In Situ Al₃BC/Al Composite Fabricated via Solid-Solid Reaction: An Investigation on Microstructure and Mechanical Behavior

Abstract

Al₃BC, with its remarkably high modulus of elasticity (326 GPa) and hardness (14 GPa), coupled with a low density (2.83 g/cc), stands out as a promising reinforcement material for Al matrix composite. To study feasibility of solid-solid reaction (SSR) by forming an in situ Al₃BC reinforcing phase within the matrix, this study developed an Al₃BC/Al composite via mechanical alloying, followed by sintering at 1000 °C/1 h, and subsequent hot pressing at 400 °C/40 MPa. The reaction kinetics and corresponding electron microscopy images suggest that the aluminum (Al)-boron (B) reacts with graphene nanoplates (GNPs) to form both clusters and a heterogeneous multi-structured Al₃BC reinforcements network dispersed within the fine-grain (FG) Al matrix. The heterostructure contributes to a good balance between strength (~284 MPa) and ductility (~17%) and stiffness (~212 GPa). Superior strain hardening ability (n = 0.3515) endorses remarkable load-bearing capacity (σ_c = 1.63) and thereby promotes excellent strength-ductility synergy in the composite. The fracture morphology reveals that reasonable ductility primarily relies on the crack deflection by the FG-Al matrix, playing a critical role in delaying fracture. The potential importance of the matrix microstructure in the overall fracture resistance of the composite has been highlighted.

1. Introduction

Al-matrix composites (AMCs) offer lightweight solutions for the transport field, enabling significant progress in energy conservation and emission reduction. In addition, they boast a superior strength-to-weight ratio, a key property deeply desired in aerospace and military applications [1,2,3]. However, the challenge lies in balancing strength with ductility, despite the lightweight and durable nature of the Al matrix, enabling a serious threat to their adaptation in those applications [4]. While strength may increase with decreasing the grain size in the matrix in many nanostructured materials, at the same time it costs reduced ductility [5]. The limited ductility associated with nanostructured grain structures leads to a reduction in composite strain hardening capacity. On this occasion, Al-matrix composites with particulate reinforcement can typically provide several key advantages of superior specific strength and stiffness [6], enhanced wear resistance [7], and good structural integrity with property design flexibility [8]. Ceramic reinforcements with high modulus and strength, such as silicon carbide (SiC) [9], alumina (Al₂O₃) [10], boron carbide (B₄C) [11], titanium boride (TiB₂) [12], aluminum nitride (AlN) [13,14], zirconium boride (ZrB₂) [15], titanium aluminide (Al₃Ti) [16], magnesium silicide (Mg₂Si) [17], etc., have been beneficial in enhancing the room temperature properties and softening resistance of these composites. According to research Al₃BC (an Al compound)-reinforced AMCs (Al₃BC/Al composites) show impressive properties: high modulus (100 GPa), strength (400 MPa), and reasonable ductility (6%) [18,19]. In addition, excellent Al₃BC-Al matrix interfacial bonding promotes the development of a heterogeneous grain structure and boosts strengthening in the Al₃BC/Al composites [20,21]. From a mechanical point of view, Al₃BC has emerged as one of the potential candidates for reinforcing Al-matrix composites.

Al₃BC is a promising reinforcement for the Al-matrix composites due to several advantages. Firstly, reinforcing Al₃BC is significantly easier than conventional reinforcements. Additionally, it requires only minimal additions of B and C to obtain a high fraction of reinforcing Al₃BC in the Al-B-C systems, even among other aluminum–boron carbide phases, e.g., Al₃BC₃, Al₈B₄C₇, AlB₂₄C₄, etc. [22]. Secondly, Al₃BC particles act as good heterogeneous nucleating sites for the matrix, as demonstrated by Ma et al. in Mg-ZA263 composites [22]. Notably, the unique distinctive crystal structure of Al₃BC, rich in Al, offers an excellent combination of high hardness, like ceramic reinforcements, yet with a more metallic behavior [21]. Furthermore, the combination of properties of the Al₃BC phase, such as low density (2.83 g/cc) compared to other reinforcements, e.g., Al₂O₃ (3.98 g/cc) and TiB₂ (4.50 g/cc), excellent thermal stability (up to 1700 K under 1.6–4.8 GPa pressure), high hardness (14 GPa), and remarkably high modulus (Young’s modulus 326 GPa and shear modulus 140 GPa), makes it a promising candidate for reinforcing Al-matrix composites [21,22]. Zhao et al. demonstrated that Al-matrix composites reinforced with Al₃BC particles show noticeably high room-temperature strength (775 MPa) and elastic modulus. Additionally, Al/Al₃BC/25p composites, prepared via liquid–solid reaction (stir-casting) possess superior strain hardening ability [21]. Growing with such promise, Tian et al. and Zhao et al. explored remarkable strengthening in uniformly dispersed nano-Al₃BC-reinforced 6061 Al-alloy and AZ63 Mg-alloys [23,24]. Due to their strength, machinability, and sustainable manufacturing demands, Al₃BC/Al composites show promise for future-generation aerospace and automotive parts (airframes, propeller blades) [25].

Reinforcing graphene platelets (GNPs) uniformly in the Al matrix has been a decades-long challenge, which limits the development of graphene-reinforced high-strength Al-matrix composites. Recently, research by Mei et al. successfully demonstrated the fabrication of Al₃BC/Al composites using a liquid reaction (stir casting) from amorphous B and GNPs [20]. This approach can potentially overcome the hurdle of graphene dispersion and pave the way a new generation of high-performance Al-matrix composites [26]. Powder metallurgy (P/M) offers a solid-solid reaction to create a bimodal heterostructure in the Al-matrix composites, using ball milling and high-press sintering to achieve homogeneous dispersion of reinforcements in the matrix [27,28,29,30,31,32,33]. The key advantage of solid–solid (pressing and sintering) reactions is the ability to control the morphology, shape, and size of the reinforcing phases that evolve during both in situ and ex situ approaches [17,34,35,36]. However, the solid-solid reaction path to develop in situ Al₃BC/Al composites remain a challenge. Kubota et al. showed difficulty controlling microstructures with Al-MgB₂ powder reactions because multiple by-products, like AlB₂ and Al₄C₃, often resulted [37]. The presence of intermediate products, like micron-sized Al₄C₃, can significantly deteriorate the mechanical properties of Al₃BC/Al composites [38]. Therefore, fabrication temperature and modulating the reaction products of Al₃BC/Al composites are crucial to obtain the desired mechanical properties in these composites because of the effect of undesired reaction products. Accordingly, in this work, we focused on a straightforward method for fabricating Al₃BC/Al composites using mechanical alloying (ball milling) combined with a solid-state reaction technique (pressing and sintering) by varying reaction time (3 to 7 h) at 1000 °C temperature to avoid unwanted byproducts with optimized mechanical properties.

2. Experimental Details

2.1. Materials

The investigated Al₃BC/Al composite was fabricated via ball milling and vacuum sintering (1000 °C/1 h) to initiate solid–solid reaction of constituent elements of commercial Al powder (particle size 250–450 μm, HiMedia ^TM, Kennett Square, PA, USA), GNPs (average diameter 1.65 nm, Alfa Aesar ^TM, Haverhill, MA, USA) aggregates, and amorphous B (Sigma Aldrich ^TM, St. Louis, MO, USA). B and GNP powders (2:1 molar ratio) were mechanically alloyed in a high-energy ball mill (Retsch^TM PM 400-MA, Haan, Germany).

2.2. Processing

Milling was carried out for 2 h at a speed of 200 rpm in an Argon atmosphere to ensure uniform elemental distribution. The ball-milling parameters have been reported elsewhere [27]. The milled B+GNPs powder was mixed with Al powder (overall molar ratio 1:2:4) for an additional 5 h using a ball mill and subsequently pre-heated (overnight at 80 °C) and green compacted (60 vol.%). To prepare the corresponding Al₃BC/Al composite, the green compact was then sintered in a vacuum sintering furnace (tubular furnace) at 1000 °C. The sintering process maintained a slower heating and pressurization process, e.g., heated at a rate of 5 °C/min, kept for 1 h (soaking time) below 5 × 10⁻² mbar pressure, and then slowly cooled to room temperature. This ensures Al will maintain formability favorably and even become liquid at 680 °C. Finally, a low-pressure (40 MPa) and high-temperature (400 °C) hot-press deformation was applied for 20 min in a graphite mold (φ 20 mm) to achieve the final consolidated specimens. During hot pressing, the heating rate was maintained at 25 °C/min, followed by water cooling. The experimental procedure employed in this research is shown schematically in Figure 1.

2.3. Characterization

The composite samples were subjected to standard metallurgical procedures, including grinding and polishing, and subsequently cut (as cylindrical compression test specimens) using wire EDM following standard procedure. The overall phase composition was evaluated with X-ray diffraction (XRD) using Rigaku^TM DMAX X-ray diffractometer [operating at 40 kV-40 mA and Cu-Kα radiation] (Tokyo, Japan). The reactions during the composite formation were investigated using differential scanning calorimetry (DSC). Microstructure and composition area mapping were then investigated using an electron microscope (SEM) equipped with energy-dispersive X-ray spectroscopy (EDS). The SEM device fitted with an electron backscattered diffraction (EBSD) detector was used to obtain inverse pole figure maps (IPF) with step size 0.5 μm by an EDAX Trident^TM system (Sparks Glencoe, MD, USA) along with system-integrated FEI Scios™ field emission gun (FEG) scanning electron microscope (FE-SEM) (Sunnyvale, CA, USA). The EBSD data were processed using an EDAX OIM Analysis™ software (https://www.edax.com.cn/products/ebsd/oim-analysis, accessed on 4 May 2025) and a detailed description of the technique has been reported elsewhere [27]. The hardness was measured using a Vickers hardness tester at a load of 100 g [ASTM E10-14 standard [39]]. The room temperature (RT) compression test (strain rate ($\dot{\epsilon }$) of 0.5 $\times$ 10⁻³ s⁻¹) specimens with dimensions 20 × 5.6 × 5 mm³ were evaluated using a servo-hydraulic UTM (100 kN load cell) [ASTM-E08 standard [40]]. The tests continued until the failure of the specimen. The fracture surfaces of the post-compressed specimens were examined using SEM to understand the mode of failure.

3. Results

3.1. Physical and Structrual Characterization

Figure 2a presents the DSC traces of the ball-milled Al and B+GNP powders conducted at a heating rate of 5 °C/min under an argon atmosphere. An exothermic peak is observed beyond the melting point of Al (660 °C) around 680 °C [41], signifying the possible chemical reaction between Al, B, and GNPs initiating solid-solid reactions [21]. No (endothermic/exothermic) peaks are observed beyond 1000 °C. This suggests that there is a simultaneous reaction of Al, B, and GNPs and the in-situ synthesis of Al₃BC. The XRD patterns (Figure 2b) of (B+GNPs)/Al milled powders reveal the presence of crystalline fcc-Al phase [space group: Fm-3m; PDF-2-2023#00-004-0787]. The peak broadening with increasing milling time indicates the possible structural refinement. Furthermore, a decrease in intensity of the fcc-Al peak (111) (w.r.t. (200) peak) demonstrates the dissolution of the fcc phase, which indicates successful alloying [42]. Figure 2c shows the XRD pattern of the as-fabricated Al₃BC/Al composite. The diffraction peaks show the presence of fcc-Al, hcp-Al₃BC [space group: P63/mmc; PDF-2-2023#00-047-1628], and hcp-AlB₂ [space group: P6/mmm; PDF-2-2023#03-65-3381] phases. Importantly, no unwanted Al₄C₃, a common by-product during solid-liquid or liquid-liquid state reaction(s), was detected. In addition, a higher intensity of Al and Al₃BC phases than AlB₂ phases indicates their dominant behavior. The lattice parameters (a), and the relative amounts of the phases in terms of mass fraction/phase volume fractions, V_f, were determined using the XRD peak profile analysis (refer to

Table 1. The summary of phase formation along lattice parameters, a; wt./mass fraction (%), $V_f$; area fraction, $A_f$; Young’s modulus; and hardness of as-evolved phases.

Phase	Vf (wt.%)	Lattice Parameter (Å)	Hardness (GPa)	Modulus (GPa)	Af (%)
Al3BC	65 ± 2	(a) 3.488 ± 0.003 (c) 11.535 ± 0.003	22 ± 7	326 [20]	52 ± 4
fcc-Al	25 ± 2	(a) 4.098 ± 0.001	3 ± 1	52	31 ± 2
AlB2	10 ± 1	(a) 2.978 ± 0.001 (c) 3.259 ± 0.001	2 ± 0.5	250	17 ± 3

for details) [42]. It shows the volume fraction of the Al matrix ($V_f^{Al}$) accounts for 25 ± 2 wt.%, primary reinforcement Al₃BC ($V_f^{{Al}_{3}BC}$) corresponds to 65 ± 2 wt.%, and AlB₂ ($V_f^{Al{B}_2}$) volume fraction is calculated to be 10 ± 1 wt.%. This suggests that the addition of GNPs and B promotes the Al₃BC formation preferentially reacting with Al over forming AlB₂.

3.2. Microstructure Evolution and Reinforcement Distribution

Figure 3a shows a typical composite microstructure of the as-fabricated Al₃BC/Al specimen with heterogeneous grain structure comprised of both coarse (CGs) and fine grains (FGs). The microstructure reveals a deformed morphology with the presence of the following phases: a denser and wider size distribution of elongated shaped (60–70 µm) agglomerates (clusters) of grey Al₃BC, darker contrast AlB₂ phases, and a string/or network of submicron grey Al₃BC phases dispersed within a highly deformed (brighter contrast) fcc-Al matrix (Figure 3b). The EDS element overlay mapping in Figure 3c identifies the following elements: B, C, and Al. The EDS results confirm the presence of the desired in situ Al₃BC phase (yellow circle), and the matrix is primarily composed of fcc-Al. Notably, network-like Al₃BC structure comprises both CG (elongated) and FG (equiaxed) grains and clusters of black, rod-like, micron-size AlB₂ phases located at the interfaces of Al₃BC agglomerates. However, obtaining a clear element overlay for the AlB₂ phase was challenging due to its brittle nature.

The standard metallographic preparation employed likely fractured the clusters, leading to the presence of inclusions. Similar morphology can be found in other relevant studies [38]. Therefore, the overall microstructure reveals a heterogeneous microstructure where in situ Al₃BC phase with varying grain sizes, from coarse-grained Al₃BC to sub-micron fine-grained Al₃BC, dispersed within fcc-Al matrix [43]. The solid-solid reaction at 1000 °C likely accelerated the agglomeration, promoting the formation of micron-scale Al₃BC/AlB₂ clusters within the composite [44]. The phase area fraction (A_f) was quantitatively measured (including more than 5000 particles) using the ImageJ software (https://imagej.net/ij/). The average A_f values were found to be $A_f^{{Al}_2{O}_3}$ = 52 ± 4%, $A_f^{Al}$ = 31 ± 2%, and $A_f^{Al{B}_2}$ = 17 ± 3%, respectively. Such heterogeneity in microstructure is in good agreement with the XRD results [$V_f^{{Al}_{3}BC}$ = ~65 wt.%, $V_f^{Al}$ = ~25 wt.%, and $V_f^{Al{B}_2}$ = ~10 wt.%] (Table 1). The presence of micro-voids (Figure 3b), cracks (Figure 3a), and fracturing in some CG-Al₃BC phases can be observed throughout the microstructure, possibly due to the low pressure and high temperature (400 °C/40 MPa) employed during deformation (hot press).

3.3. Grain Structure Evolution

Figure 4 shows the phase fraction, orientation, and inverse pole figure maps (IPF) obtained from the EBSD analysis of the corresponding SEM micrograph shown in Figure 4a. Figure 4b reveals that the matrix consists of near equiaxed FG-Al (green), FG-Al₃BC (red), and elongated CG-Al₃BC (red) phases that arrange themselves heterogeneously throughout the microstructure. Such microstructural formation is likely due to ongoing deformation, showing a heterogeneous structure. Here, grain size, smaller than 2 µm is defined as FG. FG-Al occupies most of the visual field area and becomes the matrix phase, leading to 41% of the overall composition being FG-Al, with around 19% FG/CG-Al₃BC phases being present within the FG-Al matrix. Additionally, the darker regions, along with the red areas in the map, likely correspond to the unindexed AlB₂ phase. Such heterogeneous microstructures indicates the possible dynamic recrystallization, which usually happens during thermal processing processes [45,46,47]. Comparing the orientation map with the phase map, the recrystallized fcc-Al (FG-Al) occupying the FG regions, and the CG regions occupied by the deformed Al₃BC phase can be observed.

According to the IPF image, the orientation of FG-Al (green), FG-Al₃BC (red), and elongated CG-Al₃BC (red) phases is random and the misorientation between them is high-angle. Similar phenomena were observed in Al₃BC/6061 Al composites, where a dynamic recrystallization-induced fine-grain structure was observed in the Al matrix, which is likely due to dislocation accumulation via plastic deformation by extrusion [18,19]. Therefore, the present results suggest that hot pressing significantly initiates dislocation accumulation and local plastic deformation (SGB deformation) in fcc-Al grains. Further deformation refines the Al grains into near-equiaxed FG-Al grains and brittle agglomerated Al₃BC reinforcements into FG-Al₃BC within the FG-Al matrix, thereby leading to a heterogeneous structure within the composite. Interestingly, during thermal processing, such heterogeneous hard FG-Al₃BC and FG-Al arrangements are likely to limit/hinder migration of grain boundaries and promote mechanical incompatibility at regional boundaries between FG-FGs and FG-CGs, resulting in a strain gradient.

3.4. Evolution of Mechanical Properties

3.4.1. Compressive Response

The room temperature compression test was used to calculate the Young’s modulus (E), yield strength (σ_y), ultimate compressive stress (σ_s), compressive plastic strain (ε_s), and strain hardening capacity (σ_s-σ_y) of the composite (

Table 2. Mechanical properties of the composites and their comparison with published reports.

Alloy	σy (MPa)	σs (MPa)	εs (%)	H (GPa)	E (GPa)	σc	R	Ref.
Al3BC/Al Sintering + Hot pressing	108	284	17	15 ± 5	212	1.63	5.6	Present study
Ti/1TiB2	750	856	10	-	128	0.14	-	[12]
Ti/4TiB2	576	689	27	-	103	0.20	-	[12]
Al/15Mg2Si Sintering + Hot extrusion	114	200	7	-	75	0.75	-	[17]
Al/20SiC Sintering + Hot extrusion	64.6	127	7	-	-	1.00	-	[51]
Al/20Al2O3 SPS	-	96	-	-	-	-	-	[10]
Al/AlNp RT	179	251	-	14	-	0.40	-	[13]

). The compression engineering stress-strain curve displayed in Figure 5a exhibits Young’s modulus (E) of 212 GPa, a uniform deformation with σ_y of 108 MPa and σ_s of 284 MPa without sacrificing plasticity (ε_s) of 17%. A close inspection shows a significant strain hardening capacity (σ_s-σ_y) of up to 176 MPa in the composite, enabling good load-bearing capacity during plastic deformation and delaying the failure. In addition, serrated plastic flow (Luders’ strain) can be observed in the strain-hardened regions (marked in Figure 5a). Interestingly, such observed capacity (σ_s-σ_y) is higher than the values observed for TiB₂ (1–4 vol.%)-reinforced Ti matrix composites showing (σ_s-σ_y) = 105–113 MPa [48]. The work hardening capacity (σ_c) was measured as (σ_c = (σ_s/σ_y) − 1) 1.63, indicating better load-bearing ability in the composite than other particulate reinforced composites (Table 2) [49,50]. Furthermore, σ_c is strongly correlated to its yield strength. This relationship is governed by the Hall-Petch relationship, which considers both grain size and dislocation density contrast factor between adjacent grains [49]. Therefore, it appears that an increase in the grain size, i.e., hard CG/clusters reinforcements, can significantly lower the σ_y (easier to deform initially) but has a greater potential to increase the strain hardening capacity σ_c in Al₃BC/Al composites [50]. A close-up of the engineering stress–strain diagram reveals a serrated plastic flow, which indicates potential key occurrences of the formation of high-density dislocations, and initiating dislocation plasticity events [23,24].

3.4.2. Pronounced Strain Hardening Behavior

Strain hardening behavior of the composite was analyzed and plotted in Figure 5b by measuring the strain hardening exponent (n) using Hollomon equations:

$$\sigma =K{\epsilon }^n,$$

(1)

$$n=d\left(ln\sigma \right)/d\left(ln\epsilon \right),$$

(2)

where σ is stress, ε is strain, and K is a constant [52]. The plot shows strain hardening behavior depending on the microstructure of the composite. It reveals two stages of strain hardening, e.g., an initial high rate due to hard reinforcements (n₁ = 0.4112) and a later more homogeneous deformation due to particulate matrix involvement interaction (n₂ = 0.3048), with an overall strain hardening rate observed as 0.3515. It is reasonable that particulate matrix involvement interaction (n₂ = 0.3048) influences the serrated plastic flow and promotes uniform deformation and superior ductility by delaying localized failure.

Serrated plastic flow appears either due to displacive phase transformation or the initiation of dislocation plasticity events [53,54]. As the XRD ruled out the possibilities of displacive phase transformation, dislocation plasticity events likely govern the mode of plastic deformation in Al₃BC/Al composite. Furthermore, such serrated flow resembles Type-B serrations, which seemingly pronounce dislocation interaction involvement during SGB deformation [53,54]. Vickers hardness was measured to determine the individual strength of the phases and their influence on the overall strength of the composites (Table 2). It shows significantly higher hardness for Al₃BC (CG) of H_Al3BC = 22 ± 7 GPa, compared to FG-Al matrix, H_Al = 3 ± 1 GPa, and AlB₂, H_AlB2 = 2 ± 0.5 GPa. The measured H_Al3BC values are identical to the measured hardness values of Al₃BC from other published reports, e.g., H_Al3BC = 24 ± 1 [55], ~26 [21], respectively. Surprisingly, the hardness of the FG-Al matrix within the composite surpasses that of conventional Al (~1 GPa) [18,19,20]. Voigt’s model predicted the theoretical Young’s modulus (E_Th) of the composite [56]:

$$E_{Th}=f_p{E}_p+f_m{E}_m,$$

(3)

where f_p and f_m are the volume fraction of the particles and the matrix and E_p and E_m correspond to the Young’s modulus of Al₃BC (326 GPa) [18,19] and FG-Al (52 GPa), as measured using nanoindentation, respectively. The measured Young’s modulus of FG-Al matches the literature value of ~56 GPa [21]. The fitted Young’s modulus E_Th = 224 GPa matches the experimental Young’s modulus of the Al₃BC/Al composite (E = 12 GPa). Furthermore, bulk modulus (G = 80 GPa) and shear modulus (B = 196 GPa) were calculated from the empirical equations [57].

The hardness (H_Th) of the Al₃BC/Al composite was theoretically predicted using the ROM [58]:

$$H_{Th\_up}=f_p{H}_p+f_m{H}_m,$$

(4)

$$H_{Th\_low}=f_p{H}_p+f_m{H}_m^{-1},$$

(5)

where H_p and H_m are the Vickers hardness, and f_p and f_m correspond to the mass fraction of Al₃BC particle and Al matrix (obtained using the XRD patterns), respectively. Notably, the AlB₂ phase was ignored due to low contributions in hardness (H_AlB2 = 2 ± 0.5 GPa) and mass fraction (low $V_f^{Al{B}_2}$ = 10 ± 0.1 wt.%). Such significantly improved mechanical properties surpass those of deformed Al alloys (Table 2).

3.4.3. Strengthening Parameters Measurements

Apparently, strain hardening leads to increased hardness, modulus, and overall strength, while still maintaining good ductility. In this case, the achieved mechanical property may be compared with Al₃BC/Al-Cu composites containing 13% and 26% Al₃BC particles in the Al matrix, which possess high tensile strength (300 MPa–427 MPa) but lower elongations (2–4.5%) [53]. This justifies further investigation of strengthening parameters.

• Load-bearing contribution (Δσ_Load): The strength increment due to load-bearing capacity of Al₃BC reinforcement was determined from the expression [59]:

  $$∆{\sigma }_{Load}={\displaystyle \frac{A+1}{4}}{\upsilon }_P{\sigma }_m,$$

  (6)

  where υ_p denotes the volume fraction of Al₃BC particles (υ_p = 0.65) and σ_m refers to the yield strength of the composite and the particle aspect ratio denoted by A (~1) [18,19]. Hence, the contribution Δσ_Load of Al₃BC reinforcement was predicted to be Δσ_Load = 12 MPa, comparable to Δσ_Load = 16 MPa in the AlSi₁₀Mg/SiC (10 wt.%) matrix composite, exhibiting UTS 350 MPa and elongation to failure up to ~3% [60].
• Strengthening capability contribution (R): The contribution of a specific volume fraction of Al₃BC reinforcements to the strengthening of the Al matrix in Al₃BC/Al composite, R, was evaluated from [61]:

  $$R=({\sigma }_{yc}-{\sigma }_{ym})/f_p{\sigma }_{ym}$$

  (7)

  The strengthening capability of the Al₃BC/Al composites, calculated as R = 5.6 is impressive. The value is close to FeSiB-reinforced Cu matrix composites (R = 5.9) [62] and even surpasses SiC-particle-reinforced (R = 2.5) [63] and Al₂O₃-particle-reinforced (R = 2.3) [48] Cu matrix composites by a factor of two. This finding highlights the remarkable strengthening efficiency achieved in Al₃BC/Al composites.
• Furthermore, Pugh’s ratio, which is the ratio of the calculated bulk modulus (G = 80 GPa) to shear modulus (B = 196 GPa) (B/G), can be used to determine the ductility or brittleness of the material. A Pugh’s ratio greater than 1.75 indicates that the material is more ductile in nature [64]. Based on the calculated Pugh’s ratio, a B/G = 2.45 is observed for the as-developed Al₃BC/Al composite. It can be concluded that the as-fabricated Al₃BC/Al composite has superior ductility while considering the poison’s ratio of υ = 0.32. Furthermore, if the Vickers hardness of a material falls within the range of 10–20 GPa, it is typically referred to as hard [65]. According to this information, the composite (H_Th = 15 ± 5 GPa) is categorized as a hard material.

4. Discussion

4.1. Fracture Surface Investigation

To better understand the strengthening effect of particles and the failure mechanism of the corresponding composite, the fracture surface of the sample after a compression test has been investigated. Figure 6a shows a typical compression fracture surface, which reveals primary cracks following the Al₃BC network, with limited penetration in the matrix. AlB₂ and Al₃BC phases are exposed on the fractured surface because of their low ductility, and cracks appear at interfaces with the matrix. Figure 6b reveals that Al₃BC particles have a hexagonal morphology with a sub-micron size. The broken particles highlighted by arrows further support their greater load-bearing ability, like mechanically milled and hot-extruded Mg₂Si/Al composites [17]. In addition, Figure 6b reveals that the FG-Al phases are tightly surrounded by the Al matrix (white circle). Figure 6c reveals that the traces of dimples (red arrow) significantly suggest the ductile nature of fractures in FG-Al grains. Furthermore, Figure 6c shows the tear ridges (yellow arrow), which indicate localized plastic deformation around fractured particles, further supporting some ductility, along with cleavage planes (white arrow). These observations imply a mixed mode of fracture, which identifies both tough and brittle characteristics that exist in Al₃BC/Al composites.

4.2. Understanding the Strength-Ductility Response in the Composite

The cracked Al₃BC phase (Figure 6b) supports its brittleness (Figure 3c). Therefore, the fracture morphology suggests that the large CG/clusters might be the origin of the failure in these composites. Traces of pull-out (indicating interfacial debonding, suggesting Al₃BC particles might have reached their strength limit) and tear ridges (indicating stress concentration occurring around these particles) (Figure 6b,c) suggest Al₃BC bears significant load during deformation and stress would concentrate around these particles. Brittle FG/CG-Al₃BC (Figure 6c) initiates cracks or deboning from the matrix, which ultimately leads to the formation of primary cracks (Figure 6a). In addition, Figure 6c reveals some micron-sized AlB₂ particles fractured with smooth surfaces, while some particles remain intact but deboned from the matrix, supporting the brittleness claims. Figure 6d shows that while the primary crack propagates along interfaces, a few irregular secondary cracks initiate within the FG-Al matrix and penetrate further, suggesting some resistance to fracture. These secondary cracks (Figure 6d, yellow arrow) are ultimately bridged and deflected by the surrounding FG-Al matrix, delaying the fracture. This highlights the potential role of the FG-Al matrix in enhancing fracture resistance. As the Al₃BC phase has a high-volume fraction in the composite, the inner unit stress facilitates load transfer to the adjacent units (n = 0.4112) to coordinate the deformation, likely contributing to early yielding of as-fabricated Al₃BC/Al composites.

Continued loading promotes crack bridging through superplastic flow, maintaining a global strain hardening (n = 0.3048) and resulting in adequate ductility (17%). A high-volume fraction of brittle and hard Al₃BC ($V_f^{{Al}_{3}BC}$ = 65 ± 0.1 wt.%) transfers load to the adjacent unit to coordinate the deformation, ultimately improving strength and ductility. However, a very high amount of Al₃BC ($V_f^{{Al}_{3}BC}$ = 65 ± 0.1 wt.%) leads to an increase in the probability of primary crack formation and causes the early fracture in the compression test. Such a distinctive microstructure arrangement is beneficial to strengthening the composite, which is favorable to the strength and ductility of the composite.

5. Conclusions

In summary, the as-fabricated Al₃BC/Al composites involving solid-state reactions exhibit improved mechanical properties. These reaction kinetics suggest that that prolonged milling duration (7 h) accelerates the solid–solid reaction to initiate in the modified alloy. Structure and microstructure investigations reveal that hot pressing (400 °C/40 MPa) effectively promotes a multi-scale heterostructure in the composite via size reduction by breaking down and dispersing agglomerates within the composite from coarse-grain CG-Al₃BC to sub-micron FG-Al₃BC dispersed within the FG-Al matrix. Strong interfaces between the FG-Al matrices promote crack bridging deflection through the plastic flow, reducing crack propagation, delaying fracture, and ultimately maintaining adequate ductility.