Effects of In Situ TiB₂ on the Microstructural Evolution, Mechanical Properties, and Friction Behavior of the Al-Si-Cu Alloys Processed by Laser Powder-Bed Fusion

Abstract

In the present study, the densification behavior, microstructural evolution, mechanical properties, and friction behavior of a TiB₂/Al8SiCu composite and Al8SiCu alloy manufactured by laser powder-bed fusion (PBF-LB) were systematically investigated. The results confirm that the addition of in situ TiB₂ particles into Al8SiCu alloys reduce the volumetric energy density required for a high-density TiB₂/Al8SiCu composite. The TiB₂ particles promoted a transformation of columnar to equiaxed crystals and the formation of high-angle grain boundaries. The grains on the vertical direction of the PBF-LBed TiB₂/Al8SiCu composite were much finer than those of the PBF-LBed Al8SiCu alloy. The addition of TiB₂ promoted the grain refinement of the Al8SiCu alloy, of which the average grain size decreased from 15.31 μm to 7.34 μm. The yield strength (YS), ultimate tensile strength (UTS), and elongation (El) of the PBF-LBed Al8SiCu alloy were 296 ± 6 MPa, 517 ± 6 MPa, and 11.7 ± 1.0%, respectively. The PBF-LBed TiB₂/Al8SiCu composite achieved a balance between strength and ductility with a yield strength of 328 ± 8 MPa, an ultimate tensile strength of 541 ± 3 MPa, and an elongation of 9.1 ± 0.7%. The increase in strength mainly resulted from grain boundary strengthening, dislocation strengthening, load-bearing strengthening, solid-solution strengthening, and Orowan strengthening, of which the dislocation strengthening and Orowan strengthening were critical. The enhanced hardness associated with the grain refinement and the formation of the in situ TiB₂ particles also led to an enhanced tribological performance, of which reductions in the average friction coefficient from 0.655 to 0.580 and wear rate from 1.76 × 10⁻³ mm³/Nm to 1.38 × 10⁻³ mm³/Nm were found.

1. Introduction

Additive manufacturing (AM) is a widely studied method that builds objects from the bottom up by adding materials in layers. Unlike traditional subtractive manufacturing, AM allows the direct creation of parts with intricate geometries, streamlining the production process, reducing cycle times, and cutting costs [1,2]. It also enables the production of light components without compromising strength. Laser powder-bed fusion (PBF-LB) is a key method in additive manufacturing that employs a high-energy laser beam to selectively melt metal powder layer-by-layer to create parts under a high cooling rate [3,4]. PBF-LB offers extensive design freedom and precise fabrication, making it pivotal in aerospace and automobile manufacturing [5]. Currently, PBF-LB is predominantly utilized for shaping various metal materials, including high/medium-entropy alloys [6,7], stainless steel [8], etc. Aluminum (Al) alloys are preferred materials for PBF-LBed lightweight structures due to their excellent balance, low cost and density, excellent antioxidant properties, superior corrosion resistance, and effective thermal conductivity [9,10,11].

Cast Al-Si alloys exhibit excellent fluidity in a liquid state and demonstrate relatively small shrinkage during solidification, making them very suitable for PBF-LB manufacturing [12,13]. However, the moderate mechanical properties of Al-Si alloys limit the industrial applications in an extended range. Previous studies [14,15] have shown that the addition of Cu into Al alloys exerts a positive influence in strength enhancement. For example, the hardness of PBF-LBed AlSi10Mg with Cu addition reached a hardness level that the PBF-LBed Al7075 alloy obtained after T6 treatment. It has been suggested that the incorporation of 4 wt.% Cu in PBF-LBed AlSi10Mg alloy improves the nucleation rate of Al-Si eutectics and fosters the formation of the nano-Si phase during direct ageing, which achieves the ultrahigh yield strength of 522 MPa [15]. Apart from the increased solubility of Cu when induced by a high cooling rate at 10³–10⁶ °C/s, the non-equilibrium solidification process also promotes the formation of precipitation phases and builds the foundation of strengthening in Al-Si-Cu alloys with subsequent heat treatments [16,17].

Another approach to enhance the mechanical properties of PBF-LBed Al-Si-based alloys is to use composites by adding reinforcement particles. It has been found that Al-based composites can obtain excellent balance in good plasticity, good thermal conductivity, high stiffness, and excellent wear resistance [18,19]. The available and popular ceramic particles include Al₂O₃ [20], TiC [21], and TiB₂ [22]. Of these, TiB₂ offers extraordinary properties in combination, including (i) a high melting point that allows TiB₂ to remain stable in a melt pool [23], (ii) a high laser absorption rate and hardness that can enhance the formability and wear resistance of Al-based composites [24], and (iii) the coherent interface between α-Al matrix and TiB₂ particles that promotes grain refinement and improves mechanical properties [25]. It has been found that the optimal proportion of TiB₂ into an AlSi10Mg alloy can effectively inhibit crack propagation and acquire a superior combination of strength and ductility via eliminating the texture [26]. With diminishing metallurgical defects and improving the nucleation of α-Al grains, excellent mechanical properties (e.g., UTS of 390 ± 6.5 MPa and El of 14.8 ± 1.8%) could be obtained in the PBF-LBed TiB₂/Al-Cu-Mg-Ag composites [27]. Also, the hardness and wear resistance of a PBF-LBed Al7075 alloy are significantly increased when the content of TiB₂ particles increases continuously [28]. Also, the vertical section of a PBF-LBed TiB₂/Al-12Si composite consists of fine equiaxed grains [29]. TiB₂ particles can refine the coarse columnar grains and transform columnar grains into equiaxed grains in a PBF-LBed Al-12Si composite.

Previous work [30,31,32] has mainly concentrated on the effects of TiB₂ particles on Al-Cu and AlSi10Mg alloys. The studies include the role of TiB₂ particles in promoting the heterogeneous nucleation of α-Al grains and the impact of TiB₂ on the mechanical properties and strengthening effects of aluminum alloys processed by PBF-LB. However, the work on the influence of TiB₂ particles on the microstructure and mechanical properties of PBF-LBed Al-Si-Cu alloys is still limited even though Cu is critical in Al-Si alloys for improved mechanical properties. Therefore, the current research seeks to explore the effect of in situ TiB₂ particles on Al-Si-Cu alloys. The composite was made via the salt–metal reaction method. The microstructure and the mechanical and tribological properties of the PBF-LBed TiB₂/Al8SiCu alloy were thoroughly investigated. The discussion was concentrated on the strengthening mechanisms and wear mechanisms.

2. Materials and Experimental Procedures

2.1. Powder and Material Fabrication

The introduction of in situ TiB₂ nanometer-sized particles to Al-Si-Cu alloy was achieved using the mixed salt method with K₂TiF₆ and KBF₄ as reaction salts. The pre-alloyed ingots were supplied by industry. Gas atomization was utilized to fabricate Al8SiCu and TiB₂/Al8SiCu pre-alloyed powders under argon protection. The elemental composition in

Table 1. The chemical composition of the Al8SiCu powders and TiB₂/Al8SiCu composite powders.

Alloy (wt.%)	Si	Cu	Mn	Mg	Ti	B	Al
Al8SiCu	7.48 ± 0.03	1.42 ± 0.01	0.47 ± 0.01	0.40 ± 0.01	0.09 ± 0.02	/	Bal.
TiB2/Al8SiCu	7.68 ± 0.03	1.40 ± 0.01	0.45 ± 0.01	0.39 ± 0.01	2.89 ± 0.02	1.24 ± 0.01	Bal.

was determined using inductively coupled plasma atomic emission spectrometry (ICP-AES), revealing that the proportion of TiB₂ particles was 3.98 wt.%. The morphology and particle size distribution are depicted in Figure 1. The pre-alloyed powder showed high sphericity with minimal satellite attachments, and the TiB₂ particles were dispersed within the TiB₂/Al8SiCu composite powders. The size distribution of the Al8SiCu and TiB₂/Al8SiCu powders ranged from 11 to 62 μm with the average size of 30 μm. The TiB₂ particles were quite uniformly distributed throughout the matrix with minimal agglomeration (Figure 1(e₁–e₆)), and the distribution of the Ti element served as an indicator of the presence of TiB₂ within the powder composition. The powders were dried at 75 °C for 8 h in a vacuum oven and the substrate was preheated to 70 °C before processing by PBF-LB. The samples were fabricated at room temperature using the EP-M150 system equipped with an IPG Fiber Laser. The Al8SiCu and TiB₂/Al8SiCu alloys were processed using identical processing parameters: a laser power (P) of 230, 260, 290, and 320 W; a scanning speed (v) of 600, 800, 1000, 1200, and 1400 mm/s; a hatch spacing (h) of 0.1 mm; a layer thickness (t) of 0.03 mm; and a rotation angle of 67° between layers. The dimensions of the processed cube and rectangular samples were 10 mm × 10 mm × 10 mm and 84 mm × 12 mm × 10 mm, respectively.

2.2. Microstructure Characterization

The phase composition of the PBF-LBed specimens was determined by X-ray diffraction technology (XRD) equipped with a Cu Kα radiation setting in the 2θ range from 20° to 80°. The crystallographic database used in this work refers to the International Centre for Diffraction Data (ICDD) database. Meanwhile, the Rietveld analysis method was applied to calculate the quantities of each of the phases identified in the spectra. The drainage method was employed to assess the densification behavior. The surface defects were observed under an optical microscope (OM) after the samples were polished. The microstructure of the samples etched with Keller’s reagent was observed with a scanning electron microscope (SEM) fitted with an energy-dispersive X-ray spectroscope (EDS) at an acceleration voltage of 20 kV. The grain size and orientation were investigated in terms of electron backscatter diffraction (EBSD) with a step size of 1.1 μm and an acceleration voltage of 20 kV, and the raw data were analyzed using TSL OIM 8.0 software. Disk-shaped samples prepared using the ion-thinning method were examined through transmission electron microscope (TEM) to further characterize and analyze the detailed microstructure. Additionally, the lattice strain (ε), dislocation density (δ), and lattice constant (D) of both the PBF-LBed Al8SiCu alloy and the PBF-LBed TiB₂/Al8SiCu composite listed in

Table 2. The lattice strain, dislocation density, and lattice parameter of the Al phase for the PBF-LBed Al8SiCu alloy and the TiB₂/Al8SiCu composite.

Sample	Lattice Strain	Dislocation Density	Lattice Parameter
Al8SiCu	2.87 × 10−3	8.53 × 1014 m−2	4.056 Å
TiB2/Al8SiCu	3.01 × 10−3	9.45 × 1014 m−2	4.066 Å

were estimated based on XRD results. The calculation utilized the following formulas [33]:

$$D=K\lambda /(\beta \times \cos \theta )$$

(1)

$$\delta =1/D^2,$$

(2)

$$\epsilon =\beta /4\times \tan \theta ,$$

(3)

where K is a constant (0.9), β is the maximum value of the half-maximum width, λ is a constant related to the wavelength (0.15405 nm), D is the lattice constant, ε is the lattice strain, and θ is the diffraction angle.

2.3. Mechanical and Frictional Properties

The hardness was obtained by a micro-Vickers hardness tester with a 100 g load for 10 s. The average value was derived from the measurements taken at 8 separate points on each specimen. Dog-bone-shaped tensile samples fabricated through electric discharge wire-cutting technology were subjected to testing on a universal tensile testing machine with a strain rate of 1 × 10⁻³ s⁻¹. The tensile properties were determined by averaging the results from at least five specimens. Room temperature-reciprocating friction wear tests were conducted on the PBF-LBed samples using a ball-on-disk tribometer with GCr15 balls. The GCr15 ball applied a 7 N load to the specimens and performed reciprocating motion on its surface for 30 min. The motion speed and stroke length of the GCr15 ball were 40 mm/s and 12 mm, respectively.

3. Results

3.1. Densification Behavior

The relative density is a critical indicator for evaluating the densification effectiveness of PBF-LB. The volume energy density (VED) significantly determines the relative density of the samples manufactured by PBF-LB. The VED is usually described as in [34]: VED = P/(vht). The relative densities of the PBF-LBed TiB₂/Al8SiCu composites and Al8SiCu alloys exhibited a similar trend of initial increase and then decrease when increasing the VED. The maximum relative densities of the Al8SiCu without and with TiB₂ addition were measured as 99.8% and 99.7%, respectively. The optimal VED range for achieving a relative density above 99.5% was found between 56.2 J/mm³ and 91.6 J/mm³ for the TiB₂/Al8SiCu composite and 68.7 J/mm³ and 94.8 J/mm³ for the Al8SiCu alloy. Clearly, the addition of in situ TiB₂ not only lowered the optimal VEDs for high-density Al8SiCu alloys, but it also expanded the range of the process parameters for PBF-LB process. Meanwhile, the OM results in Figure 2b indicate that low VEDs resulted in incomplete powder melting and poor density (Figure 2(b1,b4)). The powders were fully melted and exhibited good wettability as the VED increased, and only small pores were observed in the good samples (Figure 2(b2,b5)). However, further elevations of VED triggered intensive convection within the melt pool and caused element evaporation, leading to the formation of more holes. Also, the presence of TiB₂ enhanced the laser absorption rate of the pre-alloyed powders and made the melt pool more unstable at high VEDs, resulting in more pinholes in the TiB₂/Al8SiCu composite (Figure 2(b3,b6)).

3.2. Phase Analysis

The XRD spectra for the phase analysis in the PBF-LBed TiB₂/Al8SiCu composite and Al8SiCu alloy are shown in Figure 3. The α-Al phase and Si phase were detected in both samples, and the TiB₂ peak was only detected in the TiB₂/Al8SiCu composites. The quantities of the α-Al phase and Si phase in Al8SiCu were 92.5 wt.% and 7.5 wt.%, respectively. In TiB₂/Al8SiCu, the quantities of the α-Al phase, Si phase, and TiB₂ were 88.9 wt.%, 6.8 wt.%, and 4.3 wt.%, respectively. Also, Al₃Ti and Al₂B phases were not found in the composites, indicating that the TiB₂ was stable in the composite and that were was no decomposition or the formation of metastable phases. Also, no Mg- and/or Cu-enriched phases were observed in the XRD spectra, indicating the complete dissolution of these solute elements into the matrix (Figure 3a). Meanwhile, the introduction of TiB₂ induced a shift of the Al peaks toward lower diffraction angles, suggesting an expansion in the lattice constant of Al in accordance with Bragg’s law (Figure 3b). It was anticipated that the enlargement of the lattice constant was associated with the increased dissolution of Mg and Cu elements under a high cooling rate in PBF-LB [35]. Additionally, the strongest (200)_α-Al peak in the PBF-LBed Al8SiCu alloy implied the preferred crystallographic orientation during the PBF-LB process. But the intensity of (200)_α-Al peak was weakened in the TiB₂/Al8SiCu composite, suggesting that TiB₂ might inhibit the growth of Al grains in a preferred direction (Figure 3b).

According to the XRD data and Formulas (1)–(3), it was also observed that the lattice parameter, dislocation density, and lattice strain of the PBF-LBed TiB₂/Al8SiCu composites were all higher than those of the PBF-LBed Al8SiCu alloy (Table 2). The incorporation of TiB₂ led to the enlargement of the α-Al lattice parameter while also inducing an increase in the dislocation density. The increase in dislocation density enhanced the mechanical properties of the TiB₂/Al8SiCu composite.

3.3. Microstructural Characterization

Non-equilibrium solidification is the key feature in the PBF-LB process, resulting in an increase in the solubility of the solute elements and the formation of melting pools containing a cellular structure in the two experimental materials (Figure 4). Similar to the observations of the cellular structures for other alloys [36], the melt pool was composed of a fine zone (FZ), heat-affected zone (HAZ), and coarse zone (CZ) in the TiB_2/Al8SiCu composite and the Al8SiCu alloy. Also, TiB₂ particles were observed within the α-Al cells and at the cell boundaries, with sizes ranging from 100 nm to 3 μm in the TiB₂/Al8SiCu composite. During solidification under a high cooling rate, the α-Al phase was precipitated first, and then the Si and other solutes accumulated in front of the solid/liquid interface and segregated along the cell boundaries [37], as shown in Figure 4b,d. The fine precipitates were observed within the supersaturated α-Al cells. As can be seen in Figure 4, the addition of TiB₂ particles had little effect on the α-Al cellular structure but significantly refined the grains. The changes in the grain and detailed microstructural characteristics will be investigated in the following sections.

Figure 5 shows that EBSD results of the PBF-LBed Al8SiCu alloy and TiB₂/Al8SiCu composite in a building direction. It was observed that the TiB₂ promoted a columnar to equiaxed transition (CET) of the α-Al grains in the PBF-LBed Al8SiCu alloy. The texture was 2.081 in the microstructure of the Al8SiCu alloy with the feature of columnar crystals and fine equiaxed crystals distributed at the melt pool boundaries. During the PBF-LB process, the temperature gradient (G) and the solidification rate (R) affected the grain morphology and grain size. The ratio G/R was confirmed to control the grain morphology of the microstructure, and the G×R determined the grain size. At the bottom of the melt pools, a lower G/R ratio resulted in the formation of fine equiaxed crystals at the boundaries of the pool [38]. The steep thermal gradient within the melt pools led to the growth direction of the columnar crystals being the almost aligned with the vertical direction, and most of these were along the <001> direction, which was consistent with the XRD results (Figure 3a). Due to the layer-by-layer accumulation and melting, some of the coarse columnar crystals would penetrate into another melting pool. Clearly, no coarse columnar crystals were detected in the PBF-LBed TiB₂/Al8SiCu composite. The size of the columnar crystals was significantly decreased with a weak average texture intensity (1.76), suggesting the random grain orientation was strengthened. The obvious CET induced via TiB₂ addition promoted the grain refinement in the Al8SiCu alloy, of which the average grain size decreased from 15.31 μm to 7.34 μm (Figure 5c,d). Moreover, the addition of TiB₂ generated a high fraction of high-angle grain boundaries (HAGBs > 15°, Figure 5e,f). The HAGBS were known as the barriers to dislocation motion and thus could enhance the mechanical properties [39]. Further analysis of the grain aspect ratio is presented in Figure 5g,h. The ratio of length to width was used to define the aspect ratio. For clarity, grains with aspect ratios below 2 were classified as equiaxed grains, those between 2~4 as columnar grains, and those above 4 as coarse columnar grains. Clearly, the existence of TiB₂ elevated the proportion of the equiaxed crystals from 36.3% to 49.9% and decreased the fraction of the coarse columnar crystals from 21.2% to 11.6%.

Detailed microstructural characteristics of the PBF-LBed Al8SiCu alloy without and with TiB₂ addition are shown in Figure 6 and Figure 7. Bright field TEM (BF-TEM) images showed that the microstructure of the PBF-LBed Al8SiCu alloy featured cellular structures (Figure 6a), similar to the observations in Figure 4. The interaction between a high number density of precipitates and dislocations was also observed within the cellular structures (Figure 6b,c). Elemental results calibrated in Figure 6c demonstrated that the Si phases were the main precipitates, as verified by the Si segregation along the cellular structure boundaries. The β-Mg₂Si phase and the θ-Al₂Cu phase were identified by the enrichment of Mg, Si, and Cu in Figure 6 (c₂–c₄). The orientation relationships between the α-Al matrix and β-Mg₂Si phase/β″ phase was determined via high-resolution TEM (HR-TEM) and fast Fourier transform (FFT) images, as shown in Figure 6d–f. The orientation relationship between the Si phase and α-Al could be expressed as follows: (200)_α-Al//($\overline{1}11$)_Si, (220)_α-Al//(311)_Si, and [001]_α-Al//[$\overline{1}12$]_Si. The ($00\overline{1}$) plane of the β″ phase was found parallel to the (020) plane of α-Al, forming a semi-coherent interface. Also, the HR-TEM micrographs and corresponding FFT images shown in Figure 6f further confirmed the presence of the β-Mg₂Si phase at the cellular structures, of which the β phase possessed a coherent orientation with the α-Al matrix. The [$\overline{1}$12] axis of the β-Mg₂Si phase was parallel to the [001] axis of α-Al, and the summarized orientation relationships are shown below (($\overline{1}11$)_β//(020)_α−Al, (220)_β//(200)_α−Al).

Similar to the Al8SiCu alloy, the interaction between precipitates and dislocation was detected within the Si eutectics in the PBF-LBed TiB₂/Al8SiCu composite (Figure 7a–c). The EDS results shown in Figure 7(c₁–c₆) confirmed that the nanoscale precipitates were Si phases. The interaction between the tangled dislocations and β-Mg₂Si phase/θ-Al₂Cu phase was also detected at the vicinity of TiB₂ particles. The precipitates and dislocations interacted with each other, generating elastic stress fields that hindered the migration of dislocations and increased the strength of the composite. BF-TEM images of prismatic TiB₂ with a size of 100 nm and inserted SAED patterns are displayed in Figure 7d. The HAADF-STEM image in Figure 7e,f shows in situ particles within the cellular structure. The corresponding elemental distribution confirmed that the precipitates were in a Si phase and Al₂Cu phase. Also, the selected area electron diffraction (SAED) indicated that the orientation relationship between the Al and θ- Al₂Cu phases was identified as follows: [$0\overline{1}1$]_α-Al//[101]_θ and ($\overline{1}\overline{1}\overline{1}$)_α-Al//(020)_θ (Figure 7f).

The detailed characteristics of interface between the submicron TiB₂ particles and the Al matrix are displayed in Figure 8. The TiB₂ particle showed a good relationship with the Al matrix, with the enrichment of Al, Ti, and Si segregated at the transition layer (Figure 8b). The EDX line profiles in Figure 8c further confirmed the presence of the Al-Ti layer at the TiB₂/α-Al interface, which is consistent with previous research conclusions [40]. The presence of the Al-Ti layer has been confirmed to be able to promote the heterogeneous nucleation of TiB₂ particles [41]. However, the presence of Si at the interface causes the Al-Ti layer to lose its ability to promote the nucleation of TiB₂ particles. Although the poisoning effect of Si inhibits the nucleation of TiB₂ particles, most of the TiB₂ particles can still play a role during the heterogeneous nucleation of α-Al due to their high content and predominantly nanoscale size.

3.4. Mechanical Properties

Figure 9a illustrates the engineering stress–strain curves of the PBF-LBed Al8SiCu alloy without and with TiB₂ addition, and the corresponding tensile properties are illustrated in Figure 9b. The Al8SiCu alloy and the TiB₂/Al8SiCu composite initially underwent elastic deformation followed by yielding and strain hardening before fracturing. The enhanced strain hardening capability was observed in the TiB₂/Al8SiCu composite, which was primarily ascribed to the abundant dislocation loops distributed within cellular structures [42]. The yield strength (YS), ultimate tensile strength (UTS), and elongation (El) of the PBF-LBed Al8SiCu alloy were 296 ± 6 MPa, 517 ± 6 MPa, and 11.7 ± 1.0%, respectively. The addition of TiB₂ particles enhanced the strength with a slight reduction in elongation, of which the UTS, YS, and El were 328 ± 8 MPa, 541 ± 3 MPa, and 9.1 ± 0.7%, respectively.

The fractured morphology shown in Figure 10 indicated that a large number of dimple structures and a few cleavage steps were observed on the fractured surfaces of both the Al8SiCu alloy and the TiB₂/Al8SiCu composite, suggesting a mixed mode of ductile and brittle fracture patterns. Meanwhile, a large number of pores ranging from 5–15 μm were detected on the fractured surface of the TiB₂/Al8SiCu composite because of the interfacial debonding between the TiB₂ particles and the α-Al matrix during tensile process [43], while the micrometer-scale small pores at the fractured surface of the Al8SiCu alloy were highly associated with the non-escaping gas in time during the PBF-LB solidification process [44]. Clearly, a high number density of pores and partially agglomerated TiB₂ particles resulted in a low poor elongation for the TiB₂/Al8SiCu composite as a counterpart with the Al8SiCu alloy.

3.5. Tribological Behavior

The addition of ceramic particles with high hardness normally improves the wear resistance of metallic materials [45,46]. Figure 11 shows the friction coefficient, the wear-off rate, and the hardness of the PBF-LBed Al8SiCu alloy and the PBF-LBed TiB₂/Al8SiCu composite. Initially, the friction coefficient experienced notable fluctuations due to the insufficient contact between the milling ball and the samples. With the increase in friction time, the friction coefficient gradually stabilized and maintained no change for the experimental time. The calculated data revealed that the Al8SiCu alloy had an average friction coefficient of 0.655 and a wear rate of 1.76 × 10⁻³ mm³/Nm, while the TiB₂/Al8SiCu composite had a lower average friction coefficient of 0.580 and a reduced wear rate of 1.38 × 10⁻³ mm³/Nm. The TiB₂ addition reduced the average friction coefficient by 11.4% and the wear rate by 21.5%.

Figure 12 shows the morphology of the worn surfaces of the PBF-LBed Al8SiCu alloy and TiB₂/Al8SiCu composite. It was observed that the rough worn surface was featured by deep ploughing grooves in the PBF-LBed Al8SiCu alloy (Figure 12a). Some severe worn areas also showed fragmentation and peeling and the existence of grooves and pits (Figure 12b), indicating the poor wear resistance and adhesive wear mechanism of the PBF-LBed Al8SiCu alloy. In the counterpart of the TiB₂/Al8SiCu composite, only a low density of the thinner/shallower grooves and a high density of wear debris were detected at the worn surface. During the friction test, the hard TiB₂ particles that possessed a good interfacial relationship with the ɑ-Al matrix did not easily fall off during the testing. Therefore, the frictional force and the wear rate were correspondingly reduced in the TiB₂/Al8SiCu composite. Meanwhile, TiB₂ particles were disengaged from the matrix and generated wear debris on the wear track surface as the friction time was prolonged, prompting a shift to a combined wear mechanism of abrasive and adhesive wear (Figure 12d). Moreover, the incorporation of TiB₂ particles resulted in an evident grain refinement that enhanced the load-bearing capacity of the TiB₂/Al8SiCu composite, which effectively inhibited the initiation of cracks during friction.

4. Discussion

4.1. Effect of In-Situ TiB₂ Particles on Microstructural Evolution

Referring to the results mentioned above, the introduction of an addition of in situ TiB₂ offered the following positive influence: (1) The addition of TiB₂ increased the laser absorption rate of the powder and achieved a high relative density under lower VED conditions. (2) The refinement of the composite microstructure and the well-fitted interface between the TiB₂ particles and the Al matrix resulted in significant improvements in the mechanical and friction properties. (3) The TiB₂ particles provided effective nucleation sites for α-Al grains and promoted the CET. Consequently, the microstructure was refined in the PBF-LBed composite. Schematic diagrams of the microstructural characteristics of the PBF-LBed Al8SiCu alloy and TiB₂/Al8SiCu composite are shown in Figure 13. It was observed that the preliminary solidified areas underwent partial remelting during PBF-LB, which led to epitaxial grain growth in the Al8SiCu alloy, which formed columnar crystals (Figure 5a and Figure 13a). However, an obvious CET effect can be achieved after the precipitation of TiB₂, of which the TiB₂/Al8SiCu composite was characterized by equiaxed grains (Figure 5b and Figure 13b).

To achieve effective grain refinement, the ceramic particles must not only maintain a favorable interface with α-Al, but also uniformly distribute in the entire melted pools. The TiB₂ particles showed a good orientation relationship with the Al matrix (Figure 8b), and this also increased the laser absorption in improving the energy in the melt pool, which causes intensive Marangoni convection, which improves the distribution of TiB₂ particles in the melted pools [47]. It was also noted that the Ti atoms were initially adsorbed on the {0001} plane of TiB₂ during solidification, which led to the formation of the typical Al-Ti layer, as well as reductions in the lattice mismatch between the TiB₂ and the α-Al matrix (from 4.22% to 0.09%) [48,49]. Correspondingly, the uniformly distributed TiB₂ particles that possessed a low lattice mismatch with the ɑ-Al matrix were able to exert a positive influence in increasing the heterogeneous nucleation rate and promoting grain refinement [50]. Although previous studies [51] have indicated that a small fraction of TiB₂ particles are effective in heterogeneous nucleation, the majority of the TiB₂ particles effectively limited the epitaxial growth of grains, further leading to grain refinement and the formation of equiaxed crystals [52]. Last but not least, the in situ TiB₂ particles that possessed a large difference in thermal expansion coefficient with ɑ-Al matrix promoted an increase in dislocation density within the cellular structure, which builds a solid foundation for strength enhancement (Figure 9).

4.2. Effect of In-Situ TiB₂ Particles on Strengthening Mechanisms

As shown in Figure 9, the incorporation of TiB₂ into the Al8SiCu alloy enhanced the yield strength at a level of 328 MPa. The yield strength of the TiB₂/Al8SiCu composite mainly originates from the following strengthening mechanisms: grain boundary strengthening (${\sigma }_g$), dislocation strengthening (${\sigma }_{dis}$), load-bearing strengthening (${\sigma }_{load}$), Orowan strengthening (${\sigma }_{Orowan}$), and solid solution strengthening (${\sigma }_{ss}$). The formula for estimating the yield strength is as follows:

$$\sigma ={\sigma }_g+{\sigma }_{dis}+{\sigma }_{load}+{\sigma }_{Orowan}+{\sigma }_{ss}.$$

(4)

Compared to the average grain size of the Al8SiCu alloy (15.31 μm), the average grain size of the TiB₂/Al8SiCu composite was significantly reduced (7.34 μm), increasing the ability to hinder the dislocation movement. The contribution of grain boundary strengthening to the strength of the TiB₂/Al8SiCu composite can be calculated as follows:

$${\sigma }_g={\sigma }_0+K\times {d_c}^{-1/2},$$

(5)

where ${\sigma }_0$ is the lattice friction strength (20 MPa), K is the Hall–Petch coefficient (0.14 MPa m^1/2) [16], and d_c is the average grain size of the TiB₂/Al8SiCu composite (7.34 μm). Therefore, the contribution of the grain boundary strengthening to strength was approximately 71.7 MPa.

The data in Table 2 were obtained through a calculation of the XRD test results. The dislocation density of the Al8SiCu alloy was 8.53 × 10¹⁴ m⁻², while the dislocation density (${\rho }_c$) of the TiB₂/Al8SiCu composite was 9.45 × 10¹⁴ m⁻². Due to the difference in the coefficients of thermal expansion between Al (23.5 × 10⁻⁶ K⁻¹) and TiB₂ (7 × 10⁻⁶ K⁻¹), the dislocation density at the TiB₂/Al interface increased. The contribution of the dislocation strengthening to the strength of the TiB₂/Al8SiCu composite can be calculated as follows:

$${\sigma }_{dis}=M\beta G_{m}b{{\rho }_c}^{{\displaystyle \frac{1}{2}}},$$

(6)

where M is the Taylor factor (3.06 for the FCC crystal structure of Al), β is the dislocation strengthening coefficient (0.16), $G_m$ is the shear modulus (26.5 GPa), and $b$ is the Burgers vector (0.286 nm) [11]. Consequently, the strength attributed to dislocation strengthening was about 114.1 MPa.

The TiB₂ particles had a good interface relationship with the aluminum matrix, allowing the load to be transferred from the aluminum matrix to the TiB₂, which led to an increase in yield strength. The contribution of the load-bearing strengthening can be expressed as follows [29]:

$${\sigma }_{load}=1.5V_p{\sigma }_m,$$

(7)

where $V_p$ is the volume fraction of TiB₂ particles (2.3 vol%) and ${\sigma }_m$ is the bonding strength between the TiB₂ particles and the Al-Si matrix (~676 MPa [17]); thus, the contribution of the load-bearing strengthening mechanism to the strength was 23.3 MPa.

In the TiB₂/Al8SiCu composite, the reciprocal scanning of the laser during the PBF-LB process resulted in a significant precipitation of the nano-Si phases within the cellular structure. The nano-Si phases within the cells contributed to Orowan strengthening. Since most of the TiB₂ was mainly distributed at the grain boundaries or cell boundaries, the contribution of the TiB₂ particles to Orowan strengthening was minimal and can be neglected. The strength contribution from the Orowan strengthening mechanism can be calculated as follows [14]:

$${\sigma }_{Orowan}={\displaystyle \frac{\phi G_{m}b}{d_p}}{\left({\displaystyle \frac{6V_p}{\pi }}\right)}^{1/3},$$

(8)

where φ is a constant 2, $G_m$ is the shear modulus (26.5 GPa), b is the Burgers vector (0.286 nm), $V_p$ is the volume fraction of the Si particles within the coarse Si eutectic networks (3.2 vol%), and $d_p$ is the average size of the Si particles (80 nm). The values of $V_p$ and $d_p$ were obtained through the analysis of TEM images using ImageJ 1.53 software. The contribution of the Orowan strengthening to the strength was approximately 74.6 MPa.

During the PBF-LB process, the high cooling rate caused the solubility of the Si, Cu, Mg elements in the Al matrix to far exceed their equilibrium solubility limits, resulting in excellent solid solution strengthening. The contribution of the solid solution strengthening can be calculated using the following formula:

$${\sigma }_{\mathit{ss}}=k_{\mathit{Cu}}{C}_{\mathit{Cu}}^m+k_{\mathit{Si}}{C}_{\mathit{Si}}^m+k_{Mg}{C}_{\mathit{Mg}}^m,$$

(9)

where k_Cu, k_Si, k_Mg, and m are 13.8 MPa/wt.%, 11 MPa/wt.%, 18.6 MPa/wt.%, and 1, respectively [53]; and C_Cu, C_Mg, and C_Si are the concentration of Cu, Mg, and Si in the solution, respectively. According to the results of the EDS measurement, the values of C_Cu, C_Mg, and C_Si were 3.57 wt%, 1.23 wt%, and 0.26 wt%, respectively. The contribution of the solid solution strengthening to the strength was approximately 61.1 MPa.

$$\sigma ={\sigma }_g\left(71.7\right)+{\sigma }_{dis}\left(114.1\right)+{\sigma }_{load}\left(23.3\right)+{\sigma }_{Orowan}\left(74.6\right)+{\sigma }_{ss}\left(61.1\right).$$

The value of $\sigma$(344.8 MPa) calculated from the formula was slightly greater than the result measured by mechanical performance tests (328 MPa). The excessive increment may have resulted from the volume fraction of the nano-Si phase used in Formula (8) exceeding the actual volume fraction of the precipitated Si phase in the TiB₂/Al8SiCu composite, leading to an overestimation of the Orowan strengthening.

5. Conclusions

In this study, a systematic study was conducted on the parameter optimalization, microstructural evolution, mechanical properties, and friction characteristics of the Al8SiCu alloy without and with in situ TiB₂ particles. The main conclusions can be drawn as follows:

(1) The addition of in situ TiB₂ particles into Al8SiCu alloys reduces the volumetric energy density required for the high-density TiB₂/Al8SiCu composite. Meanwhile, TiB₂ particles promote the transformation from columnar to equiaxed crystals and the formation of high-angle grain boundaries. The grains on the vertical direction of the PBF-LBed TiB2/Al8SiCu composite are much smaller than those of the PBF-LBed Al8SiCu alloy.

(2) The yield strength (YS), ultimate tensile strength (UTS), and elongation (El) of the PBF-LBed Al8SiCu alloy were 296 ± 6 MPa, 517 ± 6 MPa, and 11.7 ± 1.0%, respectively. The PBF-LBed TiB₂/Al8SiCu composite achieved a balance between the strength and the ductility, with the yield strength of 328 ± 8 MPa, an ultimate tensile strength of 541 ± 3 MPa, and an elongation of 9.1 ± 0.7%. The increase in strength primarily resulted from grain boundary strengthening, dislocation strengthening, load-bearing strengthening, solid solution strengthening, and Orowan strengthening, of which the dislocation strengthening and Orowan strengthening were critical in the strengthening.

(3) The enhanced hardness resulting from the refinement of grains and the presence of in situ TiB₂ particles also led to lower friction coefficients and wear off rates, which, in turn, enhanced the wear resistance behavior.