hexagonal Q-Al5Cu2Mg8Si6 distributed at the grain boundary; (b,c) diffraction patterns of hexagonal and ellipsoidal Q-Al5Cu2Mg8Si6 at the grain boundary, respectively; (d) θ′ and Q′ nano−precipitated phases inside the grains; (e) HRTEM of θ′ and Q′ nano-precipitated phases in Figure (d); (f) FFT and IFFT images of the θ′ nano−precipitated phases in Figure (e); (g) FFT and IFFT images of the Q′ nano-precipitated phases in Figure (e).
Metals 14 00957 g007
Metals 14 00957 g008
Figure 8. TEM micrographs of M142-BNNS-TiB2 composite after T6 heat treatment: (a) TEM images of Si in the composite are shown; (b) BNNS distributed in grain boundaries and grains; (c) TEM image shows a large number of dislocations around TiB2; (d) HRTEM of the interface between Al and TiB2.
Figure 8. TEM micrographs of M142-BNNS-TiB2 composite after T6 heat treatment: (a) TEM images of Si in the composite are shown; (b) BNNS distributed in grain boundaries and grains; (c) TEM image shows a large number of dislocations around TiB2; (d) HRTEM of the interface between Al and TiB2.
Metals 14 00957 g008
Metals 14 00957 g009
Figure 9. Room temperature tensile engineering stress–strain curves of pure M142, M142-BNNS, and M142-BNNS-TiB2 composites in different states: (a) hot extrusion; (b) T6 heat treatment.
Figure 9. Room temperature tensile engineering stress–strain curves of pure M142, M142-BNNS, and M142-BNNS-TiB2 composites in different states: (a) hot extrusion; (b) T6 heat treatment.
Metals 14 00957 g009
Metals 14 00957 g010
Figure 10. Typical fracture morphologies of pure M142 and composites after heat treatment: (a1,a2) M142; (b1,b2) M142-BNNS-TiB2.
Figure 10. Typical fracture morphologies of pure M142 and composites after heat treatment: (a1,a2) M142; (b1,b2) M142-BNNS-TiB2.
Metals 14 00957 g010
Table 1. Average size of Al3CuNi in M142, M142-BNNS, and M142-BNNS-TiB2 composites at different states (μm).
Table 1. Average size of Al3CuNi in M142, M142-BNNS, and M142-BNNS-TiB2 composites at different states (μm).
Hot ExtrusionT6 Heat Treatment
M14214.3915.22
M142-BNNS3.055.79
M142-BNNS-TiB21.842.28
Table 2. Average Si size of samples at different states (μm).
Table 2. Average Si size of samples at different states (μm).
Hot ExtrudedT6Size Increase (%)
M14223.5227.8518.40
M142-BNNS17.6618.715.94
M142-BNNS-TiB214.2314.632.81
Table 3. Mechanical properties of M142, M142-BNNS, and M142-BNNS-TiB2 composites in different states.
Table 3. Mechanical properties of M142, M142-BNNS, and M142-BNNS-TiB2 composites in different states.
Name of SampleUST (MPa)Elongation Rate (%)
Hot extrudedM142256 ± 65.7 ± 0.3
M142-BNNS301 ± 51.8 ± 0.3
M142-BNNS-TiB2313 ± 42.2 ± 0.4
T6 heat treatmentM142362 ± 42.5 ± 0.2
M142-BNNS395 ± 31.7 ± 0.1
M142-BNNS-TiB2463 ± 42.6 ± 0.2
Disclaimer/Publisher’s Note: The statements, opinions and data contained in all publications are solely those of the individual author(s) and contributor(s) and not of MDPI and/or the editor(s). MDPI and/or the editor(s) disclaim responsibility for any injury to people or property resulting from any ideas, methods, instructions or products referred to in the content.

Share and Cite

MDPI and ACS Style

Wang, Y.; Wang, J.; Xu, Z.; Xu, B.; Yu, B.; Dong, J.; Li, C. Improving the Mechanical Properties of Al-Si Composites through the Synergistic Strengthening of TiB2 Particles and BN Nanosheets. Metals 2024, 14, 957. https://doi.org/10.3390/met14090957

AMA Style

Wang Y, Wang J, Xu Z, Xu B, Yu B, Dong J, Li C. Improving the Mechanical Properties of Al-Si Composites through the Synergistic Strengthening of TiB2 Particles and BN Nanosheets. Metals. 2024; 14(9):957. https://doi.org/10.3390/met14090957

Chicago/Turabian Style

Wang, Yiren, Jian Wang, Zunyan Xu, Baoqiang Xu, Bingheng Yu, Jianwu Dong, and Caiju Li. 2024. "Improving the Mechanical Properties of Al-Si Composites through the Synergistic Strengthening of TiB2 Particles and BN Nanosheets" Metals 14, no. 9: 957. https://doi.org/10.3390/met14090957

Note that from the first issue of 2016, this journal uses article numbers instead of page numbers. See further details here.

Article Metrics

Back to TopTop