Laser Powder-Bed Fusion of Ceramic Particulate Reinforced Aluminum Alloys: A Review

Aluminum (Al) and its alloys are the second most used materials spanning industrial applications in automotive, aircraft and aerospace industries. To comply with the industrial demand for high-performance aluminum alloys with superb mechanical properties, one promising approach is reinforcement with ceramic particulates. Laser powder-bed fusion (LPBF) of Al alloy powders provides vast freedom in design and allows fabrication of aluminum matrix composites with significant grain refinement and textureless microstructure. This review paper evaluates the trends in in situ and ex situ reinforcement of aluminum alloys by ceramic particulates, while analyzing their effect on the material properties and process parameters. The current research efforts are mainly directed toward additives for grain refinement to improve the mechanical performance of the printed parts. Reinforcing additives has been demonstrated as a promising perspective for the industrialization of Al-based composites produced via laser powder-bed fusion technique. In this review, attention is mainly paid to borides (TiB2, LaB6, CaB6), carbides (TiC, SiC), nitrides (TiN, Si3N4, BN, AlN), hybrid additives and their effect on the densification, grain refinement and mechanical behavior of the LPBF-produced composites.


Introduction
In many engineering solutions, product performance is determined by weight, which can be scaled down by material-efficient construction and the use of low-density alloys [1,2]. Due to exceptional strength/stiffness-to-weight ratio, low density, good damage tolerance, ability to be heat treated and the low cost, aluminum (Al) alloys are extensively used in many exclusive fields, such as: automotive, aerospace, marine navigation, rail transit, architectural construction, microelectronics and consumer applications [3][4][5][6][7].
In the meantime, owing to the moderate strength and relatively poor wear resistance of aluminum alloys, they are not applicable as structural materials for critical parts of aircrafts or satellites [8,9]; therefore, there is a need to improve the mechanical properties of aluminum alloys to be used for special applications. Along the modern industrial developments, the demand for complex-shaped products in diverse sectors is widespread. Problems related to traditional casting of aluminum alloys include coarse microstructures, a long process chain with limited flexibility [10], use of PM/casting molds [11] and a high rate of tool degradation [12].
Additive manufacturing (AM) provides an integrated way of item production [13]. Additive manufacturing, also known as 3D printing, refers to the layer-wise fabrication process of functional objects adopting nearly unlimited geometrical complexity, processing freedom, high level of accuracy and customization with elimination of traditional economyof-scale constraints [14]. Furthermore, the material efficiency and design flexibility of AM technology meet the requirements for resource optimization, mass customization and For a wide acceptance of the alloys for industrial use, the materials must ensure a number of required properties. The ideal alloy must be highly matched for the extreme thermal conditions by means of decreasing fabrication defects. Meanwhile, it is crucial for it to possess a suitable microstructure along with specific mechanical properties, which are comparable to the existing peak-aged wrought alloys, and to maintain a major part of its strength at elevated or high temperatures [30]. To further improve the mechanical performance of LPBF-prepared aluminum alloys, a substantial amount of research has been devoted to the following: (i) Studying the modification of existing compositions by minor alloying constituents to generate strengthening phases upon the fabrication process or during post-processing (heat treatment) [32]. (The effects of common modifying elements are given in Figure 1). (ii) The addition of grain refiners (stable, non-soluble solid ceramic particulates) to reduce hot-tear susceptibility, grain growth and dislocation motion by developing aluminum matrix composites (AMC) [8,33]. The latter conveys a combination of properties of two or more physically distinct phases with the aim to produce parts with far superior properties to the individual components [34]. (iii) Heat treatment [35][36][37].
Materials 2022, 15, x FOR PEER REVIEW 3 of 41 so determined "printable" alloys through LPBF possess a non-uniform microstructure and demonstrate poor mechanical performance [31]. For a wide acceptance of the alloys for industrial use, the materials must ensure a number of required properties. The ideal alloy must be highly matched for the extreme thermal conditions by means of decreasing fabrication defects. Meanwhile, it is crucial for it to possess a suitable microstructure along with specific mechanical properties, which are comparable to the existing peak-aged wrought alloys, and to maintain a major part of its strength at elevated or high temperatures [30]. To further improve the mechanical performance of LPBF-prepared aluminum alloys, a substantial amount of research has been devoted to the following: (i) Studying the modification of existing compositions by minor alloying constituents to generate strengthening phases upon the fabrication process or during postprocessing (heat treatment) [32]. (The effects of common modifying elements are given in Figure 1) (ii) The addition of grain refiners (stable, non-soluble solid ceramic particulates) to reduce hot-tear susceptibility, grain growth and dislocation motion by developing aluminum matrix composites (AMC) [8,33]. The latter conveys a combination of properties of two or more physically distinct phases with the aim to produce parts with far superior properties to the individual components [34]. (iii) Heat treatment [35][36][37].  [14,27,30,33,[38][39][40][41][42][43][44][45][46][47][48][49][50][51][52][53][54].
AM processes are categorized as master forming technologies, where customized designed objects' properties are generated by the fabrication process itself. Therefore, the composition and aluminum alloy chemistry undertake a central role during the LPBF process [1]. Combining the advantages offered by AM with the favorable mechanical properties of aluminum alloys will create viable mass-market manufacturing strategies that will increase the adoption and implementation of both across the world [7].
AM processes are categorized as master forming technologies, where customized designed objects' properties are generated by the fabrication process itself. Therefore, the composition and aluminum alloy chemistry undertake a central role during the LPBF process [1]. Combining the advantages offered by AM with the favorable mechanical properties of aluminum alloys will create viable mass-market manufacturing strategies that will increase the adoption and implementation of both across the world [7].
In this review paper, the focus is placed on the laser powder-bed fusion of the ceramic particulate (boride, carbide, nitride and hybrid additive) reinforced aluminum alloys, concentrating on the effect of additives on the microstructure and grain refinement of the produced materials. Thereafter, the mechanical properties and the mechanisms responsible for their change are confronted to lead to a deeper understanding of the possible performance of ceramic particulate reinforced aluminum matrix composites (AMCs). The list of used reinforcements and their unique features during the LPBF process, as well as diagrams relationships (ORs) between two crystals, as different lattice parameters cause distortion of the lattice, resulting in an excess strain energy, which is determined by a lattice mismatch (also called as lattice disregistry, δ) [58]. The selection of potent grain refiners with the smallest disregistry with the matrix crystal throughout a specific interface is favored [58]. If disregistry value is below 10%, both in situ formed and added inoculants have the ability to induce heterogeneous nucleation of Al grains [66].
Nucleant particles serve a dual role in the AMCs as refiners and reinforcements, and they can be classified in three categories: non-oxide ceramics, oxide ceramics and carbonbased compounds. Generally, the ceramic particulates of a high hardness, good thermal stability, relatively high laser absorptivity and compatibility with metals/alloys are suitable constituents for the preparation of high-performance AMCs [67]. To meet the demand to satisfy the "light weight and high strength" concept, novel AMCs are continuously under development [5,11,68].
For the conventional AMCs, relatively coarse ceramic particles with a size ranging from several tens to hundreds of micrometers are broadly utilized as reinforcements. However, reasoned by limited interfacial wettability between reinforcement and matrix, the large particles are susceptible to cracking during mechanical loading, causing reduced ductility and inducing premature failure of AMCs [69]. Consequently, both tensile strength and ductility of AMCs increase if the fine-sized reinforcements are used. On that account, the introduction of the nano-scaled ceramic particles can remarkably enhance the mechanical performance of AMCs [70,71].
However, the agglomeration of nanoparticles may cause unfavorable microstructural changes and affect the mechanical behavior of the composites, as well as affecting thermal and rheological behavior of the melt pool (increasing viscosity, especially in case of high volume of nanoparticles) and shifting the LPBF parameter window. The LPBF method enables effective fabrication of composites reinforced with ceramic reinforcements, taking into account the unique metallurgical nature of the process, high temperatures and thermal convection in a micron-sized molten pool [23,72,73].

Non-Oxide Additives
Non-oxide additives (borides, carbides, nitrides, etc.) are one of the most used reinforcements for Al alloys due to their high melting temperatures and chemical stability [74]. AMCs merge the ductility and toughness of aluminum with the high strength and modulus of the ceramic reinforcement [75], hence achieving an improvement of the overall characteristics and durability [12]. The low laser absorptivity of aluminum in the infrared range challenges the controlled melting, while the increase in the laser absorption of ceramic particulate decorated/mixed aluminum alloy at a laser wavelength of 1064 nm promotes the LPBF process. The introduction of ceramic particles to the pure alloy increases laser absorptivity of the overall powder mixture, as (i) non-oxide ceramic particles display high laser absorptivity and (ii) the added ceramic particles increase surface roughness of decorated powder, promoting multiple reflections of the laser in the powder bed [28]. As shown in Figure 2a-c, the ray absorption of the SiC/AlSi10Mg and TiB 2 /AlSi10Mg powder mixtures is higher compared to pure AlSi10Mg alloy. There is a lower intensity of interactions between laser rays and particles of pure AlSi10Mg compared to SiC and TiC added composite powder. (Figure 2d-g) [76].  Illustration of track spot of each laser ray on the particle surface of AlSi10Mg (d), SiC/AlSi10Mg (e), TiB2/AlSi10Mg (f) (side view) and numerical representation of laser-particle interactions (g) (reproduced with permission from [76]).

Borides: Grain Refining and Strengthening Effect of TiB2, LaB6, CaB6
As one of the proven highly effective grain refiners for Al alloy, TiB2 particles exhibits good thermal stability, good wettability and interfacial compatibility, in addition to the acknowledged crystallographic orientation relationship with Al matrix, contributing to a comprehensive mechanical performance of AMCs [59,73]. The addition of TiB2 to AlSi10Mg increases the laser absorptivity of the powder bed by almost 1.5 times [76]. To provide even distribution, small particle size and adequate interfacial bonding of the TiB2 particles, in situ fabrication approaches have been implemented, offering the advantages of a clean interface between ceramic particles and matrix alloy and fine morphology of in situ formed particles [5]. Both in situ and ex situ fabrication of TiB2 reinforced Al alloys are discussed below.
In Ref. [77], 0.5-8 wt.% nano-sized TiB2 particles were introduced into AlSi10Mg, which resulted in the elimination of columnar grains and refined elongated dendritic structures from 4.6 to 2 µm, as shown in Figure 3a-d and Table 1. Similar results were obtained in Refs. [59,73], as the introduction of 1-5 wt.% and 5.3 wt.% (3.4 vol%) TiB2 to AlSi10Mg, respectively, led to remarkable grain refinement down to 1.55 µm (Figure 3eg,i,j). However, the incorporation of only 1 wt.% TiB2 into AlSi10Mg [78] did not demonstrate a dramatic difference between reinforced and pure alloy parts; however, the grain size distribution became distinctly narrow (Figure 3h). As one of the proven highly effective grain refiners for Al alloy, TiB 2 particles exhibits good thermal stability, good wettability and interfacial compatibility, in addition to the acknowledged crystallographic orientation relationship with Al matrix, contributing to a comprehensive mechanical performance of AMCs [59,73]. The addition of TiB 2 to AlSi10Mg increases the laser absorptivity of the powder bed by almost 1.5 times [76]. To provide even distribution, small particle size and adequate interfacial bonding of the TiB 2 particles, in situ fabrication approaches have been implemented, offering the advantages of a clean interface between ceramic particles and matrix alloy and fine morphology of in situ formed particles [5]. Both in situ and ex situ fabrication of TiB 2 reinforced Al alloys are discussed below.
In Ref. [77], 0.5-8 wt.% nano-sized TiB 2 particles were introduced into AlSi10Mg, which resulted in the elimination of columnar grains and refined elongated dendritic structures from 4.6 to 2 µm, as shown in Figure 3a-d and Table 1. Similar results were obtained in Refs. [59,73], as the introduction of 1-5 wt.% and 5.3 wt.% (3.4 vol%) TiB 2 to AlSi10Mg, respectively, led to remarkable grain refinement down to 1.55 µm (Figure 3e-g,i,j). However, the incorporation of only 1 wt.% TiB 2 into AlSi10Mg [78] did not demonstrate a dramatic difference between reinforced and pure alloy parts; however, the grain size distribution became distinctly narrow (Figure 3h).
Partial melting of TiB 2 was reported in Ref. [73] despite the fact that TiB 2 is considered a refractory material. Adding 5 vol.% (or 8.3 wt.%) TiB 2 to an Al-Cu alloy [81] resulted in a remarkable grain size reduction from 23 to 2.5 µm. In Ref. [82], the in situ TiB 2 (4 wt.%) reinforced Al-Cu-Ag-Mg-Ti alloy had fine equiaxed grains with~0.64 µm average size without preferential orientation (Figure 3p). The reported grain size was smaller than that stated in Refs. [73,80]. In Refs. [64,83], the addition of 2 wt.% TiB 2 to an Al-12Si alloy produced a textureless microstructure with an average grain size of 5 µm, meaning that in case of similar content of incorporated TiB 2 , coarser grains were grown in the Al-12Si alloy than in AlSi10Mg (Figure 3m,n). For comparison, a hot-pressed sample's EBSD image is shown in Figure 3o, which, interestingly, showed a higher degree of grain refinement.
For a bare minimum boride additive range, at least 2 wt.% TiB 2 is sufficient to significantly alter the final morphology and crystallographic texture of LPBF-processed materials [64,73,77,82,83].   The grain refining (columnar to equiaxed transition) effect of TiB 2 (Figure 4a,b) is ascribed to its good stability in a melt pool, supplying numerous low-energy barrier nucleation sites (crystal embryos) and a reduction in the critical amount of total undercooling required to initiate the formation of equiaxed crystals [77]. The particles pushed to the grain boundaries pin and stabilize grain boundaries and limit grain growth along the heat flux direction [59]. Furthermore, due to a lower thermal conductivity of TiB 2 (~77.8 W/mK) as compared to Al (~108 W/mK) [73], TiB 2 particles prevent heat flux at a high temperature, reducing the temperature gradient. The latter results in the formation of fine equiaxed grains, weakening the texture and anisotropy of fabricated AMCs [59]. Overall, grain refinement is justified with a combination of high cooling rates during LPBF, an increased number of nucleation sites and limitations on grain growth [73,80], which lie beneath three main mechanisms: constitutional supercooling, heterogeneous nucleation and Zener pinning. Meanwhile, random orientations of TiB 2 particles provide the randomization of Al grain orientation and texture elimination [77]. produced a textureless microstructure with an average grain size of ~5 µm, meaning that in case of similar content of incorporated TiB2, coarser grains were grown in the Al-12Si alloy than in AlSi10Mg (Figure 3m,n). For comparison, a hot-pressed sample's EBSD image is shown in Figure 3o, which, interestingly, showed a higher degree of grain refinement.
The grain refining (columnar to equiaxed transition) effect of TiB2 (Figure 4a,b) is ascribed to its good stability in a melt pool, supplying numerous low-energy barrier nucleation sites (crystal embryos) and a reduction in the critical amount of total undercooling required to initiate the formation of equiaxed crystals [77]. The particles pushed to the grain boundaries pin and stabilize grain boundaries and limit grain growth along the heat flux direction [59]. Furthermore, due to a lower thermal conductivity of TiB2 (~77.8 W/mK) as compared to Al (~108 W/mK) [73], TiB2 particles prevent heat flux at a high temperature, reducing the temperature gradient. The latter results in the formation of fine equiaxed grains, weakening the texture and anisotropy of fabricated AMCs [59]. Overall, grain refinement is justified with a combination of high cooling rates during LPBF, an increased number of nucleation sites and limitations on grain growth [73,80], which lie beneath three main mechanisms: constitutional supercooling, heterogeneous nucleation and Zener pinning. Meanwhile, random orientations of TiB2 particles provide the randomization of Al grain orientation and texture elimination [77]. The grain refining effect of TiB2 is also reported to be a result of the formation of Al3Ti and the crystallographically coherent interface between Al3Ti and TiB2, which promotes the nucleation of Al3Ti on the surface of TiB2 particles in an Al melt. Without the Al3Ti layer, TiB2 additives are easily contaminated by impurities with a high tendency to form a eutectic microstructure with Al and, therefore, being insufficient in nucleating α-Al grains [85]. However, in Ref. [81], a preferable natural stacking sequence of Al atoms on The grain refining effect of TiB 2 is also reported to be a result of the formation of Al 3 Ti and the crystallographically coherent interface between Al 3 Ti and TiB 2 , which promotes the nucleation of Al 3 Ti on the surface of TiB 2 particles in an Al melt. Without the Al 3 Ti layer, TiB 2 additives are easily contaminated by impurities with a high tendency to form a eutectic microstructure with Al and, therefore, being insufficient in nucleating α-Al grains [85]. However, in Ref. [81], a preferable natural stacking sequence of Al atoms on TiB 2 and direct refining are reported. Meanwhile, in Ref. [82], it was highlighted that the absence of the Al 3 Ti layer does not prove a lack of nucleation, since the Al 3 Ti layer can fully transform into α-Al during the cooling process via a peritectic reaction.
Besides TiB 2 , other borides, such as CaB 6 and LaB 6, had shown a promising refining capability. The addition of 0.05-2 wt.% LaB 6 to AlSi10Mg resulted in grain refinement down to 1.6 µm (Figure 3q-t). LaB 6 particles form a highly coherent interface with the Al matrix. A higher amount of LaB 6 nanoparticles (>0.5 wt.%) did not further provide grain refinement and restricted longitudinal elongation due to the weakening of melt pool boundaries by segregation of the excess LaB 6 nanoparticles [84]. The addition of 2 wt.% CaB 6 nanoparticles to the high-strength 2024 aluminum alloy resulted in an equiaxed, crackfree microstructure with an average grain size of 0.91 ± 0.32 µm and a highly coherent interface with Al ( Figure 3u,v and Figure 5a,b) [28]. No decomposition of CaB 6 was observed. However, not every CaB 6 nanoparticle functions as a nucleant; a large quantity of them is acquired in the liquid phase between the growing grains, and they are forced to the grain boundaries where they stabilize the microstructure via Zener pinning. boundaries by segregation of the excess LaB6 nanoparticles [84]. The addition of 2 wt.% CaB6 nanoparticles to the high-strength 2024 aluminum alloy resulted in an equiaxed, crack-free microstructure with an average grain size of 0.91 ± 0.32 µm and a highly coherent interface with Al (Figures 3u,v and 5a,b) [28]. No decomposition of CaB6 was observed. However, not every CaB6 nanoparticle functions as a nucleant; a large quantity of them is acquired in the liquid phase between the growing grains, and they are forced to the grain boundaries where they stabilize the microstructure via Zener pinning. (HAADF-STEM stands for high-angle annular dark-field scanning transmission electron microscope, ADF for annular dark-field)(reproduced with permission from [28]).
In Ref. [77], the addition of 0.5-8 wt.% TiB2 to AlSi10Mg resulted in increased strength (up 544 MPa) and hardness (with 20%); however, the high content of TiB2 (>2%) resulted in a reduced ductility (6.2%), which was still higher than for a reference AlSi10Mg. Simultaneous enhancement of strength (up to 537 MPa and 530 MPa) and ductility (16.5% and 15.5%) was achieved in Refs. [79,80], respectively, when 6.5 wt.% and 11.6 wt.% TiB2 were introduced to AlSi10Mg. The increased strength was mainly attributed to the Hall-Petch relationship, loading-bearing and Orowan strengthening mechanisms. The grain boundary modification by TiB2 nano-particulates and the promoted dislocation plasticity by nano-Si precipitates improved ductility. LaB6 addition resulted in a subtle improvement of strength and ductility; however, the reinforcing effect was not as pronounced, as in the case of TiB2.
The highest elongation (~17.7%) was recorded in Ref. [82], when the Al-Cu alloy was reinforced with 4 wt.% TiB2; however, the alloys exhibited a significantly lower strength and hardness. The addition of 2 wt.% CaB6 [28] resulted in an increased elongation of 2024 alloy, up to 12.6%, and improved tensile and yield strength (Table 1).

Titanium Carbide: TiC
TiC exhibits several favorable characteristics required for Al alloys reinforcement; among them, there are moderate density (4.91 g/cm 3 ), high hardness (28-32 GPa) [86], high modulus of elasticity (up to 440 GPa) [87], good wettability, good laser absorptivity (higher than TiB2) and low lattice mismatch (6.9%) with Al. TiC particle reinforced AMCs In Ref. [77], the addition of 0.5-8 wt.% TiB 2 to AlSi10Mg resulted in increased strength (up 544 MPa) and hardness (with 20%); however, the high content of TiB 2 (>2%) resulted in a reduced ductility (6.2%), which was still higher than for a reference AlSi10Mg. Simultaneous enhancement of strength (up to 537 MPa and 530 MPa) and ductility (16.5% and 15.5%) was achieved in Refs. [79,80], respectively, when 6.5 wt.% and 11.6 wt.% TiB 2 were introduced to AlSi10Mg. The increased strength was mainly attributed to the Hall-Petch relationship, loading-bearing and Orowan strengthening mechanisms. The grain boundary modification by TiB 2 nano-particulates and the promoted dislocation plasticity by nano-Si precipitates improved ductility. LaB 6 addition resulted in a subtle improvement of strength and ductility; however, the reinforcing effect was not as pronounced, as in the case of TiB 2 .
The highest elongation (~17.7%) was recorded in Ref. [82], when the Al-Cu alloy was reinforced with 4 wt.% TiB 2 ; however, the alloys exhibited a significantly lower strength and hardness. The addition of 2 wt.% CaB 6 [28] resulted in an increased elongation of 2024 alloy, up to 12.6%, and improved tensile and yield strength (Table 1) TiC exhibits several favorable characteristics required for Al alloys reinforcement; among them, there are moderate density (4.91 g/cm 3 ), high hardness (28-32 GPa) [86], high modulus of elasticity (up to 440 GPa) [87], good wettability, good laser absorptivity (higher than TiB 2 ) and low lattice mismatch (6.9%) with Al. TiC particle reinforced AMCs have a high strength, stiffness and modulus, good corrosion and wear performance [22,72]. However, when formed in situ in the melt pool, the TiC phase possesses unstable chemical composition (portrayed as TiC x , where x is in 0.48-1 range) due to the generation of carbon atom vacancies. Consequently, the nucleating behavior of TiC x for α-Al is not consistent, since the TiC x +Al→Al 4 C 3 reaction is favored, which results in weakened grain refining performance [88].
In Ref. [89], an increase in the TiC content from 1 to 10 wt.% when added to the Al-15Si alloy resulted in an increase in melt pool fluidity and a decrease in the undercooling degree, leading to significant grain coarsening ( Figure 6). Ultimately, with the added threshold limit of TiC (10 wt.%), the primary Si particles precipitate out and distribute on the surface of the Al matrix ( Figure 6d). consistent, since the TiCx+Al→Al4C3 reaction is favored, which results in weakened grain refining performance [88].
In Ref. [89], an increase in the TiC content from 1 to 10 wt.% when added to the Al-15Si alloy resulted in an increase in melt pool fluidity and a decrease in the undercooling degree, leading to significant grain coarsening ( Figure 6). Ultimately, with the added threshold limit of TiC (10 wt.%), the primary Si particles precipitate out and distribute on the surface of the Al matrix ( Figure 6d). Alternatively, the fabrication of AlSi10Mg/5 wt.%-nano-TiC [70] under an increased laser energy caused the nano-TiC particles to accumulate in clusters, forming the micronsized agglomerates. However, the dispersion of reinforcement became more uniform, as shown in Figure 7a-d. Alternatively, the fabrication of AlSi10Mg/5 wt.%-nano-TiC [70] under an increased laser energy caused the nano-TiC particles to accumulate in clusters, forming the micronsized agglomerates. However, the dispersion of reinforcement became more uniform, as shown in Figure 7a [70]).
An increase in energy input resulted in change in TiC appearance, from aggr ring (circular) structures, due to intensive Marangoni flow (Figure 8a-d) in AlSi10Mg/3 wt.%TiC composites [71]. The formation of ring-structured TiC was reported in Ref. [22] as well. At 5 wt.% TiC addition, at elevated Marangoni force and a lower viscous drag fo ceramic particulates are captured in the circular melt motion (Figure 9b,c) and g distinct circular structures in solidified build (Figure 9e-g). The circular-structu agglomerate formation was not found in Ref. [70] when 5 wt.% TiC was used, wh probably be justified by the application of different process parameters. The formation of ring-structured TiC was reported in Ref. [22] as well. At 5 and 7.5 wt.% TiC addition, at elevated Marangoni force and a lower viscous drag force, the ceramic particulates are captured in the circular melt motion (Figure 9b,c) and generate distinct circular structures in solidified build (Figure 9e-g). The circular-structured TiC agglomerate formation was not found in Ref. [70] when 5 wt.% TiC was used, which can probably be justified by the application of different process parameters. The presence of in situ formed D022-Al3Ti inoculants (with tetragonal structure) was revealed in Ref. [31] for the AlSi10Mg/5 wt.%TiC composite. Heterogeneous nucleation of α-Al on the D022-Al3Ti nanoparticles (Figure 10c-f) occurred, leading to (i) columnar-toequiaxed transition with subsequent grain refinement from ~80 µm to ~1 µm ( Figure  11a,b), and (ii) the preferred orientation of the of α-Al (200) phase was removed ( Figure  10a,b). In situ formed Al3Ti served as a more effective nucleant as compared to TiC, mainly due to the small lattice mismatch between Al and Al3Ti, which was reduced to 0.09%. . Velocity vector plots around a TiC reinforcing particle in the melt pool (the dashed circles highlight the circular motion) and micrographs demonstrating typical morphology of LPBF-processed AlSi10Mg/TiC nanocomposites with different TiC contents: 2.5 wt.% (a,d), 5 wt.% (b,e) and 7.5 wt.% (c,f). Schematics of the formation mechanism of novel circular TiC configurations during fusion process at fixed E v = 571.43 J/mm 3 (g) (reproduced with permission from [22]).
The presence of in situ formed D0 22 -Al 3 Ti inoculants (with tetragonal structure) was revealed in Ref. [31] for the AlSi10Mg/5 wt.%TiC composite. Heterogeneous nucleation of α-Al on the D0 22 -Al 3 Ti nanoparticles (Figure 10c-f) occurred, leading to (i) columnar-toequiaxed transition with subsequent grain refinement from~80 µm to~1 µm (Figure 11a,b), and (ii) the preferred orientation of the of α-Al (200) phase was removed (Figure 10a,b). In situ formed Al 3 Ti served as a more effective nucleant as compared to TiC, mainly due to the small lattice mismatch between Al and Al 3 Ti, which was reduced to 0.09%. processed AlSi10Mg/TiC nanocomposites with different TiC contents: 2.5 wt.% (a,d), 5 wt.% (b,e) and 7.5 wt. % (c,f). Schematics of the formation mechanism of novel circular TiC configurations during fusion process at fixed Ev = 571.43 J/mm 3 (g) (reproduced with permission from [22]).
The presence of in situ formed D022-Al3Ti inoculants (with tetragonal structure) was revealed in Ref. [31] for the AlSi10Mg/5 wt.%TiC composite. Heterogeneous nucleation of α-Al on the D022-Al3Ti nanoparticles (Figure 10c-f) occurred, leading to (i) columnar-toequiaxed transition with subsequent grain refinement from ~80 µm to ~1 µm ( Figure  11a,b), and (ii) the preferred orientation of the of α-Al (200) phase was removed ( Figure  10a,b). In situ formed Al3Ti served as a more effective nucleant as compared to TiC, mainly due to the small lattice mismatch between Al and Al3Ti, which was reduced to 0.09%. Figure 10. Diffractograms of the LPBFed AlSi10Mg (a) and AlSi10Mg/5 wt.%TiC (b) specimens, HRTEM image of the D022-Al3Ti/Al matrix (c) and interface (d), SAED patterns taken at the D022-Al3Ti along (010) Al3Ti (e), FFT patterns of the D022-Al3Ti/Al matrix interface (f) (SAED stands for selected area electron diffraction and FFT for fast Fourier transform, (reproduced with permission from [31]).
Yet, another variable parameter centers on powder production for the LPBF process. In Ref. [90], the LPBF of the ball-milled composite powder of AlSi10Mg/5 wt.%TiC is reported. After printing, the TiC particles maintained their nanoscale nature and were not subjected to a significant coarsening, which resulted in an increased hardness of the alloy from 140 to 185 HV0.1 and the tensile strength from 400 to 482 MPa ( Table 2). The elongation of the composite part (10.8%) was similar to the elongation measured for the Yet, another variable parameter centers on powder production for the LPBF process. In Ref. [90], the LPBF of the ball-milled composite powder of AlSi10Mg/5 wt.%TiC is reported. After printing, the TiC particles maintained their nanoscale nature and were not subjected to a significant coarsening, which resulted in an increased hardness of the alloy from 140 to 185 HV 0.1 and the tensile strength from 400 to 482 MPa ( Table 2). The elongation of the composite part (10.8%) was similar to the elongation measured for the pure AlSi10Mg alloy. This can be explained by various effects: (i) an increased dislocation density near reinforcement/matrix interface, (ii) TiC nanoparticles acting as a barrier for dislocation movement, (iii) delaying crack propagation, thus improving the tensile strength. Alternating the TiC concentration, laser energy density and powder processing technique yield different composite attributes, as shown in Table 2.     pure AlSi10Mg alloy. This can be explained by various effects: (i) an increased dislocation density near reinforcement/matrix interface, (ii) TiC nanoparticles acting as a barrier for dislocation movement, (iii) delaying crack propagation, thus improving the tensile strength. Alternating the TiC concentration, laser energy density and powder processing technique yield different composite attributes, as shown in Table 2. Figure 11. EBSD color maps for LPBF-prepared Al alloys and AMCs reinforced with carbides, carbonitride, carbide/hydride and carbide/boride additives (a)-(n) (reproduced with permission from [31,88,[91][92][93]). While using a single carbide reinforcement has proven to be an effective way for grain refinement, the use of a second additive was shown to complement the effects of a single species. In Ref. [92], the dual reinforcing phases were used, resulting in a crack-free sample produced from the 2024 alloy/1 wt.%TiC+1 wt.%TiH 2 powders mixture. It was shown that unreinforced alloy contained columnar microstructure (Figures 11g and 12a-c), while the 2024 alloy/1 wt.%TiC+1 wt.%TiH 2 composite was composed of superfine equiaxed grains (Figures 11h and 12d-h).
While using a single carbide reinforcement has proven to be an effective way for grain refinement, the use of a second additive was shown to complement the effects of a single species. In Ref. [92], the dual reinforcing phases were used, resulting in a crack-free sample produced from the 2024 alloy/1 wt.%TiC+1 wt.%TiH2 powders mixture. It was shown that unreinforced alloy contained columnar microstructure (Figures 11g and 12ac), while the 2024 alloy/1 wt.%TiC+1 wt.%TiH2 composite was composed of superfine equiaxed grains (Figures 11h and 12d-h). Ti-rich particles (TiC and Al3Ti) with irregular or cubic shape are present in the grains exhibited in Figure 13a,b. The L12-Al3Ti with a face-center-cubic (FCC) structure is a result of TiH2 decomposition (TiH2→Ti+H2) and reaction between Ti and Al melt. It is worth mentioning that in Ref. [31], a formation of the D022-Al3Ti phase with a tetragonal structure was reported. A highly coherent interface between L12-Al3Ti and α-Al was observed (with 0.24% lattice mismatch) (Figure 13c), indicating that L12-Al3Ti might serve as substrate for heterogeneous α-Al nucleation; however, a coherent interface was not generated between TiC and Al (Figure 13d). Following the "Ti transition zone" theory (demonstrated in Figure 12), Ti-covered TiC nanoparticles, and then TiC particles themselves, become the effective nucleation substrates for α-Al as well. Ti-rich particles (TiC and Al 3 Ti) with irregular or cubic shape are present in the grains exhibited in Figure 13a,b. The L1 2 -Al 3 Ti with a face-center-cubic (FCC) structure is a result of TiH 2 decomposition (TiH 2 →Ti+H 2 ) and reaction between Ti and Al melt. It is worth mentioning that in Ref. [31], a formation of the D0 22 -Al 3 Ti phase with a tetragonal structure was reported. A highly coherent interface between L1 2 -Al 3 Ti and α-Al was observed (with 0.24% lattice mismatch) (Figure 13c), indicating that L1 2 -Al 3 Ti might serve as substrate for heterogeneous α-Al nucleation; however, a coherent interface was not generated between TiC and Al (Figure 13d). Following the "Ti transition zone" theory (demonstrated in Figure 12), Ti-covered TiC nanoparticles, and then TiC particles themselves, become the effective nucleation substrates for α-Al as well. On account of the inhibition of columnar grains, elimination of cracks, a refined microstructure and Orowan strengthening, the 2024 alloy/TiC+TiH2 AMC showed a simultaneous enhancement of tensile strength and ductility.
Another study on the fabrication of double TiB2-TiC reinforced AMCs [93] revealed that the addition of dual ceramic phases improved laser absorptivity by almost two-fold, On account of the inhibition of columnar grains, elimination of cracks, a refined microstructure and Orowan strengthening, the 2024 alloy/TiC+TiH 2 AMC showed a simultaneous enhancement of tensile strength and ductility.
Another study on the fabrication of double TiB 2 -TiC reinforced AMCs [93] revealed that the addition of dual ceramic phases improved laser absorptivity by almost twofold, substantially refining the Al grains (Figure 11i,k) and resulting in the increment in tensile strength (552 MPa) and elongation (12%) ( Table 2). It was revealed that the dual reinforcement more remarkably affected the mechanical performance, improved densification and grain refinement compared to the single reinforcement with the same total content (Table 2 and Figure 11j,l).
Double or triple reinforcements formed during in situ chemical reactions generate a composite material highly coherent with the metal matrix. When 0-17.2 wt.% (Ti-B 4 C) mixture was added to AlSi10Mg [94], the full densification of samples and in situ formation of ceramic phases were reported due to the combined LPBF and combustion synthesis (CS) process. Silicon atoms released from the alloy combine with Ti and C atoms, yielding the formation of transitional ternary carbide Ti 3 SiC 2 , while the remaining B 4 C and Ti are responsible for the formation of TiB 2 and TiC particulates ( Figure 14). The generation of the Ti 3 SiC 2 phase resulted in a significant drop in porosity of the fabricated sample. The heat released during the combustion reaction allowed for carrying out the fabrication in low laser energy regime.

Silicon Carbide: SiC
The SiC particle reinforced AMCs are applied in aerospace and electronic encapsulation, both in military and civilian fields, due to their high specific strength and stiffness, in addition to abrasion resistance. SiC has a much higher laser absorptivity (≈78%) than aluminum (≈7%), moderate density (3.21 g/cm 3 ), and it increases the laser absorptivity of the blended mixture [13,34,97,98]. During laser irradiation, SiC particles tend to heat up to extremely high temperature, leading to rapid reaction rates. Hence, the decrease in thermal conductivity results in further rise in temperature, the lifetime and fluidity of the melt pool. Meanwhile, an increase in SiC content in the initial feedstock and, hence, in the blend melt pool, increases the viscosity of a liquid melt and results in a lower fluidity. Therefore, both thermo-kinetic factors should be considered before selecting the content and size of the reinforcing SiC [11,13].
The chemical reaction between silicon carbide and aluminum melt at temperatures exceeding 940 K may result in SiC decomposition according to 4Al(l)+3SiC(s)→Al4C3(s)+3Si(s) reaction. Al4C3 compound is known to be brittle and unstable, causing degradation of the mechanical properties of the AMCs. It is reactive with H2O in humid conditions and might form amorphous aluminum hydroxide. This process is followed by a volume increase and can induce the residual stresses into the surrounding aluminum matrix. Therefore, the inhibition of the Al4C3 formation is a crucial issue to be overcome [11,34].

Silicon Carbide: SiC
The SiC particle reinforced AMCs are applied in aerospace and electronic encapsulation, both in military and civilian fields, due to their high specific strength and stiffness, in addition to abrasion resistance. SiC has a much higher laser absorptivity (≈78%) than aluminum (≈7%), moderate density (3.21 g/cm 3 ), and it increases the laser absorptivity of the blended mixture [13,34,97,98]. During laser irradiation, SiC particles tend to heat up to extremely high temperature, leading to rapid reaction rates. Hence, the decrease in thermal conductivity results in further rise in temperature, the lifetime and fluidity of the melt pool. Meanwhile, an increase in SiC content in the initial feedstock and, hence, in the blend melt pool, increases the viscosity of a liquid melt and results in a lower fluidity. Therefore, both thermo-kinetic factors should be considered before selecting the content and size of the reinforcing SiC [11,13].
The chemical reaction between silicon carbide and aluminum melt at temperatures exceeding 940 K may result in SiC decomposition according to 4Al(l)+3SiC(s)→Al 4 C 3 (s)+3Si(s) reaction. Al 4 C 3 compound is known to be brittle and unstable, causing degradation of the mechanical properties of the AMCs. It is reactive with H 2 O in humid conditions and might form amorphous aluminum hydroxide. This process is followed by a volume increase and can induce the residual stresses into the surrounding aluminum matrix. Therefore, the inhibition of the Al 4 C 3 formation is a crucial issue to be overcome [11,34].
At a processing temperature above 1670 K, Al 4 SiC 4 (ternary carbide) is formed following the 4Al (l) +4SiC (S) →Al 4 SiC 4(S) +3Si reaction [13]. Al 4 SiC 4 , due to its high hardness of 1200 HV, low brittleness, remarkable chemical stability in wet conditions, is a favored reinforcement for aluminum [11]. At temperatures above 2800 • C, SiC particles partially or fully decompose into silicon and carbon vapor [34,97]. The increase in applied energy results in a high degree of SiC decomposition, causing surface turbulence, melt pool instability, non-continuous scan tracks and, consequently, an uneven surface finish.
It should be noted that the size of used SiC reinforcing particles ranges from tens of micrometers down to nanoscale, and the resultant mechanical properties of AMCs are significantly affected by particle size [8,13]. In Refs. [8,34], the LPBF of AlSi7Mg/2 wt.% nano-SiCp (40 nm) and Al-12Si/10 vol.%SiC (≈11.7 wt.%) (SiC ≈ 25 µm), respectively, were reported. Nano SiC in AlSi7Mg matrix serves as a grain refinement agent (Figure 11m,n) due to the nucleation of numerous heterogenous sites and formation of nanosized Al 4 C 3 (Figure 15b,c). The use of nano-SiC yielded low porosity, near-full densification and improvement in tensile strength without sacrificing ductility. However, inferior densification was observed in Ref. [34] when a micron size reinforcement was used. The successful fabrication of AlSi10Mg/2 vol.% nano-SiC (~2.4 wt.%) composite reinforced by Al4SiC4 phase was reported in Ref. [95]. With an increase in laser power, the eutectic structure gradually changed from thick flakes to network shapes and then to a fine structure, as shown in Figure 16.  The successful fabrication of AlSi10Mg/2 vol.% nano-SiC (~2.4 wt.%) composite reinforced by Al 4 SiC 4 phase was reported in Ref. [95]. With an increase in laser power, the eutectic structure gradually changed from thick flakes to network shapes and then to a fine structure, as shown in Figure 16.
At low applied energy, the eutectic structure represents a collection of thick flakes. In contrast, high energy input provides sufficient wettability between SiC and Al, promoting the reaction product transformation into Al 4 SiC 4 and a homogeneously dispersed eutectic structure (Figure 17), which positively affects the mechanical properties of the AMC. Despite the analogous content of nano SiC added to the Al alloy, the mechanical properties of the samples in this work are far inferior to those reported in Ref. [8].
the illustration of the formation route of different phases during the LPBF process (c) (reproduced with permission from [8]).
The successful fabrication of AlSi10Mg/2 vol.% nano-SiC (~2.4 wt.%) composite reinforced by Al4SiC4 phase was reported in Ref. [95]. With an increase in laser power, the eutectic structure gradually changed from thick flakes to network shapes and then to a fine structure, as shown in Figure 16. At low applied energy, the eutectic structure represents a collection of thick flakes. In contrast, high energy input provides sufficient wettability between SiC and Al, promoting the reaction product transformation into Al4SiC4 and a homogeneously dispersed eutectic structure (Figure 17), which positively affects the mechanical properties of the AMC. Despite the analogous content of nano SiC added to the Al alloy, the mechanical properties of the samples in this work are far inferior to those reported in Ref. [8]. Microstructure changes of the composites at low to high energy application (reproduced with permission from [95]).
An increase in SiC content up to 10 wt.% resulted in increased tensile and yield strength; however, the SiC, Si and in situ formed Al4SiC4, reduce the elongation of the composites [96]. When comparing the properties of AlSi10Mg/15-20 wt.% SiC composites [11,13,97,98], it should be mentioned that the highest hardness (316.2HV0.2) and density (98.9%) were achieved for AlSi10Mg/15 wt.%SiC, when the SiC particle size was 1200 mesh [98] ( Table 2). The larger SiC particles reduced tensile strength as compared to a pure alloy [97]. The use of finer SiC particulates yields to a higher degree of densification, elevated microstructural uniformity and simultaneous improvement in compressive strength, hardness and strain [11,98]. In Refs. [11,13], the in situ formed Al4SiC4 is shown to serve as a transition zone, limiting the interaction of SiC and aluminum crystals simultaneously with reinforcing capacity for the Al.

Titanium Nitride: TiN
Besides the favorable characteristics of ceramic materials, TiN (titanium nitride) also demonstrates excellent light absorptivity. TiN has good coherency with Al, owing to small difference (4.72%) in lattice parameters (aAl = 0.4049 nm and aTiN = 0.4240 nm). Meanwhile, the laser reflectivity (at 1064 nm laser wavelength) of the AlSi10Mg/TiN composite powder is around 25%, which is much lower than that of AlSi10Mg powder (62%) [99].  [95]).
An increase in SiC content up to 10 wt.% resulted in increased tensile and yield strength; however, the SiC, Si and in situ formed Al 4 SiC 4 , reduce the elongation of the composites [96]. When comparing the properties of AlSi10Mg/15-20 wt.% SiC composites [11,13,97,98], it should be mentioned that the highest hardness (316.2HV 0.2 ) and density (98.9%) were achieved for AlSi10Mg/15 wt.%SiC, when the SiC particle size was 1200 mesh [98] ( Table 2). The larger SiC particles reduced tensile strength as compared to a pure alloy [97]. The use of finer SiC particulates yields to a higher degree of densification, elevated microstructural uniformity and simultaneous improvement in compressive strength, hardness and strain [11,98]. In Refs. [11,13], the in situ formed Al 4 SiC 4 is shown to serve as a transition zone, limiting the interaction of SiC and aluminum crystals simultaneously with reinforcing capacity for the Al.

Nitrides: Grain Refinement and Strengthening Effect 2.3.1. Titanium Nitride: TiN
Besides the favorable characteristics of ceramic materials, TiN (titanium nitride) also demonstrates excellent light absorptivity. TiN has good coherency with Al, owing to small difference (4.72%) in lattice parameters (a Al = 0.4049 nm and a TiN = 0.4240 nm). Meanwhile, the laser reflectivity (at 1064 nm laser wavelength) of the AlSi10Mg/TiN composite powder is around 25%, which is much lower than that of AlSi10Mg powder (62%) [99].
In Refs. [99,100], when fabricating AlSi10Mg/2 wt.%TiN composite, the mutual diffusion and in situ reaction between the TiN clusters and aluminum generates a graded interfacial layer composed of Al 3.21 Si 0.47 and (Ti,Al)N ( Figure 18). The formed layer is of central importance to the enhancement in microhardn to an improved interface bonding and a precipitation of stiff (Al,Ti)N. The co influence of superfine grains (0.284 µm), uniform particle dispersion, formed no and high densification significantly improve the mechanical and wear characte the fabricated AMCs. The Al matrix-Mg2Si-TiN coherent interfaces lead to a prec strengthening, benefiting the enhancement in strength [100].
An increase in TiN content (0-6 wt.%) improves strength, ductility and har nano-TiN particle reinforced AlSi10Mg [101]. It was shown that 4 wt.% TiN is threshold to inhibit porosity. The composites had a relatively random grain orie and the grain size decreased from 3.86 to 1.19 µm when the content of TiN increas 0-6 wt.% due to intensive heterogenous nucleation (Figures 19a-d and 20, Table The formed layer is of central importance to the enhancement in microhardness due to an improved interface bonding and a precipitation of stiff (Al,Ti)N. The combined influence of superfine grains (0.284 µm), uniform particle dispersion, formed novel layer and high densification significantly improve the mechanical and wear characteristics of the fabricated AMCs. The Al matrix-Mg 2 Si-TiN coherent interfaces lead to a precipitation strengthening, benefiting the enhancement in strength [100].
An increase in TiN content (0-6 wt.%) improves strength, ductility and hardness of nano-TiN particle reinforced AlSi10Mg [101]. It was shown that 4 wt.% TiN is a critical threshold to inhibit porosity. The composites had a relatively random grain orientation, and the grain size decreased from 3.86 to 1.19 µm when the content of TiN increased from 0-6 wt.% due to intensive heterogenous nucleation (Figures 19a-d and 20, Table 3).
An increase in TiN content (0-6 wt.%) improves strength, ductility and hardness nano-TiN particle reinforced AlSi10Mg [101]. It was shown that 4 wt.% TiN is a criti threshold to inhibit porosity. The composites had a relatively random grain orientati and the grain size decreased from 3.86 to 1.19 µm when the content of TiN increased fro 0-6 wt.% due to intensive heterogenous nucleation (Figures 19a-d and 20, Table 3).  As shown in Figure 20, only a fraction of TiN serves as heterogenous nucleation substrates, and the majority of particles are dispersed along the grain boundaries owing to the pushing effects of the solidification front. It was found that all the specimens were dominated by high-angle grain boundaries (HAGBs), and with an increase in TiN content, the volume of low-energy HAGBs increased. TiN nanoparticles also promote recrystallization and possesses a crucial role in recrystallized nucleation during the LPBF process, as shown in Figure 19e-h.   As shown in Figure 20, only a fraction of TiN serves as heterogenous nucleation substrates, and the majority of particles are dispersed along the grain boundaries owing to the pushing effects of the solidification front.
It was found that all the specimens were dominated by high-angle grain boundaries (HAGBs), and with an increase in TiN content, the volume of low-energy HAGBs in-creased. TiN nanoparticles also promote recrystallization and possesses a crucial role in recrystallized nucleation during the LPBF process, as shown in Figure 19e-h.
The use of hybrid Ti-TiN reinforcements for 7050 Al alloy was reported in Ref. [66], exhibiting significant synergistic grain refinement and a higher strengthening as compared to pure 7050 Al alloy and a single reinforced 7050-TiN and 7050-Ti. Although both single-Tireinforced and hybrid-reinforced alloy possessed a crack-free microstructure (Figure 21g-l), the hybrid reinforcement provided greater grain refinement (Figure 19k,l). Meanwhile, the 7050 and 7050-0.18%TiN specimens are prone to cracking, consist of columnar grains and possess relatively high porosity (Figures 19 i,j and 21a-f). The reason for grain refinement, when Ti is added to pure alloy and to 7050-TiN, is the formation of L12 structured Al3Ti, which promotes heterogeneous nucleation and contributes to the rapid formation of constitutional supercooling zones (Figure 21n). Besides Al3Ti, fine MgZn2 phase was formed with coherent interface with Al; however, the in situ formed Al2CuMg showed non-coherent interface with Al. Ultrafine grains (775 nm) were reported in the LPBF-prepared 7050-2 wt.%(Ti+TiN) composite, vastly benefiting from the Ti/TiN synergism.
It can be concluded that the addition of 2-4 wt.% TiN-Ti hybrid additives notably improved the quality of LPBF-fabricated AMCs.

Aluminum Nitride: AlN
AlN is one of the favorable reinforcing candidates for aluminum alloys due to its superior combination of high thermal conductivity (~250 W/mK) [105] and high hardness The reason for grain refinement, when Ti is added to pure alloy and to 7050-TiN, is the formation of L1 2 structured Al 3 Ti, which promotes heterogeneous nucleation and contributes to the rapid formation of constitutional supercooling zones (Figure 21n). Besides Al 3 Ti, fine MgZn 2 phase was formed with coherent interface with Al; however, the in situ formed Al 2 CuMg showed non-coherent interface with Al. Ultrafine grains (775 nm) were reported in the LPBF-prepared 7050-2 wt.%(Ti+TiN) composite, vastly benefiting from the Ti/TiN synergism.
It can be concluded that the addition of 2-4 wt.% TiN-Ti hybrid additives notably improved the quality of LPBF-fabricated AMCs.

Aluminum Nitride: AlN
AlN is one of the favorable reinforcing candidates for aluminum alloys due to its superior combination of high thermal conductivity (~250 W/mK) [105] and high hardness (~12 GPa) [106]. AlN shows high chemical stability, good compatibility with Al alloy combined with a good interfacial adherence without any interfacial reaction [107]. Besides, due to a low thermal expansion coefficient (similar to Si), AlN has been broadly employed in the aviation and transportation and is shown to be an appropriate reinforcement for aluminum alloys [102].
In a series of works [67,107,108], it was observed that the applied energy had a dramatic effect on the AlN particle distribution. At low energy, random AlN distribution occurred due to the relatively consistent pressure around the introduced particles (Figure 22a,c); and at high laser energy, a circular-structured AlN distribution was compelled by the centripetal force (Figure 22b,d).
Materials 2022, 15, x FOR PEER REVIEW 28 of 41 employed in the aviation and transportation and is shown to be an appropriate reinforcement for aluminum alloys [102].
In a series of works [67,107,108], it was observed that the applied energy had a dramatic effect on the AlN particle distribution. At low energy, random AlN distribution occurred due to the relatively consistent pressure around the introduced particles ( Figure  22a,c); and at high laser energy, a circular-structured AlN distribution was compelled by the centripetal force (Figure 22b,d). However, excessive energy results in particles coarsening and a deconstruction of the circular-structured AlN. In Ref. [58], the preparation of an almost fully densified composite with 1 wt.% AlN and refined grains of increased wear resistance has been reported. In Ref. [102], it was shown that during LPBF of AlSi10Mg-2 wt.%AlN powders mixture, the solidified material undergoes various microstructural transformations from the first to the fourth layer (directional columnar microstructure to coarse cellular microstructure), affirming the importance of added particles, solidification rate, the lifespan of the melt pools and subsequent crystal growth rate. However, excessive energy results in particles coarsening and a deconstruction of the circular-structured AlN. In Ref. [58], the preparation of an almost fully densified composite with 1 wt.% AlN and refined grains of increased wear resistance has been reported. In Ref. [102], it was shown that during LPBF of AlSi10Mg-2 wt.%AlN powders mixture, the solidified material undergoes various microstructural transformations from the first to the fourth layer (directional columnar microstructure to coarse cellular microstructure), affirming the importance of added particles, solidification rate, the lifespan of the melt pools and subsequent crystal growth rate.

Boron Nitride: BN
The high tensile strength and low density (2.1 g/cm 3 , which is close to that of pristine Al), makes hexagonal boron nitride (h-BN) an effective reinforcing agent for the AMCs [109]. It was revealed that even 1 wt.% addition of BN micro-flakes to AlSi10Mg increased the tensile strength and hardness as compared to a pure alloy due to the formation of AlN and AlB 2 phases via solid-state Al-BN reaction [103].

Silicon Nitride: Si 3 N 4
A whole basket of favorable properties of Si 3 N 4 (silicon nitride), including remarkable strength, high hardness, high elastic modulus, lower CTE, superior hardness compared to other ceramics, [110][111][112], similar density with aluminum, which will ensure homogeneous dispersion, and high wettability with the aluminum matrix [104] makes it a promising reinforcing agent. The enhanced strength and elastic modulus of the LPBF-prepared AlSi10Mg-Si 3 N 4 composite, owing to the impeded dislocation motion during deformation and load-bearing effect of added reinforcing Si 3 N 4, are achieved. The mutual diffusion of Al and Si atoms and the absence of in situ formed brittle phases increased the Al matrix-Si 3 N 4 particles bonding strength [104]. The addition of Si 3 N 4 to the Al alloy, however, reduces process stability and thus narrows the optimal range of process parameters [104].

Comparison of Ceramic Reinforcements' Influence on LPBF Process and the Properties of the AMCs
As shown above, even small portions of ceramic or hybrid additives (metal-ceramic), such as 0.5-0.7 wt.%, are able to dramatically improve the performance of the AMCs. Accordingly, matching ceramic additives with an optimized fraction and particle size provides good wettability, compatible interfaces and a strong bonding between the constituents, which hinder crack propagation and contribute to a hardening and strengthening of AMCs.
When analyzing SiC reinforced AlSi10Mg, huge attention was given to applied energy, as under low energy, brittle Al4C3 is formed. However, higher energy promotes the formation of Al4SiC4, along with a well-dispersed eutectic structure, hence prohibiting the premature failure of the composite [95]. Similar to SiC (Figure 24g,h), in Si3N4 reinforced AMC (Figure 24k,l), the nature of the fracture is ductile brittle, dominated by brittle, whereas pure AlSi10Mg (Figure 24i,j) shows a ductile-brittle composite fracture dominated by ductile. Due to Si3N4, crack propagation is suppressed when the tip meets the Si3N4-AlSi0Mg interface. However, because of the irregular distribution of Si3N4 and the changes in propagation path of the connected cracks, more cleavage steps were formed [104]. When TiN nanoparticles are added to AlSi10Mg, the fracture behavior of the alloy remains in mixed failure mode; however, large-size agglomerates formed during excess addition of TiN, decreasing both strength and ductility [101].
The degree of improvement depends on additive content and composition of the Al alloy. Table 4 briefly summarizes the influence of the reported ceramic additives on the LPBF process and their content limitation. Table 4. The effect of reinforcing compounds on the fabrication and properties of AMCs and their optimal content limit.
The tensile fracture of the AlSi10Mg-6.5 wt.%TiB 2 composite showed that the fracture path of the AMC is not flat, as in the case of AlSi10Mg, but rather random for both horizontal and vertical samples (Figure 24a,b) [79]. Generally, the reinforced composites with refined microstructure have high ductility due to less stress concentration. Based on the fine-sized equiaxed dimples (Figure 24e,f), the failure mode of the AMC is a ductile fracture, stating improved ductility. However, the holes and the tears on the fracture surface might have led to premature failure of the AMC (Figure 24c,d). Similarly, in the AlSi10Mg-0.2 wt.%LaB 6 composite, cracking predominantly occurred within the melt pool boundaries, and the LaB 6 nanoparticles led to more ductile fracture of the composite, owing to fine equiaxed dimples [84]. Ductile-type failure was reported for AlSi10Mg with homogeneously dispersed circular-structured TiC (3 wt.%). The latter contributed to the improvement of tensile strength without sacrificing ductility [71]. The dual TiB 2 and TiC reinforced AMC's tensile fracture (Figure 24m,n) possesses fewer pores and deeper dimples as compared to AlSi10Mg (Figure 24o,p) and shows mixed ductile and brittle fracture mode. The relatively hard intragranular TiB 2 and TiC particles accommodate the dislocations in the grains, contributing to strain hardening and uniform elongation [93]. Both brittle and ductile fractures were observed in the case of 0.7 wt.% hybrid Ti-B 4 C addition. However, the further increase in additive content led to fracture changes from ductile to brittle [94].
When analyzing SiC reinforced AlSi10Mg, huge attention was given to applied energy, as under low energy, brittle Al 4 C 3 is formed. However, higher energy promotes the formation of Al 4 SiC 4 , along with a well-dispersed eutectic structure, hence prohibiting the premature failure of the composite [95]. Similar to SiC (Figure 24g,h), in Si 3 N 4 reinforced AMC (Figure 24k,l), the nature of the fracture is ductile brittle, dominated by brittle, whereas pure AlSi10Mg (Figure 24i,j) shows a ductile-brittle composite fracture dominated by ductile. Due to Si 3 N 4 , crack propagation is suppressed when the tip meets the Si 3 N 4 -AlSi0Mg interface. However, because of the irregular distribution of Si 3 N 4 and the changes in propagation path of the connected cracks, more cleavage steps were formed [104]. When TiN nanoparticles are added to AlSi10Mg, the fracture behavior of the alloy remains in mixed failure mode; however, large-size agglomerates formed during excess addition of TiN, decreasing both strength and ductility [101].
Analyses show that the highest hardness was shown by 15 wt.%SiC reinforced AMCs, followed by the 17.2 wt.% hybrid B 4 C-Ti and 11.6 wt.%TiB 2 reinforced materials (Figure 25a). Hardness values of TiC and Si 3 N 4 reinforced AMCs are comparable with TiB 2 . Meanwhile, ceramic reinforced 2024, Al-12Si and Al-Cu alloys show inferior hardness compared to AlS10Mg with similar additives (Figure 25b).
The AMCs reinforced with TiB 2 , TiC, hybrid TiN-Ti and TiC-TiH 2 additives are subjected to in situ formation of L1 2 -Al 3 Ti or D0 22 -Al 3 Ti (Table 4), which serve as active nucleation sites and promote grain refinement in the 0.5-2 µm range (Figure 26a,b). The substantial grain refinement, down to submicron level, is achieved by the incorporation of TiN and CaB 6 into AMCs, resulting in both significantly enhanced hardness and tensile strength (Figure 26a,b). Table 4. The effect of reinforcing compounds on the fabrication and properties of AMCs and their optimal content limit.

SiC
Use of fine (nanosized or few-micron-sized) SiC results in grain refinement, decrease in porosity, enhancement of hardness, tensile strength and ductility but, depending on the process parameters, can cause in situ formation of Al 4 C 3 or Al 4 SiC 4 phases.
Up to 2 wt.% The degree of improvement depends on additive content and composition of the Al alloy. Table 4 briefly summarizes the influence of the reported ceramic additives on the LPBF process and their content limitation.

Summary and Outlook
LPBF technologies are now commercially available and attract a huge deal of attention in research community. Although the number of aluminum alloys suitable for AM through LPBF is quite limited, the process keeps evolving, and, in the nearest future, a widespread application of AM of high-strength aluminum alloys is expected to occur in the aerospace market.
The cost of industrial metal printers remains the chief capital expenditure of AM parts to achieve economies-of-scale cost reduction. Although the industry has suffered due to COVID-19, the reverse has now begun. In light of current metal printers' high prices, they are mostly used in high-value industries, such as aerospace, defense and medical.
Other fields, such as energy, are starting to show interest in powder bed fusion technology, although developing economically viable applications requires sufficient time.
A 2.6 percent annual growth rate is predicted for aluminum consumption globally up to 2029. In 2021, global aluminum consumption is projected at 64.2 million metric tons alone ( Figure 27). However, fuel efficiency and low carbon emission are the mantra for new-era airliners, which have groundbreaking design equipped with composite materials comprising 50 percent of the primary structure, hence eliminating the use of numerous aluminum parts [114]. In addition, the world's biggest aluminum producers are limiting the production of Al, planning to reduce energy consumption and encourage the producers to develop green and low-carbon technologies and produce high-quality, highstrength and long-life aluminum products through innovations [115]. This means that there is a need for revolutionary actions to keep additive manufacturing of aluminum alloys on track.
Over the next decade, the development of new 3D printable Al alloys is expected to bring down the cost and enlarge the materials' capacity and portfolio. For example, the lightweight aluminum-lithium alloys could contribute to reducing aircraft weight, also benefiting from excellent fatigue resistance and cryogenic toughness in addition to light weight and high specific modulus.
As numerous reinforcements are used to further enhance the properties of Al alloys, one big step ahead will be using different reinforcing particles (ceramics) and covering them with compatible coatings to provide suitable wettability and interface, or incorporating the reinforcing particles into Al alloy particles to provide a homogeneous distribution. Another main challenge is the recycling of the used feedstock and the utilization of the spattered debris to prepare new powders for further use.
As the design of new alloys applicable for the LPBF process is time and cost consuming, a high-throughput and reliable technique is needed to experimentally validate the custom alloys and effectively introduce them into the market. Therefore, a deep understanding of the impact of the alloying constituents on the processability of the feedstock by LPBF and, ultimately, the properties of the produced items in application, is of a crucial importance.
In this review paper, the effect of non-oxide ceramic (borides, nitrides, carbides) and hybrid reinforcing additives on the densification, grain refinement and respective mechanical characteristics of LPBF-fabricated AMCs was discussed. A comprehensive analysis of research studies on densification, compositional and microstructural characteristics of the in situ and ex situ reinforced aluminum alloys produced by LPBF method was accomplished to demonstrate the capability of different ceramic additives to tailor the mechanical properties with application to a wide variety of process parameters.

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Generally, an incorporation of the ceramic particles into Al alloys results in a significant improvement in strength, ductility and hardness of the fabricated parts accompanied by a refined microstructure and with randomization of crystallographic orientation of reinforced AMCs. However, fuel efficiency and low carbon emission are the mantra for new-era airliners, which have groundbreaking design equipped with composite materials comprising 50 percent of the primary structure, hence eliminating the use of numerous aluminum parts [114]. In addition, the world's biggest aluminum producers are limiting the production of Al, planning to reduce energy consumption and encourage the producers to develop green and low-carbon technologies and produce high-quality, high-strength and long-life aluminum products through innovations [115]. This means that there is a need for revolutionary actions to keep additive manufacturing of aluminum alloys on track.
Over the next decade, the development of new 3D printable Al alloys is expected to bring down the cost and enlarge the materials' capacity and portfolio. For example, the lightweight aluminum-lithium alloys could contribute to reducing aircraft weight, also benefiting from excellent fatigue resistance and cryogenic toughness in addition to light weight and high specific modulus.
As numerous reinforcements are used to further enhance the properties of Al alloys, one big step ahead will be using different reinforcing particles (ceramics) and covering them with compatible coatings to provide suitable wettability and interface, or incorporating the reinforcing particles into Al alloy particles to provide a homogeneous distribution. Another main challenge is the recycling of the used feedstock and the utilization of the spattered debris to prepare new powders for further use.
As the design of new alloys applicable for the LPBF process is time and cost consuming, a high-throughput and reliable technique is needed to experimentally validate the custom alloys and effectively introduce them into the market. Therefore, a deep understanding of the impact of the alloying constituents on the processability of the feedstock by LPBF and, ultimately, the properties of the produced items in application, is of a crucial importance.
In this review paper, the effect of non-oxide ceramic (borides, nitrides, carbides) and hybrid reinforcing additives on the densification, grain refinement and respective mechanical characteristics of LPBF-fabricated AMCs was discussed. A comprehensive analysis of research studies on densification, compositional and microstructural characteristics of the in situ and ex situ reinforced aluminum alloys produced by LPBF method was accomplished to demonstrate the capability of different ceramic additives to tailor the mechanical properties with application to a wide variety of process parameters.

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Generally, an incorporation of the ceramic particles into Al alloys results in a significant improvement in strength, ductility and hardness of the fabricated parts accompanied by a refined microstructure and with randomization of crystallographic orientation of reinforced AMCs.

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Most of the AMCs can be densified to over 99% relative density; moreover, non-oxide ceramic additives significantly improve laser absorptivity of a powder feedstock.

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The addition of ceramic particulates shifts the process window to a higher energy regime; however, an applied excess energy may result in the evaporation or decomposition of ceramics particles (mainly SiC).

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The application of a laser re-melting strategy can further increase the densification degree and the surface quality of AMCs; however, it also can cause the evaporation and loss of ceramic particles. • Hybrid reinforcements are proven to be the effective additives, providing the formation of a wide variety of reinforcing phases with a coherent interface with matrices.

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The use of ceramics with a fine-particle size results in an increased degree of densification, microstructural and compositional uniformity, as well as an apparent grain refinement.

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The addition of TiB 2 , CaB 6 , TiC, TiN to Al alloys leads to a considerable grain refinement, down to the submicron level, due to the intensive heterogeneous nucleation and grain growth inhibition. • An addition of matching ceramics prevents the hot tearing and gives the prospect to consolidate crack-susceptible Al alloys by a laser powder-bed fusion technique.

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The highest elongation of 17.7% is demonstrated by the AlSi10Mg/TiB 2 composite; however, the highest strength of 613 MPa is recorded for the hybrid TiN-Ti reinforced AMCs.

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The highest hardness of 316 HV is estimated for SiC reinforced AMCs, which possess a relatively high strength and moderate ductility.