Analysis of the interactions between nonoxide reinforcements and Al–Si–Cu–Mg matrices

2.1.1 SiC in Al matrices

SiC in contact with liquid Al is expressed through the following reaction [74,78,79,80,81,82]:

Salvo et al. [82] established that the extension of this reaction depends on both the alloy composition and the processing conditions, mainly temperature and time of contact between liquid metal and the reinforcement. Al₄C₃ obtained through this reaction is brittle and insoluble and is present as either isolated precipitates or continuous layers [83]. When the Al₄C₃ layer is not present, SiC is partially bonded to the matrix, and decohesion could occur at the interface, propagating and growing cracks. This is observed in the transmission electron microscopy (TEM) and Figure 2a–d shows images of a 6061 Al–1.0 Mg–0.6 Si–0.25 Cu alloy reinforced with SiC particles. Figure 2a shows a clean and smooth interface, without the formation of Al₄C₃ or other interfacial products [84]. Nevertheless, some decohesion seems to be present (arrowed) in this interface, which (i) limits load transfer and (ii) leads to crack propagation. It is a desired perfect and complete contact between SiC and Al, as shown in Figure 2b, which corresponds to a high-resolution TEM (HRTEM) image of the SiC/Al interface and its selected area electron diffraction (SAED) pattern. Figure 2c shows the indexed pattern, while Figure 2d shows the inverse fast Fourier transform (IFFT) obtained from area 3 in Figure 2b, with an orientation relationship (0006) SiC || (111) Al. Other reported orientation relationship is (10–10) SiC || (−111) Al) [85]. Such semi-coherent interfaces can effectively transfer the load from matrix to SiC. Nevertheless, the zones without effective bonding could provoke failures in the composite.

Figure 2

TEM (a) and HRTEM (b) images of the interface SiC/Al without interfacial products for an SiCp/Al–1.0Mg–0.6Si–0.25Cu composite, (c) indexed pattern of the SAED presented in Figure 2b, and (d) IFFT of the area 3 in Figure 2b (Reproduced with permission from [84].).

Otherwise, when the Al₄C₃ layer is present, the interfacial bond is stronger, and fracture propagates through the reinforcement instead of through the interface [83]. This makes a more efficient load transfer across the interface possible and leads to an increase in mechanical properties, although ductility decreases. Tham et al. [83] reported that values of ultimate tensile and yield strengths increased, respectively, 10 and 20% in composites with the presence of the Al₄C₃ layer, compared to composites without it. Nevertheless, if thickness of this interface is larger than a critical size, mechanical properties could be reduced. This can also occur when Al₄C₃ is obtained as isolated particles. Figure 3a and b shows an example of the formation of Al₄C₃ as a reaction product at the Al/SiC interface, for a 2009 Al alloy reinforced with 15 vol% SiC particles [86]. HRTEM in Figure 3b shows that the crystallographic orientation for Al₄C₃ corresponds to (−11 to 24)Al₄C₃ || (−101 to 6)SiC, without crystallographic orientation Al₄C₃/Al matrix. The composite obtained under these conditions presented excellent performance because SiC is completely bonded to the matrix, avoiding decohesion and improving load transfer, which raised the UTS. These authors also reported a crystallographic orientation (1–102)Al₄C₃ || (1–100)SiC. These different orientations depend on the nucleation conditions for Al on SiC and are focused to reduce the free energy of the interfaces. Aluminum has a face-centered cubic (fcc) lattice, with a = 0.405 nm. Otherwise, SiC can be present as 6H α-SiC, which has a hexagonal structure, with lattice parameters such as a = 0.30807 nm and c = 1.51174 nm, or as β-SiC, with a cubic structure (a = 0.43581 nm) [87]. On the other hand, Al₄C₃ crystallizes in the hexagonal unit cell, with crystallographic parameters a = 0.333 nm and c = 2.498 nm [88]. These characteristics make possible different orientation relationships for hexagonal SiC and Al₄C₃, with small lattice mismatches. The Al₄C₃ reaction layer has the tendency to form semicoherent interfaces and orientation relationships with the aluminum matrix and the SiC particles [83]. Figures 2 and 3 show the effect that can exert not only the reinforcing type but also the addition of alloying elements on the crystallography of the interfaces. According to the matrix–reinforcement interaction, reaction products can be favored or avoided. These intermetallics commonly grows through mechanisms involving diffusion of elements from the reinforcement to the matrix and vice versa, presenting planes with the same crystallographic orientations. These interlayers increase the strength of the composites due to their strong bonding.

Figure 3

TEM (a) and HRTEM (b) images of the interface SiC/Al with the presence of the Al₄C₃ interlayer for a 2009 Al alloy reinforced with 15% SiC particles (Reproduced with permission from [86].).

The occurrence of severe interfacial reactions can be detected not only by TEM but also by using optical microscopy (OM) or scanning electron microscopy (SEM), with ledges and pits at the interfaces, as showed in the study by Sritharan et al. [89]. This is shown in Figure 4a and b for an A359 aluminum casting alloy (9.27 Si and 0.55 Mg, in wt%) reinforced with 20 vol% of SiC particles. Differences in the SiC particles are observed for this composite before (Figure 4a) and after (Figure 4b) melting and held at 650°C for 48 h.

Figure 4

SEM images of an A359 + 20 vol% SiC particles AMC. (a) Before elevated temperature exposure, with SiC particles presenting relatively flat surfaces and interfaces. (b) After being held at 650 C for 48 h, with SiC surfaces presenting pits and ledges (Reproduced with permission from [89].).

Because of the oxidation of SiC, the formation of SiO₂ occurs according to [90,91]:

The thickness of the SiO₂ layer is related to temperature: higher temperature increases layer thickness due to a higher diffusivity of oxygen, also improving wetting [91]. Its formation can modify the reaction to form Al₄C₃ because SiO₂ acts as a barrier and avoids further oxidation. Amorphous SiO₂ layers cover SiC, which acts as an intermediate to form stable interfaces and avoid Al₄C₃ formation [87]. The kinetics to obtain SiO₂ follows a linear-parabolic model, stabilizing the SiO₂ layer thickness when passive oxidation occurs [91]. After its formation, SiO₂ layer thickness could decrease through a reaction with molten Al, obtaining Al₂O₃ that reinforces the composite [92,93,94,95]. This reaction is expressed as follows:

2.1.2 SiC in Al matrices with Si

The addition of Si to an Al alloy creates an equilibrium among SiC, Al₄C₃, and liquid Al, which is the possible modification of equation (1) due to the differences in the liquid composition [96,97]:

This can diminish or even avoid the formation of Al₄C₃ at the interface [96]. The required Si content in the matrix to prevent it increases with the increase in the melting temperature [78]. Schwabe et al. [97] found that at 750°C, this content was 6 at%, while at 1,300°C, the content was 16 at%. As shown in Figure 2a–d, Wang et al. [85] obtained a clean surface, without interfacial reactants, using an Al alloy with 19.5% Si to avoid Al₄C₃ formation. The dependence between Si content and temperature for the formation of Al₄C₃ was graphically presented in the study by Salvo et al. [82]. This study could be consulted for establishing manufacturing conditions according to the Si content.

2.1.3 SiC in Al matrices with Mg

It has been reported that the addition of Mg to Al alloys reinforced with SiC modifies equation (1), where Al₄C₃ formation is now accompanied with the formation of Mg₂Si [79,80,81]. This can occur according to equations (5) and (6) [79,80,81,68]:

Due to the formation of SiO₂ through the oxidation of SiC (according to equation (2)), the following reactions have been reported when Mg is added to Al matrices [94,95]:

These reactions show that Mg improves wetting between Al and SiC particles [93,98,99]. Mahendra et al. [100] reported that the addition of 1.5 wt% of Mg to a pure Al matrix was enough for improving wetting, obtaining better distribution of this reinforcement, and enhancing the mechanical properties.

Another effect of the addition of Mg is the formation of not only MgO but also spinel MgAl₂O₄, as reported by Gonzalez et al. [101]. The bonding strength between spinel phase and Al matrix is more than the double of the strength for the bond Al/Al₄C₃ [99]. Crystallographic studies using TEM and electron probe micro analysis demonstrated that transformation from MgO to MgAl₂O₄ is associated with a volume reduction of 11%. This led to an incomplete coverage of the SiC particles, which presented degradation. These interfacial layers were discontinuous, with MgAl₂O₄ crystals significantly bigger than the MgO ones [101]. MgO formation is favored for higher Mg contents at the SiC/Al interface, while MgAl₂O₄ is more stable for low Mg contents. Li et al. [102] found that spinel improved the bonding strength of the interface through the change of the fracture mechanism from pull-out (when MgO was at the interface) to tensile loading induced. In addition, Gonzalez et al. [101] found that the addition of Mg made the alloy age hardenable due to the rejection of combination with Si as a consequence of Al₄C₃ formation at the interface (equations (5)–(7)). The following reactions to obtain spinel MgAl₂O₄ can occur [94,103,104,105]:

MgAl₂O₄ can decompose after its formation, according to [101,102,103,104]:

2.1.4 SiC in Al matrices with Si and Mg

The addition of Si and Mg as alloying elements could lead to different modifications as the effect of Si stops Al₄C₃ formation and Mg reduces SiO₂ and forms Mg₂Si. As for the case of Al–Mg alloys, in the case of Al–Mg–Si system, there is the formation of MgO and MgAl₂O₄. This could modify Si and/or Mg quantities reported as optima for improving mechanical properties of the matrix. Sritharan et al. [106] found that matrices for optimum strengthening when SiC is used as reinforcement were Al–Mg and Al–Mg–Si with high Mg content (alloys 5XXX and 6XXX, with Mg contents from 0.5 to 5.5 wt%), and this is due to the formation of Mg₂Si precipitates. An example of the formation of beneficial interfacial layers in Al alloys with Si and Mg can be found in the study by Gao et al. [107]. These authors reported a continuous and well-bonded coherent SiC/MgAl₂O₄/Al interface, without any gaps, which is beneficial for load transfer and strength improvement. This interface of ∼10 nm can be observed in the TEM images of Figure 5a and b, where the HRTEM image of Figure 5b shows the coherency for the interfacial bond SiC/MgAl₂O₄ [107]. As mentioned earlier, the formation of MgO or MgAl₂O₄ as reaction layers depends on the alloying element content. For Al–Si–Mg alloys, the combination of Mg and Si to precipitate as Mg₂Si decreased the quantity of Mg, leading to obtain MgAl₂O₄ instead of MgO. This can be seen in the HRTEM images of Figure 5c and d for Al–Mg–Si/SiC composites, where MgO and MgAl₂O₄ interfaces look clean, indicating a good wetting and bonding [102]. MgO (of about 10 nm) was formed for a supersaturated solid solution, where the content of Mg in solution is high, while MgAl₂O₄ (50–100 nm) formation was for a matrix with Mg₂Si precipitates, which led to decrease in the content of Mg in the solid solution. This demonstrates the effect that can exert the reinforcement on the matrix and both matrix and reinforcement on the interfacial layer. Contrarily, Tekmen and Cocen [73] found the modification of the precipitation process due to the addition of SiC. They reported that for Al–Mg–Si matrices reinforced with SiC particles, the hardness peak was obtained at lower duration than that for the matrix alloy. The increase in the SiC content also led to obtaining higher hardness. These authors found that for a composite with 20 vol% SiC, the increase in Mg and Si amounts led to obtaining higher hardening [73]. This demonstrates that both the alloying element content and the reinforcement volume fraction influence matrix precipitation. A good bonding achieved at the SiC/MgAl₂O₄ interface with semicoherency has been reported by Luo [87], with four types of orientation relationships, which was favorable to the composite mechanical properties. Spinel MgAl₂O₄ has a face-centered cubic structure with a lattice parameter a = 0.8075 nm, while MgO is also fcc, with a = 0.421 nm [87]. These authors reported that due to the cubic cell, MgO may present the same orientation relationship of the spinel phase with SiC. The lattice mismatch Al/MgAl₂O₄ is 0.0028 (by counting 1/2 for the spinel phase), also presenting local structural rearrangements to achieve some degree of semicoherency at their interface.

Figure 5

TEM images for pre-oxidized SiCp/Al–Cu–Mg-Si-Mn composites showing: (a) the SiC/MgAl₂O₄/Al interface and (b) HRTEM of the interface (area B in subsection b) (Reproduced with permission from [107].) (c and d) HRTEM of the interfaces obtained for Al–Mg–Si/SiC composites after a solution heat treatment (c), where MgO is obtained and overaged, (d) where MgAl₂O₄ can be observed (Reproduced with permission from [102].).

Ren et al. [98] studied the combined effect of Si and Mg for Al alloys with different contents of Si (6, 12, and 18 wt%) and Mg (4, 8, and 12%), obtaining by infiltration composites reinforced with 55% of SiC particles. They found that Si contents less than 6% led to poor wettability and high porosity, not completely infiltrating SiC, while higher Si contents avoided the formation of Al₄C₃, and the composites presented higher elastic moduli, thermal dimensional stabilities, and thermal conductivities. Otherwise, the addition of Mg avoided the formation of Al oxides at the Al/SiC interface, lowering viscosity and surface tension [95]. The work of Sukumaran et al. [108] agrees with these results, demonstrating that both Si and Mg need to be added as alloying elements for obtaining good wettability SiC matrix, with concomitant effects essential to obtain composites with improved mechanical properties. Research including matrices combining these elements is insufficient and needs more attention.

2.1.5 SiC in Al matrices with Cu

Related to the effect of Cu on SiC–Al matrix interactions, Feng et al. [109] coated SiC particles with Cu using electroless plating, depositing a layer of Cu via dipping. They found a supersaturated solid solution containing 37 wt% of Cu surrounding SiC, which strengthened the composite due to the lattice distortion. Another consequence of this coating was the precipitation of Al₂Cu around the SiC particles, hindering the dislocation movement and strengthening the composite. The presence of Cu additionally improved mechanical properties by increasing the interfacial bonding SiC–Al matrix, modifying it from SiC/Al to SiC/Cu and Cu–Al. This allowed a better load transfer from the Al matrix to the reinforcement [110]. These results show the effect of Cu improving mechanical properties through different mechanisms, which are examples of how a composite can present higher mechanical properties not only due to the addition of reinforcement but also due to other interactions matrix/reinforcement.

2.1.6 SiC in Al–Si–Cu–Mg matrices

The addition of more than two alloying elements to Al matrices contributes to stimulate interactions in the composites. An et al. [36] found that for these alloys Si reduced the interfacial free energy, improving the wettability of SiC. The content of this element was higher in the interfacial zone, diminishing the formation of Al₄C₃. The enrichment of the interface was not detected for Cu or Mg additions. Besides, Cu originated a slight deterioration of the wettability, also decreasing the formation of Al₄C₃. Moreover, Mg favored the wettability disrupting the oxide film covering Al [36]. For the combined study of these three elements, Durbadal and Srinath [111] found a thick reaction layer for a 2,124 Al–4.4 Cu–1.5 Mg–0.8 Si alloy reinforced with 10 wt% SiC. The interfacial products obtained when this alloy was re-melted at 900°C were associated with a high concentration of Cu, Mg, Si, and C, reporting the reaction of Al + SiC to form Al₄C₃ and Si. This indicated that SiC was dissociated, while Mg participated in the formation of spinel MgAl₂O₄.

Related to the effect of the reinforcement on the precipitation process for these alloys, Song et al. [112] found that for a 2,024 Al–4 Cu–1.5 Mg–0.5 Si alloy the amount of S′(Al₂CuMg) precipitates decreased due to the addition of SiC. This behavior is observed in Figure 6a–d. For the unreinforced alloy, an extensive precipitation was obtained, as shown in Figure 6a, while Figure 6b and c show that precipitation considerably decreased with the increase in the SiC content from 5 to 15 wt%, even presenting precipitation free zones in Figure 6d [112]. These modifications have a significant effect on the mechanical properties of the composites.

Figure 6

TEM images of a 2024 aluminum alloy reinforced with (a) 0, (b) 5 wt%, (c) 10 wt%, and (d) 15 wt% of SiC, after being initially aged at 110°C for 8 h, and then aged at 190°C for 24 h. S′ precipitates and SiC reinforcements can be observed (Reproduced with permission from [112].).

Examples as shown in Figure 6a–d and the explanations of these behavior were explained earlier for binary and ternary alloy systems reinforced with SiC. Independently of the alloying elements, the addition of reinforcements may lead to modify precipitation sequences and kinetics. Other precipitates have been observed in the interfacial region for this alloy system reinforced with SiC, as demonstrated in the study by Gatea et al. [113]. They found Al₂Cu, Al₄Cu₂Mg₈Si₇, and MgAl₂ at the SiC/Al alloy interface for an Al–1 Cu–1.2 Mg–0.8 Si alloy reinforced with SiC particles. On the other hand, for an Al–3.8 Cu–1.2 Mg–0.25 Si alloy, Guo et al. [114] observed that Mg was preferentially localized close to the SiC/matrix interfaces due to the formation of Mg₂Si. This led to the increase of the matrix hardness near the reinforcements. Pal et al. [110] reported that for an Al–5.4 Cu–1.5 Mg–0.3 Si alloy reinforced with SiC particles, the aging kinetic for the matrix was slower compared to the unreinforced alloy. This behavior was attributed to the segregation of alloying elements from the matrix to the interface. The decrease of Cu and Mg contents in the matrix reduced solute supersaturation, leading to this modification of the precipitation kinetic [110,115,116]. The nucleation of precipitates such as Al₂Cu can occur in the presence of Si, released by the interfacial reactions. Si–vacancies pairs favor homogeneous nucleation through the formation of dislocation loops [68]. Varma et al. [115] reported that the size of the reinforcement also affects precipitation hardening. They found that the use of particles of ∼1 µm in diameter led to a higher precipitation hardening at a lower time compared to composites containing coarser reinforcements, which was attributed to faster diffusion of the alloying elements in Al.

Other properties could be also affected depending on the alloying elements content. For example, Wang et al. [117] found that the wear resistance of AMC with Al–Cu, Al–Mg, and Al–Si alloy matrices reinforced with SiC fibers decreased in that specific order, although hardness behavior of the matrices was inverse (48 HV for Al–Cu, 57 HV for Al–Mg, and 67 HV for Al–Si). This was attributed to the effect of interfacial reactions between fibers and matrices, not depending on the matrices hardness: a smooth Al₄C₃ coherent interface for Al–Cu; a nonsmooth interface, enough to increase both the interfacial bonding strength and wear resistance for Al–Mg; and no interfacial reactions for the composite with Al–Si matrix, presenting the lowest wear resistance between these composites. These results show that the formation of an interfacial bond particle–matrix is essential for increasing the wear resistance of these composites.

After the analysis of the interactions between SiC and Al matrices with Si, Cu, and/or Mg, it was found that the effect of each one of these elements is different. Adding Si may lead to minimize or even avoid the reduction reaction of SiC with Al, lead to a higher elastic modulus, and decrease properties such as wear resistance, and this is due to low interfacial strength. On the other hand, Mg modifies wettability, but optima additions are not well specified or did not coincide in the literature. This element forms different interfacial compounds due to its interaction with SiC. Finally, interfacial compounds were not found when Cu was added to Al matrices apart from the characteristic precipitates obtained according to the matrix composition.

2.2.1 WC in Al matrices with Si, Cu, and/or Mg

According to Sharath Kumar et al. [39], WC offers strong bonding with aluminum. These authors did not find any interfacial product for the reinforcement of an Al–5 Si–3 Cu–0.5 Mn alloy adding 7.5 wt% of WC to the molten alloy at 800°C. Nevertheless, it is reported that at higher temperatures the reaction between Al and WC particles could lead to the formation of Al₁₂W and Al₄C₃, in addition to the eutectic structure of the Al–12 Si alloy [120]. For an Al–2.5 Mg–1.6 Cu–1.5 Zn–0.5 Fe–0.4 Si alloy reinforced with WC using stir casting, the formation of Al₅W was found, needing a more in-depth analysis of the interface [121]. Zhang et al. [122] reported that intermetallics such as Al₅W and Al₁₂W can improve properties including mechanical strength, oxidation resistance, and thermal stability. Otherwise, Lekatou et al. [38] reported the formation of Al₁₂W as coarse polygonal particles and Al₅W acicular plates for an Al1050 alloy reinforced with 1 vol% of WC processed at 830°C. These authors also reported the formation of Al₄C₃ through the following reactions depending on the contact temperature:

At 830°C:

On cooling at 697°C:

On cooling at 661°C:

The presence of intermetallic particles (e.g., Al₁₂W) may reduce plastic deformation in the matrix due to crack growth mechanisms. The use of techniques implying elevated temperatures leads to the formation of different interfacial products, which could be detrimental to the compactness behavior of the MMC, as reported by Li et al. [123]. They reinforced a 6061Al alloy (0.8–1.2Mg, 0.40–0.8Si, 0.15–0.40Cu) with an Al–11.52Si wire and WC powder using laser welding and found rod-like Al₄W, W₂C, and Al₄C₃. The consequences of this process are shown in Figure 7a and b, where there are observed irregular WC particles before laser welding (Figure 7a); the microstructure is obtained after the manufacturing process (Figure 7b), consisting of Al matrix and the aforementioned rod-like phases. The X-ray diffraction (XRD) diffractogram depicted in Figure 7c demonstrates the formation of these new phases. No interfacial products were detected for low-temperature manufacturing methods, although excellent interfacial bonding is reported [124,125], contrary to the observed for other reinforcements added by these methods where decohesion occurs (e.g., TiO₂ [126]). Nevertheless, Singh and Sharma [127] studied an Al–18.5 Si–1.6 Cu–1.1 Mg–0.5 Fe alloy reinforced with 3–12 wt% WC particles using stir casting, revealing that Al₅W and Al₁₂W precipitated, but only after T6 heat treatment. Although this alloy has high alloying elements content, no other intermetallics were detected. An example of the strong bonding for WC/Al composites is the manufacturing process for reinforcing Al with 3 vol% of WC particles through roll bonding at 250°C, where it was possible to obtain excellent WC particle distribution, without interfacial reactions [128]. A TEM image in Figure 7d shows this composite and WC particles with smooth and clean interfaces, without visible decohesion. This led to a composite with enhanced tensile, hardness, and wear properties, also attributed to a uniform distribution of the reinforcements and a strong interfacial bonding.

Figure 7

SEM images of (a) WC particles used as reinforcement and (b) the composite obtained after adding the WC particles to a 6061 Al alloy (0.8–1.2 Mg, 0.40–0.8 Si, 0.7max.Fe, 0.15–0.40 Cu), using laser welding. (c) XRD diffractogram of the obtained WC/Al composite. (Reproduced with permission from [123].) (d) TEM image of an Al/3vol%WC composite (Reproduced with permission from [128].).

Other elevated temperature processes that reveal the formation of different intermetallics have been reported, as shown in the work of Prasad et al. [129], coating an Al356.2 alloy with WC-Co, which led to the formation of W₂C. This resulted in the loss of reinforcement, embrittling the matrix and decreasing the mechanical properties. Moreover, Staia et al. [130] studied the superficial modification of an Al–7.0 Si–0.3 Mg alloy reinforcing with WC particles by laser alloying and identified W₂C produced by the oxidation of WC. This led to the release of C into the matrix and to the formation of Al₄C₃. Laser processing and the subsequent solidification process led to the formation of W₆C_2.54 and WC _x due to the partial dissolution of WC in the melting pool. WC is one of the reinforcements, which need more attention and research, including the analysis of the possible reaction products. This is shown in Figure 1.

Although the earlier analyzed alloys are rich in Si, Cu, and/or Mg, information was not found related to the formation of other phases with their presence. This shows the inertness of WC and the possibilities of its use. Nevertheless, more research and in-depth analysis of the interfacial regions are needed to elucidate if these elements have any effect on the interfaces, such as wettability modification, segregation, or effects on the mechanical properties.

2.3.1 B₄C in Al matrices

It is reported that when B₄C is in contact with molten Al above 1,000°C, the formation of boron oxide (B₂O₃) films occurs around B₄C, increasing its wettability [132,133]. AlB₄₀C₄, Al₂B₄₈C₈, Al₂B₅₁C₈, Al₃B₄₈C₈, Al₈B₄C₇, and B₁₂(B,C,Al)₃ have been reported as interfacial reaction products [131,134,135,136]. Below 1,000°C, wettability is poor, being added different compounds (e.g., K₂TiF₆) to the removal of the oxide film from the aluminum surface, which reactivates wetting [131,134,135]. Some reaction products formed at these lower temperatures include Al₃BC, AlB₂, Al₂B₄C₄, and Al₃B₄₈C₂. Al₃BC formation occurs as follows [137,138,139,140,141,142]:

If the temperature is increased to 1,200–1,400°C, the formation of Al₄C₃, AlB₂₄C₄, Al₈B₄C₇, and AlB₁₂C₂ can occur [141]. The main reaction for the formation of Al₄C₃ is as follows [142]:

Although the extensive reaction between Al and B₄C occurs above 1,100°C, it has been reported that it can even start at temperatures as low as 450°C. At temperatures higher than 600°C, the depletion of the aluminum content in the composite begins due to the formation of Al₂B [143].

2.3.2 B₄C in Al matrices with Mg

The presence of Mg in Al matrices results in the formation of Al _x Mg_(1−x)B₂, Al₄C₃, Al _x (B–C–O)_1−x , and MgB₆ [104]. This element significantly increases wetting due to its ability to form strong bonds between matrix and B₄C by producing layers of (Al,Mg)B₂ [137,138,139]. The formation of Mg(Al)B₂ was also reported by Zhou et al. [140], who found that the in situ formation of this reinforcement led to a higher strength. Initially, its formation started in equation (17) and then in the following equation [137,138,139,140]:

For Al-(2.5–4.5) wt% Mg reinforced with B₄C, fabricated by pressureless infiltration at 800°C, Lee et al. [144] found the formation of AlB₂, AlB₁₂ (Al₃B₄₈C₂), AlB₁₀ (AlB₂₄C₄), and Al₃BC. They also detected the presence of MgAl₂O₄, phase reported for composite systems such as Al/SiC, Al/Al₂O₃, Al/BN, and Al/Si₃N₄ [145,146,147]. Otherwise, Lucas et al. [142] reported that reaction products for an Al–7Si–0.35Mg alloy reinforced with B₄C were Al₄C₃, MgB₂, and Al _x Mg_(1−x)B₂ (where 0 < x < 1). These authors found a higher Mg concentration in the interfacial reaction zone of B₄C, compared to the same region when the reinforcement was SiC. The reaction Mg/B₄C for Mg alloys is as follows [148]:

Similar to the formation of Al₄C₃ and Mg₂Si for the reaction of Al and Mg with SiC (in equation (5)), when B₄C is used as reinforcement, the following reaction occurs:

The transformation of MgB₂ into Al _x Mg_(1−x)B₂ can also occur due to the diffusion of Al atoms into MgB₂, substituting Mg atoms [149]. This phase, reported as Al0_0.29Mg0_0.71B₂ by Suárez et al. [149], increased the microhardness of these composites. The diffusion of Al to form this phase can be presented by the following transformation:

It can be more generally presented as follows:

2.3.3 B₄C in Al matrices with Si

When an Al molten alloy in contact with B₄C contains Si, the formation of Al₄C₃ can be avoided due to the reaction between Si (dissolved) and C (in B₄C) [33]. This originates a reduction of the carbon activity by depleting its content in the carbide phase and leads to the formation of SiC instead of Al₄C₃. The presence of SiC improves the mechanical properties of these composites. Frage et al. [33] found that the Si content threshold value (necessary for preventing the formation Al₄C₃) is about 12 wt%, then favoring the formation of SiC.

Different intermetallics have been reported for composites obtained through the infiltration of B₄C with liquid Al and Al alloys (13 and 40 wt% Si) at 1,100°C, followed by a heat treatment for 10 h at 1,200°C [35]. AlB₁₂ was obtained after infiltration, and ternary compounds such as AlB₁₂C₂, Al₈C₇B₄, and Al₄C₄B_x are obtained after heat treatment. Al₄C₃ was obtained for low Si contents, while SiC was favored when the Si content was high.

2.3.4 B₄C in Al–Si–Cu–Mg matrices

For quaternary matrices, such as those reported in the research of Li et al. [150] for an Al–1.0 Mg–0.65 Si–0.25 Cu alloy reinforced with 20 wt% of B₄C, some of the reaction products are Al₃BC, MgO, MgAl₂O₄, MgB₇, Mg_0.78Al_0.75B₁₄, Al₁₂BC₂, and Al₄SiC₄. These authors divided reactions involving Mg and Si into two series according to their interactions with B₄C: first, the reactions involving B and Mg, and once they ended, the reactions involving Al, Si, and C stars (at temperatures higher than 620°C). These reactions initiated with equation (17), including the reaction of equation (11) to form MgAl₂O₄ and the following reactions:

These reactions lead to severe depletion of Mg in the Al matrix, decreasing its precipitation hardening. The deterioration of B₄C also leads to a decrease in the mechanical properties. Xue et al. [151] found that for these alloys, the addition of B₄C inhibits the formation of Guinier Preston zones and changes the order of precipitation. With the increase of the B₄C volume fraction, time for reaching hardening-peak decreased, but its intensity increased. Zhou et al. [152] studied an Al–1.0 Mg–0.65 Si–0.25 Cu (wt%) matrix reinforced with 26 wt% B₄C and found Al₃BC nanoparticles at the interface, surrounded by a thin layer of Cu. This led to the absence of the typical Q (Al₅Cu₂Mg₈Si₆) precipitates in the zone near B₄C particles. Precipitates were MgB₂ instead of Q. Al₃BC and MgB₂ can be obtained through the reaction of equation (18), followed by:

Moreover, Lee et al. [144] found MgO and TiB₂ instead of AlB₂ and Al₃BC for an Al–0.529 Si–0.954 Mg–0.172 Cu alloy reinforced with B₄C, with good interfacial. Besides, they reported the diffusion of Al toward B₄C during the stir casting process, followed by the interfacial reaction to generate an Al₄C₃ layer. Pozdniakov et al. [153] studied an Al–5% Cu alloy reinforced with 2, 5, and 7 wt% of B₄C manufactured using squeeze casting and found that reactions alloy/B₄C only led to the formation of Al₃BC and AlB₂, which assured good wettability. No reactions between Cu and B₄C were found, which agrees with the aforementioned reports. These authors reported the acceleration of the Al₂Cu precipitation due to the addition of B₄C. Manning and Gurganus [154] also found the null effect for Cu. Its addition did not cause significant modifications on the contact angle between B₄C and the molten Al alloy. Nevertheless, for high contents of Mg and Si, this angle significantly decreased, leading to good wettability and homogeneous distribution of B₄C particles. An example of well-distributed B₄C particles is shown in Figure 8a for an Al–1.0 Mg–0.65 Si–0.25 Cu (wt%) alloy reinforced with 20 wt% of B₄C [150]. Interfacial reactions products for this composite were Al₃BC, Al₄SiC₄, Mg₂B₂O₅, Mg₂AlO₄, and MgB₇. This is shown in the XRD diffractogram of Figure 8b for the composite processed under different variables. Otherwise, reaction products surrounding B₄C can be observed in the high-angle annular dark field (HAADF)-STEM image of Figure 8c [152], for a 6061 Al alloy reinforced with 26 wt% of B₄C particles. These products formed a continuous layer consisting of polygonal nanoparticles (arrowed) at the interfacial region. Figure 8d shows that these nanoparticles correspond to Al₃BC, identified using SAED. Nanometric MgB₂ particles were also found (arrowed nanorod), formed in the matrix close to the B₄C reinforcement. HAADF-STEM and energy-dispersive X-ray spectroscopy (EDS) elemental mapping analyses of the Al₃BC nanoparticles showed a thin layer of Cu surrounding the Al₃BC particle, as it is shown in Figure 8e. A high-resolution HAADF image of Figure 8f shows the bright contrast of this layer, which implies the presence of heavy Cu atoms. Cu and Mg were also segregated at the MgB₂/matrix interface. Segregation occurs to reduce the interfacial energy, enhancing nucleation rate and avoiding precipitate coarsening [152]. Al₃BC is reported as a hexagonal structure with lattice parameters a = 0.6046 nm and c = 1.154 nm [145]. This compound has been reported together with AlB₂, which also has a hexagonal structure, with a = 0.311 nm and c = 0.498 nm [145]. The crystal structure of B₄C is rhombohedra1 or hexagonal (a = 0.561 nm and c = 1.214 nm), differing from face-centered cubic structure of Al, a fact that diminish the possibility of getting a good bond directly between them, but the formation of these interfacial compounds favors bonding, significantly influencing the loading transfer between reinforcements and matrix, and consequently improving the mechanical properties of the composites [154]. This interlayer needs to be more studied to analyze their growth mechanism, orientation relationships, and other crystallographic characteristics.

Figure 8

(a) OM image of an Al–1.0 Mg–0.65 Si–0.25 Cu (wt%) alloy reinforced with 20 wt% of B₄C. (Reproduced with permission from [150].) (b) XRD diffractograms of this composite under different processing variables. (Reproduced with permission from [150].) (c) HAADF-STEM image of the B₄C reinforcement surrounded by reaction products. (d) TEM image of the Al₃BC nanoparticles at the interface Al/B₄C (SAED pattern inserted). (e) HAADF-STEM image for the analysis of the Al₃BC/matrix interface. EDS mappings show the elements distribution. (f) High-resolution HAADF image of the Al₃BC particle, surrounded by a 0.7 nm Cu layer ((c)–(f) Reproduced with permission from [152].).

As for the case of Al alloys reinforced with SiC, important research was found related to the interactions of these matrices with B₄C. These results show that again, Cu is the alloying element with the lowest interaction and reactions products when it is added to Al alloys. Its role and participation in different reactions are barely known, except its presence in precipitation processes. On the other hand, the use of Mg excels, increasing wettability and forming a wide variety of compounds, which increase the mechanical properties of the composite. Related to the addition of Si to these matrices, it is also beneficial because the formation of Al₄C₃ is avoided, obtaining SiC as a reaction product, which improves mechanical properties.

2.4.1 BN in Al matrices

For Al matrices reinforced with BN, it has been reported the formation of AlN and AlB₂ in the interfacial region, by the following direct reaction [159]:

This reaction occurs for temperatures <985°C, while at higher temperatures it is reported [159]:

2.4.2 BN in Al matrices with Mg

Lee et al. [160] reported that the presence of Mg origins the formation of reaction products containing Al, Mg, and B, named Al(Mg)–boride. This is for an Al–4.5% Mg alloy reinforced by infiltration at 800°C with 5 and 7.5% h-BN particles. Besides, they reported the reaction between Mg and N to form Mg₃N₂ surrounding BN, which enhanced infiltration. Khatavkar et al. [161] reported that the wettability of h-BN was improved due to the addition of 1% Mg for a 2024 Al alloy. Reaction to form Mg₃N₂ can be summarized as follows [160,161]:

After this reaction, Mg₃N₂ reacts with Al to obtain AlN according to:

These reactions were also reported by Woo et al. [162], who besides found MgAlB₂.

2.4.3 BN in Al matrices with Mg and Si

Konopatsky et al. [163] found that for an Al–10 Si–Mg alloy with 1, 3, and 5 wt% of BN used for 3D printing (∼2,627°C were reached during ∼0.15 ms), there was the formation of AlN at the BN surface, while AlB₂ nanoparticles (∼10 nm) were observed in the matrix near the interface. These authors did not report the effect of Si and Mg, or the formation of products due to reactions with these elements.

2.4.4 BN in Al matrices with Cu

The precipitation process can be also modified for Al alloys reinforced with BN, as it is observed in the TEM images of Figure 9a and b for an Al–5Cu alloy, and for an Al–5Cu/5%BN composite, respectively, both aged at 200°C for 9 h [157]. As it is clearly noted, differences in the precipitates are significant: θ′-phase averaging ∼200 nm in length and 20 nm in width precipitated for the Al–5 Cu alloy (Figure 9a), corresponding to age-peak, while the stable θ-phase with 500 nm in length and 200 nm in width precipitated for the Al–5Cu/5%BN composite (Figure 9b), corresponding to the overaging state. Besides, they found the formation of strong bonded Al/BN interfaces, without gaps; while AlB₂ was obtained as an interfacial product, without reactions involving the presence of Cu. An example of a well-bonded BN/Al matrix interface can be observed in Figure 9c and d for an Al–5.3Cu/3%hBN composite (in wt%) processed at 640°C through semi-solid powder metallurgy technique. Figure 9c shows a TEM image of the composite, while the HRTEM image of Figure 9d shows that BN was strongly bonded with the Al matrix, through a Al₂O₃ nanolayer formed due to the oxidation of Al. The presence of BN and the interfacial characteristics allowed to obtain elevated hardness, compressive strength, and fracture strain [156]. The aforementioned AlN and AlB₂ interfaces can be obtained even at temperatures slightly below the Al melting point, following two mechanisms: (i) transformations inside Al grains (AlB₂) or (ii) along the Al grain boundaries (AlN). Besides, Al₂O₃ often is obtained due to the substitution of N atoms by O in the growing AlN phase, being more favorable the formation of Al₂O₃ than AlN. h-BN has hexagonal crystal lattice, with a = 0.250 nm and c = 0.662 nm, with atoms of N of 0.146 nm, very similar to the ionic radius of O (0.140 nm), occurring this substitution. These factors also favor the formation of AlB₂, which has a hexagonal lattice with a = 0.301 nm and c = 0.325 nm, while AlN has a hexagonal structure, with a = 0.311 nm and c = 0.498 nm [160]. The preferential formation of the AlB₂ phase compared to AlN in the solid state could be explained because B atoms can more easily penetrate into the metallic Al lattice compared to N. Besides, the lattice mismatch for c parameter for BN/AlB₂ is 0.006 (by counting 1/2 for BN), which could lead to achieve some degree of semicoherency at their interface. At higher temperatures, the preferential formation of AlN occurs due to a more negative value of the Gibbs free energy [163].

Figure 9

TEM images after aging 9 h at 200°C for: (a) an Al–5% Cu alloy and (b) an Al–5% Cu/5% BN composite. (Reproduced with permission from [157].) (c) TEM image for an Al–5.3 Cu/3 BN composite and (d) HRTEM image of this composite showing an Al/Al₂O₃/BN interface (Reproduced with permission from [156].).

Contrary to that observed for the aforementioned carbides, for BN, the information is limited, even for the interaction of this reinforcement with Al. Only research works with reactions involving Mg were reported, reporting that this element improves wettability and mechanical properties of the composite. The effects of Si and Cu were not found, although some works reported the use of Al matrices containing these elements. This shows that more research is needed for the use of BN as reinforcement and also shows its interactions with different Al matrices.

2.5.1 Si₃N₄ in Al matrices

For the case of the interaction between molten Al and Si₃N₄, the following reaction has been reported at temperatures ranging from 827 to 1,027°C [97]:

This reaction is followed by the dissolution of Si in liquid aluminum according to the equilibrium of equation (4). As happens for the equilibrium SiC-Si, in this case, Si₃N₄–Si can control the direction of this reaction, bonding directly an interface Si₃N₄/Al, or with the interfacial layer Si₃N₄/AlN/Al. Lu et al. [166] reported the formation of AlN even for an alloy with 10% Si at temperatures above 750°C, improving the mechanical properties of the composite until reaching 850°C. In this temperature range, a thin reaction layer led to increase in the interfacial bonding, improving hardness, fracture toughness, and flexural strength. The thickness of the reaction layer increased with the increase in the temperature due to a higher reactivity, but at temperatures above 850°C, the interface embrittlement deteriorated the mechanical properties. At this point, the fracture modes changed from ductile intergranular to brittle intergranular. At temperatures higher than 950°C, intense interfacial reaction occurred, and the composites became AlN and Si due to the complete consumption of Si₃N₄.

2.5.2 Si₃N₄ in Al matrices with Mg

Akhtar and Guo [167] studied a Si₃N₄/Al–2Mg composite manufactured by pressureless melt infiltration and reported the formation of AlN and Al_13.27Si_0.47. De la Peña and Pech-Canul [168] also reported these reaction products, and the formation of MgO and MgAl₂O₄, which reduced wetting through reactions involving Mg₃N₂, described as follows:

This reaction is followed by other reactions that lead to the lack of wetting:

Besides, it is reported that the surface of Si₃N₄ can be covered by Si₂N₂O due to the interaction of the reinforcement with O₂. Then, the reaction Al alloy/Si₃N₄ is substituted by a reaction including Si₂N₂O instead of the original reinforcement, being as follows [168]:

Koike et al. [169] reported partial melting at matrix/Si₃N₄ interfaces and at the grain boundaries of the matrix for temperatures as low as 450°C. This occurred for an Al–2.46 Mg alloy reinforced with 20% Si₃N₄ particles. The cause of this behavior was the segregation of Mg and Si to interfaces and grain boundaries, decreasing solidus temperature. The effect of Mg additions was also studied by Wang et al. [170], who analyzed the infiltration kinetics of Si₃N₄/Al–Mg composites fabricated by pressureless infiltration. These authors found that the increase in the Mg content from 2 to 10 wt% significantly improved the infiltration process. Nevertheless, for higher Mg contents, they reported an adverse effect due to interfacial reactions, forming Mg₂Si, MgO, and Al₂O₃. As can be seen, a nonheat treatable Al–Mg alloy was modified to a heat-treatable Al–Si–Mg alloy due to the Si released (equations (35), (36) and (39)), being possible for the precipitation of Mg₂Si according to the following equation:

In addition, the presence of oxides increases melt viscosity and deteriorates the wettability of the system matrix/reinforcement [171]. The reactions to obtain these different oxides include equations (37)–(39), in addition to [170,172]:

2.5.3 Si₃N₄ in Al matrices with Cu and Mg

Figure 10a shows a SEM image with the microstructure of a squeeze cast Al–5Cu/7% Si₃N₄ alloy composite, where Si₃N₄ particles (dark gray) and Al₂Cu phase (light) are clearly differentiated [157]. Besides, AlN (intermediate grey scale) is observed surrounding Si₃N₄, formed according to the reaction of equation (35). Its presence was corroborated through the XRD analysis shown in Figure 10b, where peaks belonging to Al, Si₃N₄, Al₂Cu, and AlN can be identified. This or other interfaces can be formed depending on the presence of Mg. Jeong et al. [173] studied the mechanical properties and interfacial characteristics of Al–Cu and Al–Cu–Mg alloys reinforced with Si₃N₄. They found that the composites with Al–Cu–Mg matrices presented significantly higher elongations than Al–Cu composites, and this is due to the absence of cavities at the interfacial regions and a stronger interface for Al–Cu–Mg matrices. These authors used TEM to analyze the interfaces Al/Si₃N₄ after tensile deformation and found significant differences between them. As shown in Figure 10c, Al/Si₃N₄ interface for the composite with Al–Cu matrix presented flat edges without interfacial reaction products. Its analysis using HRTEM showed a thin SiO₂ amorphous layer of about 1 nm covering the interface (arrowed). The formation of this layer suggests no reaction Al/Si₃N₄. Otherwise, Figure 10d shows that interfaces for the composite with Al–Cu–Mg matrix presented curved surfaces and interfacial reaction products, indicating a strong reaction. In this case, the SiO₂ amorphous layer was not reported. Reactions products (arrowed) grew with definite crystallographic relationships to Si₃N₄ and were discontinuously located surrounding Si₃N₄. EDS and crystallographic analyses showed that the reaction phase was a fcc. Al-based solid solution was supersaturated by Si and Mg. These reaction phases grew with definite crystallographic relationships to the Si₃N₄ crystal, which has a cubic lattice (a = 0.779 nm), whit a lattice mismatch Al/Si₃N₄ of 0.015 (by counting 1/2 for Si₃N₄), because for Al, a = 0.405 nm. This interface relaxed the stress concentration and suppressed the extensive development of microcracks and cavities. These results show the importance of Mg in these composites. HRTEM images were essential to analyze these crystallographic characteristics of the interfaces and revealed their effect on the mechanical properties of the composites.

Figure 10

(a) SEM image of a squeeze cast Al–5 Cu/7% Si₃N₄ composite and (b) XRD analysis of this composite with the presence of different peaks. (Reproduced with permission from [157].) (c and d) TEM and HRTEM images of Si₃N₄ particles for tensile deformed. (c) Al–Cu/Si₃N₄ composite and (d) Al–Cu–Mg/Si₃N₄ composite (Reproduced with permission from [173].).

2.5.4 Si₃N₄ in Al–Si–Cu–Mg matrices

Feng et al. [174] studied an Al–11.5Si–1.3Mg–1.3Cu–1.3Ni (max. wt%) alloy reinforced with 20% of Si₃N₄ whiskers, manufactured by squeeze casting, and found that the addition of the reinforcement modified α-Al morphology from dendritic to equiaxed. They also reported the modification of the eutectic Al–Si from netlike to particulate. In addition, MgO and MgAl₂O₄ were obtained as reaction products. These authors found that the increase in the manufacturing temperature led to stronger interfaces due to the segregation of solute atoms near this zone.

Chen et al. [175] reported that the addition of 45% Si₃N₄ accelerated the precipitation process of Al₂MgCu for an Al 2024 alloy (4.79 Cu–1.49 Mg–0.168 Si), although the precipitation sequence was not altered. Qu et al. [171] also reported a more accelerated kinetics for Mg₂Si precipitation, for an Al–11.5 Si–1.0 Mg–0.5 Cu–0.5 Ni alloy reinforced with 20 vol% of Si₃N₄ whiskers.

As for the other nonoxide reinforcements analyzed earlier, the study of Mg outstands as the alloying element for Al matrices reinforced with Si₃N₄, increasing wettability and forming compounds such as MgO and spinel MgAl₂O₄. Although some works analyze Si-rich matrices, the effect of this element on the equilibrium needs more attention. Otherwise, the presence of Cu as the reaction product was not found.

2.6.1 CNTs

CNTs have low density and high strength-to-weight ratio and elevates ductility and corrosion resistance. They are used for reinforcing AMC in applications such as gears, cylinder liners, and sport utilities [17]. Polymer composites reinforced with CNT have been widely studied [179,180], but research related to the use of CNT reinforcing aluminum and its alloys is limited [19,43,44,45]. Among the topics needing attention for these composites are the interfacial bonding CNT/Al and the effect of the alloying elements on other interactions. It has been reported that atoms in graphitic structures in general present strong covalent s-bonds within a basal plane, also forming covalent bonds with metals [176]. In the case of CNTs, there are two different interfaces depending on the contact area of the CNT with the metal: (i) the contact with the ends of the CNT, involving covalent bonds at the interface, and (ii) a side-contact generating a weak bond [176].

2.6.1.2 CNT in Al matrices with Si

Several studies have shown that matrix composition has a significant effect on the reaction products and on the kind of interfacial layer formed. Bakshi et al. [46] reported the formation of two different carbides due to the reaction between CNT and molten metal: Al₄C₃ for an Al–11.6 wt% Si alloy, and SiC for an Al–23 wt% Si alloy. For Al–Si alloys, SiC is formed through the reaction:

(43) Si (l)  + C (s)  = SiC (s)

The mechanism of growth of SiC is similar to the formation of Al₄C₃ and begins with C atoms reacting with Si, forming a very thin SiC layer [43]. Its growth also includes lateral growth on the CNT surface and perpendicular growth to the formed carbide layer. The formation of Al₄C₃ or SiC depends on both temperature and Si content, as it is reported in different works [43,46,184]. Bakshi et al. [46] reported that for a 11.6 wt% Si Al alloy, the critical carbide thickness for the formation of Al₄C₃ is lower than for the formation of SiC, while it is inverse for a 23 wt% Si alloy. They presented a graph with the equilibrium fraction of Al₄C₃/(Al₄C₃ + SiC) vs Si wt% in the alloy, at various temperatures, for composites reinforced with 10 wt% of CNTs. This diagram can be used to predict the interfacial carbide formed according to alloy matrix at a given processing temperature. Although the formation of small amounts of Al₄C₃ is helpful in stress transfer to CNTs by pinning the CNT to the matrix [185,186], it is preferred the formation of SiC. Critical layer thickness for SiC formation is lower than that for Al₄C₃ formation, which leads to a better wettability of molten Al-Si matrix with CNT. Al/SiC has a higher work of adhesion, indicating a lower contact angle for wetting. Besides, interfacial strength is higher when SiC is formed, presenting a more efficient load transfer for Al₄C_3, preventing delamination [43]. Figure 11a and b shows high-resolution TEM images for the analysis of the formation of these carbides [46]. As shown in Figure 11a, a thin (2–5 nm) layer of SiC is present on the CNT surface, for an Al–23 wt% Si alloy reinforced with 10 wt% CNTs. Otherwise, the formation of Al₄C₃ occurred for a matrix with lower Si content, as shown in Figure 11b for an Al–11.6 wt% Si alloy. Crystallographic planes corresponding to Al, Si, and Al₄C₃ can be identified in Figure 11b. It has been reported a defined orientation relationship between Al₄C₃ and C in other C materials, but this does not occur for CNTs due to the structural configuration of the carbide. This Al₄C₃ layer has a rhombohedral structure, with alternate Al₂C and Al₂C₂ layers, with a = 0.3335 nm, b = 0.3335 nm, and c = 0.85422 nm. Thus, Al₄C₃/CNT interface is expected to be strained due to the higher C–C distance in Al₄C₃ (3.33 Å) compared to this distance in CNT (1.42 Å) [46].

Figure 11

High-resolution TEM images of (a) an Al–23 wt% Si alloy reinforced with 10 wt% CNTs (inset shows the XPS validating SiC) and (b) an Al–11.6 wt% Si alloy containing 10 wt% CNT (Reproduced with permission from [46].).

2.6.2 Other carbon material reinforcements

CNTs are the most used C material for reinforcing Al, but there are few works about the use of graphene and graphite. That is why we are including a brief analysis of our search for these reinforcements.

Graphene has high mechanical, electrical, and thermal properties, reinforcing AMC and leading to better functional and mechanical properties. Graphene/Al composites have applications in catalysis, electronics, energy storage and conversion, sensing, and biotechnology [48]. Although this composite is more feasibly manufactured by powder metallurgy, avoiding interfacial reactions [20], Chak and Chattopadhyay [49,50] reported the use of stirring-assisted liquid route for manufacturing AMC of graphene nanoplatelets reinforcing an Al matrix. Graphene addition acted as grain refiner and improved mechanical and wear properties, although agglomeration occurred for graphene wt% >0.5. The bonds between Al, Cu, and other metals with graphene are weak Van der Waals forces, leading to the agglomeration of the reinforcements. These agglomerated layers form thick sheets and generate interfacial defects. Besides, Al₄C₃ may be obtained [196], decreasing the composite properties. This reinforcement has been less studied than other C materials and needs more attention. No works were found related to the formation of other carbides or reaction products due to the addition of Si, Cu, or Mg to Al matrices reinforced with graphene.

On the other hand, graphite has low density, high thermal conductivity, and good mechanical strength, reinforcing Al alloys in applications such as heat sinks, brake discs, and pistons [17]. The interaction of Al with graphite also origins the formation of Al₄C₃. This reaction is severe at temperatures higher than 627°C, degrading the reinforcement and leading to a decrease in strength [17,196]. Liu et al. [196] found that the addition of small amounts of Mg prevents the formation of graphite clusters, leading to composites with improved physical and mechanical properties. They obtained a composite with graphite particles uniformly distributed in an Al–0.6 wt% Mg matrix, but when Mg content increased to 1 wt%, agglomeration took place. Pelleg et al. [197] investigated the effect of alloying elements on the Al/graphite interface and found that Si presented the highest effectivity as the alloying element for avoiding or reducing the formation of Al₄C₃. They reported that Mg is also efficient in reducing the solubility of C in Al. Although for graphite exists more information than for graphene, this reinforcement also needs attention.

2.7.1 Fe in Al matrices

Some reported reactions when molten Al comes in contact with solid Fe are as follows [58,198,199]:

The Fe–Al system can form six nonstoichiometric intermetallic compounds: AlFe₃, AlFe, Al₂Fe, Al₃Fe₂, Al₅Fe₂, and Al₃Fe [200]. Some of them have been reported for these syntactic foams for different matrices. Lee et al. [199] studied the reinforcement of Al by 10% of Fe particles through the friction stir process, identifying that reaction products were Al₁₃Fe₄ (also denoted as Al₃Fe) and a low quantity of Al₅Fe₂. This last compound was obtained due to the reaction between Al₁₃Fe₄ and Fe. Other Fe–Al and ternary intermetallics have been reported, but mainly for Si-rich matrices, as will be analyzed next.

2.7.2 Fe in Al matrices with Si

The use of Fe reinforcing AMCs has been mainly researched for Al–Si alloys. The same occurs for MMSF, where almost all the studies are related to the infiltration method. Si imparts excellent castability to Al alloys, also leading to the formation of different intermetallics Al–Fe–Si. Ten different ternary intermetallics have been reported for this system [201]. Guo et al. [202] reported an intermetallic layer for an Al–7 Si alloy reinforced with Fe hollow spheres, with three different compositions in the interfacial region: Al₅Fe₂, Al₃Fe, and Al₈Fe₂Si. The increase in the interface thickness led to the fast fracture of the Al/Fe bonding. A brittle-hard needle-like phase (Al₅FeSi) in the Al matrix around Al/Fe interface was the main cause of the lower shear strength. Rabiei and Vendra [59] studied MMSF consisting of low-carbon steel hollow spheres, reinforcing an Al-7Si alloy. They used casting processing, reporting the formation of an Fe₂₅Al₆₀Si₁₅ interfacial layer and of hard Fe₂Al₇Si plates. The formation of the interfacial layer increased the bonding strength between spheres and matrix, improving the mechanical properties of the material. For Al–7Si matrices, Vendra and Rabiei [203] and García-Avila and Rabiei [204] found different intermetallics and the same interfacial layer. This can be observed in the SEM images of Figure 12a and b. As shown in Figure 12a, needle-shaped FeAl₄Si is present in the Al–Si matrix, while Figure 12b shows a continuous interfacial layer of Fe₂₅Al₆₀Si₁₅ surrounding the sphere and an intermediate discontinuous layer of block-like Fe₂Al₇Si particles. Due to the presence of such intermetallic phases, these syntactic foams present brittle fracture. Al and Si interdiffusion from the matrix to the sphere walls have been reported, also diffusing Fe and other alloying elements from the sphere toward the matrix. This leads to the formation of different ternary intermetallics with hexagonal or more complex crystal structures. The crystallographic analysis in the literature of Fe₂₅Al₆₀Si₁₅ (τ10-phase Fe₅Al₁₂Si₃) showed a hexagonal lattice for this intermetallic, with a = 0.753 nm and c = 0.763 nm [205]. Fe has bcc crystal lattice, with a = 0.2866 nm, while as commented earlier for Al a = 0.405 nm. This could lead to poor coherency among Al, Fe, and the intermetallic layer, with few orientation relationships. These characteristics need more attention for these materials, including their study through techniques such as HRTEM, because this and other interfaces are only shown using SEM or OM, and identified using XRD, without a more in-depth analysis [58,202,203].

Figure 12

SEM images of an Al–12Si alloy reinforced with hollow steel spheres, processed by casting. (a) Spheres with their walls, and matrix with FeAl4Si are clearly observed. (b) Sphere/matrix interface with the presence of different phases (Reproduced with permission from [59].).

2.7.3 Fe in Al–Si–Cu–Mg matrices

Bálint and Szlancsik [206] studied syntactic foams formed by matrices Al99.5, Al–12 Si, Al–1 Mg–1 Si, and Al–5 Cu, reinforced with Fe hollow spheres. They found that spheres had no significant impact on the microstructure of the matrix material, but coherent thin interfacial layers were observed for all the matrices. These layers presented strong bonding, which were beneficial for the mechanical properties. Near the spheres, the number of Si-, Cu-, and Mg-rich precipitates increased (for Al–12 Si, Al–5 Cu, and Al–1 Mg–1 Si matrices, respectively), while for composites with Al–5 Cu and Al–1 Mg–1 Si matrices, the sphere walls were enriched with Cu and Mg. Cu was found to be the element that most strengthened the foams, forming Al₂Cu.

Otherwise Orbulov and Ginsztler [207] found a significant effect of the alloying elements on the mechanical properties of MMSF with different matrices, being the best mechanical properties in the ascendant order Al99.5, Al–1 Si–1.2 Mg, Al–12 Si and Al–4.5 Cu. These authors also found that T6 heat treatments led to strength increments much smaller than expected due to the lower precipitation. The cause of this modification in the aging process was the decrease in the alloying elements in the matrix, originated by the interfacial reactions and the diffusion of alloying elements to the spheres.

Due to the hollow interior of the spheres, their interfaces with the matrices represent a high fraction of the solid material of a syntactic foam, as reported by Vendra and Rabiei [208]. They found that for an Al 356 alloy, reinforced with 59 vol% of steel hollow spheres, the volume of the Al–Fe–Si intermetallic layer represented 7% of the whole composite. That is why in this kind of composites, the effect of the interface is more determinant on the mechanical properties.

As can be seen, enough information was found in the literature for the reactions between Fe spheres and Al or Al–Si matrices, with different reaction products. Nevertheless, it was noted that more research is needed about the effect of the manufacturing process on the interfacial reactions. For alloys with the presence of Cu and/or Mg, the information is few and mainly focused on the modification of the precipitation processes due to the addition of the reinforcements. That is why, for the use of Fe as reinforcement, it is also essential to increase research in relation to the addition of these alloying elements and direct or indirect effects on microstructure and mechanical properties. Although Fe–Al intermetallics and interfaces are well known, including their crystallographic analyses, more attention is necessary for their formation and characteristics for the specific case of syntactic foams and their manufacturing processes.