**3. Ex situ MMCs produced by laser additive manufacturing**

### **3.1. Aluminium matrix composites**

#### *3.1.1. Aluminium matrix composites reinforced with carbides*

A major concern during the fabrication of aluminium (Al) matrix composites reinforced with carbide particles is to avoid the formation of aluminium carbide Al4C3 through interfacial reactions between the liquid Al and the particles [6, 13, 42]. Indeed, the presence of Al4C3 in the composite is potentially detrimental to both its mechanical and its corrosion resistance, as it is well known that Al4C3 is extremely brittle [13] and susceptible to hydrolysis in humid environment, thus giving rise to high corrosion rate [6]. A number of different strategies have been devised with various levels of success in order to solve this problem, including the careful selection of the carbide particles used as reinforcement [2, 13], the chemical modification of the matrix alloy [6, 12] or the optimisation of the processing conditions, e.g. to promote the formation of less deleterious aluminium carbides such as the mixed aluminium-silicon carbide Al4SiC4 [2, 19, 43, 44].

In their investigation into the fabrication by DMLS of AlSi10Mg matrix composites with 10 wt % of SiC particles, Manfredi et al. [12] relied on the earlier observation that the addition of a high amount (i.e. of at least 7 at% [42]) of silicon in the matrix alloy could suppress the reaction between the Al alloy and the SiC particles leading to the formation of Al4C3 during the fabrication of Al matrix composites by conventional processes such as melt infiltration. The experimental parameters for the DMLS process were set as follows: the scan speed was varied between 500 and 700 mm/s, the laser power between 180 and 195 W, the layer thickness was set at 30 μm and the hatching distance at 0.17 mm. Although the hardness of the composite materials was significantly increased in comparison with AlSi10Mg processed in the same conditions, analysis of the AlSi10Mg/SiC composites by X-Ray Diffraction (XRD) showed that the SiC particles had almost fully disappeared while a significant amount of Al4C3 had been formed [12]. The behaviour of the DMLS AlSi10Mg/SiC composites was thus found to differ significantly from their conventional counterparts.

Also investigating the fabrication of AlSi10Mg/SiC composites by SLM, Chang et al. [2] used the same layer thickness of 30 μm as used by Manfredi et al. [12], with a smaller hatch spacing of 0.05 mm, a lower laser power of 100 W and a lower scan speed of 100 mm/s. Dissolution of the SiC particles was also observed in their study, but this time, the reaction product was the mixed Al4SiC4 carbide, that has the advantages of being chemically inert in humid environment and of being less brittle than Al4C3. By considering the energy density per unit volume for both studies, it is found that the set of experimental parameters used by Manfredi et al. [12] resulted in a much lower value of the energy density of 77 J/mm3 when compared to the set of processing parameters of Chang et al. [2] giving rise to an energy density of 666 J/mm3 . As a consequence, the samples of Chang et al. [2] experienced a greater heat accumulation during fabrication and a higher temperature when compared with the samples of Manfredi et al. [12]. The processing strategy of Chang et al. [2] thus proved successful in favouring the formation of Al4SiC4 that is expected to form at temperatures above 1350–1400°C, whereas the formation of the delete‐ rious Al4C3 cannot be avoided at lower processing temperatures [2, 19], in spite of the high silicon content of the matrix alloy [12].

major characteristics of these MMCs is proposed in Section 5. Special care is also taken to identify current fundamental issues that should more particularly be the object of future work.

A major concern during the fabrication of aluminium (Al) matrix composites reinforced with carbide particles is to avoid the formation of aluminium carbide Al4C3 through interfacial reactions between the liquid Al and the particles [6, 13, 42]. Indeed, the presence of Al4C3 in the composite is potentially detrimental to both its mechanical and its corrosion resistance, as it is well known that Al4C3 is extremely brittle [13] and susceptible to hydrolysis in humid environment, thus giving rise to high corrosion rate [6]. A number of different strategies have been devised with various levels of success in order to solve this problem, including the careful selection of the carbide particles used as reinforcement [2, 13], the chemical modification of the matrix alloy [6, 12] or the optimisation of the processing conditions, e.g. to promote the formation of less deleterious aluminium carbides such as the mixed aluminium-silicon carbide

In their investigation into the fabrication by DMLS of AlSi10Mg matrix composites with 10 wt % of SiC particles, Manfredi et al. [12] relied on the earlier observation that the addition of a high amount (i.e. of at least 7 at% [42]) of silicon in the matrix alloy could suppress the reaction between the Al alloy and the SiC particles leading to the formation of Al4C3 during the fabrication of Al matrix composites by conventional processes such as melt infiltration. The experimental parameters for the DMLS process were set as follows: the scan speed was varied between 500 and 700 mm/s, the laser power between 180 and 195 W, the layer thickness was set at 30 μm and the hatching distance at 0.17 mm. Although the hardness of the composite materials was significantly increased in comparison with AlSi10Mg processed in the same conditions, analysis of the AlSi10Mg/SiC composites by X-Ray Diffraction (XRD) showed that the SiC particles had almost fully disappeared while a significant amount of Al4C3 had been formed [12]. The behaviour of the DMLS AlSi10Mg/SiC composites was thus found to differ

Also investigating the fabrication of AlSi10Mg/SiC composites by SLM, Chang et al. [2] used the same layer thickness of 30 μm as used by Manfredi et al. [12], with a smaller hatch spacing of 0.05 mm, a lower laser power of 100 W and a lower scan speed of 100 mm/s. Dissolution of the SiC particles was also observed in their study, but this time, the reaction product was the mixed Al4SiC4 carbide, that has the advantages of being chemically inert in humid environment and of being less brittle than Al4C3. By considering the energy density per unit volume for both studies, it is found that the set of experimental parameters used by Manfredi et al. [12] resulted

the samples of Chang et al. [2] experienced a greater heat accumulation during fabrication and

when compared to the set of processing

. As a consequence,

**3. Ex situ MMCs produced by laser additive manufacturing**

**3.1. Aluminium matrix composites**

192 New Trends in 3D Printing

Al4SiC4 [2, 19, 43, 44].

*3.1.1. Aluminium matrix composites reinforced with carbides*

significantly from their conventional counterparts.

in a much lower value of the energy density of 77 J/mm3

parameters of Chang et al. [2] giving rise to an energy density of 666 J/mm3

Ghosh et al. [43, 44], on the other hand, investigated the fabrication of Al-4.5Cu-3Mg/SiC composites by DMLS using a pulsed Nd:YAG laser while varying the SiC volume fraction and particles size. However, due to the strongly non-uniform nature of the temperature distribu‐ tion created by the pulsed laser, their composites were subjected to high internal stresses that resulted in extensive cracking in all their specimens [44, 45].

In a similar way to silicon carbides, tungsten carbides (WC) present a strong tendency to dissolve in Al alloys during LD, leading to the formation of Al–W intermetallics and of the detrimental Al4C3 [6]. In order to avoid the formation of the latter phase, Li et al. [6] attempted to modify the chemical composition of their Al-11 wt% Si matrix alloy by adding pure elemental Ti powder together with the WC particles. The addition of Ti in the melt pool was found successful in inhibiting the formation of the deleterious Al4C3 by promoting the precipitation of titanium carbides (TiC) both as dispersed phase within the Al matrix and as a passivation layer at the interface between the WC particles and the Al matrix. Mixed Al–Ti (Al3Ti or AlTi) and Al–W (Al12W or Al4W) intermetallic compounds were also detected by XRD.

On the other hand, titanium carbides (TiC) present a much lower reactivity with Al than SiC and WC [13]. Indeed, TiC possesses a low solubility in Al, and its dissolution rate is also slower than for SiC. Coatings with a matrix of Al-12 wt% Si and 40 wt% TiC as reinforcement were successfully produced by LC [13]. Very limited dissolution of the TiC particles into the Al matrix led to the formation of a small amount (about 3%) of finely dispersed Ti3SiC2, whereas the formation of the deleterious Al4C3 was completely avoided. Both the hardness and the sliding wear resistance of the composite coating were significantly improved by the TiC addition.

In view of their excellent thermodynamic stability, TiC nanoparticles have also been used as reinforcement in AlSi10Mg matrix composites [46–49]. Due to their strong tendency to form clusters and agglomerates [22, 29], nanoparticles may reflect poorly on the flowability and/or the spreadability of the powders, and nanostructured or nanocomposite powders of micronsize are usually preferred for the additive manufacturing of nanocomposites. In various studies [46–49], careful mixing of TiC nanoparticles with an average size of 50 nm and of AlSi10Mg particles with an average size of 30 μm in a planetary ball mill resulted in the uniform distribution of the TiC nanoparticles around the surface of the AlSi10Mg particles, thus forming nanostructured particles suitable for further processing by SLM.

In a first set of studies [46, 47], the effect of the linear laser energy density on parts density, microstructure, hardness and wear behaviour was investigated by adjusting the scan speed at a constant laser power of 110 W. The second phases present in the composites were identified as TiC and Mg2Si, the latter formed by natural aging of the AlSi10Mg matrix alloy. Moreover, the highest investigated linear laser energy density of 733 J/m was found to result in parts with an optimum relative density of 98%, uniform distribution of the TiC nanoparticles throughout the composite and optimised hardness and wear resistance. Dai and Gu [48] reported on the optimisation of the surface finish of AlSi10Mg/TiC nanocomposites by varying the scan speed at a constant laser power of 150 W. In this case, medium values of the linear laser energy density gave the best result. Lower linear laser energy densities resulted in an excessive viscosity and poor spreading of the melt pool on the underlying solid surface [46, 47]. Higher linear laser energy density, on the other hand, led to excessive vaporisation of the melt [48]. Both situations resulted in a smaller melt pool and in a poorer surface finish. Finally, in an attempt to optimise the tensile properties of TiC/AlSi10Mg nanocomposites produced by SLM, Gu et al. [49] reported on an investigation of the influence of the processing parameters on the dispersion of the nanoparticles. Gu et al. [49] demonstrated the feasibility of tailoring the spatial distri‐ bution of the TiC nanoparticles in the composite by means of a careful control of the balance of the torque forces in the melt pool and of repulsive capillary forces arising between the nanoparticles under specific processing conditions.

#### *3.1.2. Aluminium matrix composites reinforced with borides*

In opposition to carbides and particularly to SiC or WC whose reactions with liquid Al may result in the formation of brittle and deleterious compounds, titanium diboride (TiB2) is attracting a growing interest for use as reinforcement in Al matrix composites due to its very low reactivity with Al. Anandkumar et al. [19] thus investigated the LC of an Al-12 wt% Si alloy with 40 wt% TiB2. Using processing parameters similar to those reported for the LC of Al-12 wt% Si/TiC composites [13], the Al-12 wt% Si/TiB2 composites did not present any sign of dissolution of the TiB2 particles, nor of interfacial reactions between the particles and the Al matrix [19]. Moreover, when compared with the matrix alloy, the wear resistance of the Al-12 wt% Si/TiB2 composite coating, as characterised by dry sliding wear tests using a counterbody of quenched and tempered AISI440C tool steel, was greatly increased by the addition of the TiB2 particles. Being much harder than the counterbody (2550 HV vs 800 HV), protruding TiB2 particles proved very efficient in supporting the contact stresses with the counterbody and protecting the Al-12 wt% Si matrix from continued intense plastic deformation.

#### *3.1.3. Aluminium matrix composites reinforced with oxides*

Little information is available in literature on the laser additive manufacturing of Al matrix composites with ex situ oxide particles. Manfredi et al. [12] investigated the fabrication of AlSi10Mg matrix composites with nanoparticles of the aluminium-magnesium spinel MgAl2O4 using DMLS, with limited success. In order to avoid processability issues caused by the tendency of nanoparticles for clustering [22, 29], a nanostructured powder was prepared by mixing the AlSi10Mg and the MgAl2O4 powders using a ball milling system for 48 hours prior to conducting the DMLS experiments. Even the best of the produced composite speci‐ mens still exhibited a non-negligible volume fraction of residual porosity ranging from 2.2 to 3.5%. The microstructure of the composites was also found to be much more inhomogeneous than for the AlSi10Mg alloy processed under the same conditions without spinel nanoparticles, thus suggesting that the spinel nanoparticles might affect the solidification process during DMLS. And finally, the hardness of the AlSi10Mg/spinel nanoparticles composite was decreased by 11% in comparison with the AlSi10Mg DMLS samples without nanoparticles.

#### **3.2. Ferrous matrix composites**

an optimum relative density of 98%, uniform distribution of the TiC nanoparticles throughout the composite and optimised hardness and wear resistance. Dai and Gu [48] reported on the optimisation of the surface finish of AlSi10Mg/TiC nanocomposites by varying the scan speed at a constant laser power of 150 W. In this case, medium values of the linear laser energy density gave the best result. Lower linear laser energy densities resulted in an excessive viscosity and poor spreading of the melt pool on the underlying solid surface [46, 47]. Higher linear laser energy density, on the other hand, led to excessive vaporisation of the melt [48]. Both situations resulted in a smaller melt pool and in a poorer surface finish. Finally, in an attempt to optimise the tensile properties of TiC/AlSi10Mg nanocomposites produced by SLM, Gu et al. [49] reported on an investigation of the influence of the processing parameters on the dispersion of the nanoparticles. Gu et al. [49] demonstrated the feasibility of tailoring the spatial distri‐ bution of the TiC nanoparticles in the composite by means of a careful control of the balance of the torque forces in the melt pool and of repulsive capillary forces arising between the

In opposition to carbides and particularly to SiC or WC whose reactions with liquid Al may result in the formation of brittle and deleterious compounds, titanium diboride (TiB2) is attracting a growing interest for use as reinforcement in Al matrix composites due to its very low reactivity with Al. Anandkumar et al. [19] thus investigated the LC of an Al-12 wt% Si alloy with 40 wt% TiB2. Using processing parameters similar to those reported for the LC of Al-12 wt% Si/TiC composites [13], the Al-12 wt% Si/TiB2 composites did not present any sign of dissolution of the TiB2 particles, nor of interfacial reactions between the particles and the Al matrix [19]. Moreover, when compared with the matrix alloy, the wear resistance of the Al-12 wt% Si/TiB2 composite coating, as characterised by dry sliding wear tests using a counterbody of quenched and tempered AISI440C tool steel, was greatly increased by the addition of the TiB2 particles. Being much harder than the counterbody (2550 HV vs 800 HV), protruding TiB2 particles proved very efficient in supporting the contact stresses with the counterbody

and protecting the Al-12 wt% Si matrix from continued intense plastic deformation.

Little information is available in literature on the laser additive manufacturing of Al matrix composites with ex situ oxide particles. Manfredi et al. [12] investigated the fabrication of AlSi10Mg matrix composites with nanoparticles of the aluminium-magnesium spinel MgAl2O4 using DMLS, with limited success. In order to avoid processability issues caused by the tendency of nanoparticles for clustering [22, 29], a nanostructured powder was prepared by mixing the AlSi10Mg and the MgAl2O4 powders using a ball milling system for 48 hours prior to conducting the DMLS experiments. Even the best of the produced composite speci‐ mens still exhibited a non-negligible volume fraction of residual porosity ranging from 2.2 to 3.5%. The microstructure of the composites was also found to be much more inhomogeneous than for the AlSi10Mg alloy processed under the same conditions without spinel nanoparticles, thus suggesting that the spinel nanoparticles might affect the solidification process during

nanoparticles under specific processing conditions.

194 New Trends in 3D Printing

*3.1.2. Aluminium matrix composites reinforced with borides*

*3.1.3. Aluminium matrix composites reinforced with oxides*

#### *3.2.1. Ferrous matrix composites reinforced with carbides*

Silicon carbides have been extensively used as reinforcement in mild [10], medium-carbon [50] or stainless steels [11, 51–53]. A compositionally graded mild steel/SiC composite coating produced by LC was found to exhibit enhanced hardness, wear and corrosion resistance when compared to the bare mild steel substrate [10]. The improvement of the mechanical properties was ascribed not only to the addition of the SiC particles but also to the secondary precipitation of iron silicide Fe2Si consequent to the partial dissolution of the SiC particles. Depending on the processing conditions, Fe3Si or Fe3Si5 have also been observed [50, 51], and, when SiC particles were dispersed in stainless steel, chromium carbides such as Cr3C2 [11, 51, 53] or Cr7C3 [52] were also formed. This is a major issue for the laser processing of stainless steel/ carbides composites, since an excessive precipitation of chromium carbides might deplete the chromium content of the stainless steel matrix hence compromising its corrosion resistance [3]. However, provided the precipitation of chromium carbides remains limited, Dutta Majumdar and Li [53] have demonstrated that stainless steel 316L/SiC composites could achieve a corrosion resistance that is equivalent or better when compared with conventional stainless steel 316L, in combination with improved hardness and wear resistance [11, 53].

Tunsgten carbides (WC) have also been widely used as reinforcement in steels, and like SiC, they exhibit a strong tendency to dissolve in the ferrous matrix [3, 35, 51, 54], giving way to the secondary precipitation of finely dispersed mixed carbide phases as, e.g. M6C, M23C6 or M7C3 where M stands for Fe, W or Cr (in stainless steel matrix composites). WC particles were found extremely efficient in increasing the hardness and wear resistance of steel matrix composites due to the combined hardening effect of the surviving particles, of the finely dispersed secondary precipitates and of the solid solution strengthening of W dissolved in the steel matrix [3, 35, 51]. Since it is known that the uniform distribution of very fine hard phases may be beneficial in view of improving the cavitation erosion resistance of stainless steel, Lo et al. [35] added deliberately small (~1μm) WC particles in a stainless steel 316 composite coating, so as to favour the complete dissolution of the WC particles and the reprecipitation of finely dispersed carbides. However, as mentioned earlier, secondary precipitation of mixed carbides may prove detrimental to the corrosion resistance of stainless steel/WC composites if the depletion of chromium in the stainless steel matrix reach such an extent as to reduce the capacity of the material to withstand corrosion, as observed by Betts et al. [3] in the case of a stainless steel/WC coating.

Titanium carbides (TiC), on the contrary, are characterised by a great chemical stability in presence of ferrous melts. Cheng et al. [51], for example, did not observe any significant dissolution or reaction during the LC of a stainless steel/TiC composite coating. However, decreasing the size of the TiC particles could result in their partial dissolution [55]. Due to the difference in density between TiC and stainless steel, the TiC particles tend to float at the surface of the stainless steel melt pool. As a consequence, special care must be taken when injecting the TiC particles in the melt, for fear that the TiC particles may distribute unevenly in the composite [51, 55]. When added to a Fe-36% Ni Invar36 alloy, on the other hand, TiC particles dissolved at least partially [56]. Enrichment of the Invar36 matrix with Ti resulted in the stabilisation of the body-centred cubic (BCC) α structure at the expense of the face-centred cubic (FCC) γ phase, significant increase of the coefficient of thermal expansion of the matrix alloy and the loss of its "invar" property whereas the wear resistance of the composite exhibited very little improvement.

Chromium carbides, finally, tend to dissolve at least partially in ferrous matrix, giving rise to the secondary precipitation of small amount of mixed M7C3 or M23C6 carbides [3, 51]. Cr3C2 particles did not appear as efficient as WC particles in view of enhancing the hardness and wear resistance of stainless steel matrix composites, but they also proved less deleterious to the corrosion resistance [3]. Although secondary precipitation in stainless steel/Cr3C2 compo‐ site might still have caused some depletion of Cr in the stainless steel matrix, this phenomenon was less severe than in stainless steel/WC composites.

#### *3.2.2. Ferrous matrix composites reinforced with oxides*

Attempts to reinforce ferrous matrix composites with oxide particles have met little success so far. In their investigation into the deposition of a stainless steel AISI304/Al2O3 composite coating by LC, Xu et al. [34] found that the Al2O3 particles were melted upon heating in the laser beam. Al atoms liberated by the dissociation of Al2O3 then dissolved in the stainless steel matrix, with positive effect on the hardness of the coating due to the combined effect of the solid solution strengthening of the austenitic stainless steel matrix by Al and of a partial transformation of the austenite into hard martensite. Betts et al. [3] also reported an extensive dissociation of Al2O3 particles during the direct LD of AISI316/Al2O3 composite coating. Dissociation of Al2O3 was complete at higher laser power, whereas the lower laser power resulted only in a low survival rate of the Al2O3. The hardness of the AISI316/Al2O3 composite was actually decreased in comparison with the AISI316 base material. The LC of Cr2O3 in an austenitic stainless steel matrix [51] also proved unsatisfactory: the composite coating exhib‐ ited extensive cracking and a poor bonding with the stainless steel substrate.

#### *3.2.3. Ferrous matrix composites reinforced with other types of particles*

The LC of chromium borides CrB2 in austenitic stainless steel UNS S31603 resulted in a significant increase of the composite layer hardness and of its cavitation erosion resistance, in comparison with the base material [51]. Improved properties were linked not only to the presence of surviving CrB2 particles but also to the secondary precipitation of finely dispersed chromium borides (CrB and Cr2B) and iron borides (Fe2B) following the partial dissolution of the initial CrB2 particles. Li et al. [57] added CeO2 nanopowder in laser-cladded FeCrBSi/NbC composites: CeO2 nanoparticles acted as heterogeneous nuclei, favouring higher solidification rates and enhancing the in situ synthesis of finely dispersed NbC (see also Section 4.2).

Hydroxyapatite (HA), in micro- or nano-size, has been used as addition in stainless steel matrix composites produced by SLM for biomedical applications [58, 59]. Stainless steel 316L presents good mechanical and chemical properties for use in load-bearing implants, but it does not promote bone adhesion or tissue regrowth on the implant. HA, on the other hand, lacks in ductility and toughness for load-bearing applications, but it is very efficient in favouring bone and tissue regrowth. By combining the properties of the two materials, load-bearing and bioactive implants have been fabricated. Since the density of HA is significantly lower than the density of SS 316L, the ceramic particles tend to float at the surface of the melt pool and to aggregate, resulting in insufficient wetting or "balling", and in the increased formation of porosities and cracks in the SS316L/HA composites. Consequently, a careful optimisation of the SLM processing parameters was necessary in order to avoid these problems. Optimised SLM SS316L/nano-HA composites were found to exhibit a higher tensile strength than their bulk SS316L counterparts, in combination with a good ductility [59].

### **3.3. Nickel and nickel alloys matrix composites**

surface of the stainless steel melt pool. As a consequence, special care must be taken when injecting the TiC particles in the melt, for fear that the TiC particles may distribute unevenly in the composite [51, 55]. When added to a Fe-36% Ni Invar36 alloy, on the other hand, TiC particles dissolved at least partially [56]. Enrichment of the Invar36 matrix with Ti resulted in the stabilisation of the body-centred cubic (BCC) α structure at the expense of the face-centred cubic (FCC) γ phase, significant increase of the coefficient of thermal expansion of the matrix alloy and the loss of its "invar" property whereas the wear resistance of the composite exhibited

Chromium carbides, finally, tend to dissolve at least partially in ferrous matrix, giving rise to the secondary precipitation of small amount of mixed M7C3 or M23C6 carbides [3, 51]. Cr3C2 particles did not appear as efficient as WC particles in view of enhancing the hardness and wear resistance of stainless steel matrix composites, but they also proved less deleterious to the corrosion resistance [3]. Although secondary precipitation in stainless steel/Cr3C2 compo‐ site might still have caused some depletion of Cr in the stainless steel matrix, this phenomenon

Attempts to reinforce ferrous matrix composites with oxide particles have met little success so far. In their investigation into the deposition of a stainless steel AISI304/Al2O3 composite coating by LC, Xu et al. [34] found that the Al2O3 particles were melted upon heating in the laser beam. Al atoms liberated by the dissociation of Al2O3 then dissolved in the stainless steel matrix, with positive effect on the hardness of the coating due to the combined effect of the solid solution strengthening of the austenitic stainless steel matrix by Al and of a partial transformation of the austenite into hard martensite. Betts et al. [3] also reported an extensive dissociation of Al2O3 particles during the direct LD of AISI316/Al2O3 composite coating. Dissociation of Al2O3 was complete at higher laser power, whereas the lower laser power resulted only in a low survival rate of the Al2O3. The hardness of the AISI316/Al2O3 composite was actually decreased in comparison with the AISI316 base material. The LC of Cr2O3 in an austenitic stainless steel matrix [51] also proved unsatisfactory: the composite coating exhib‐

The LC of chromium borides CrB2 in austenitic stainless steel UNS S31603 resulted in a significant increase of the composite layer hardness and of its cavitation erosion resistance, in comparison with the base material [51]. Improved properties were linked not only to the presence of surviving CrB2 particles but also to the secondary precipitation of finely dispersed chromium borides (CrB and Cr2B) and iron borides (Fe2B) following the partial dissolution of the initial CrB2 particles. Li et al. [57] added CeO2 nanopowder in laser-cladded FeCrBSi/NbC composites: CeO2 nanoparticles acted as heterogeneous nuclei, favouring higher solidification rates and enhancing the in situ synthesis of finely dispersed NbC (see also Section 4.2).

ited extensive cracking and a poor bonding with the stainless steel substrate.

*3.2.3. Ferrous matrix composites reinforced with other types of particles*

very little improvement.

196 New Trends in 3D Printing

was less severe than in stainless steel/WC composites.

*3.2.2. Ferrous matrix composites reinforced with oxides*

Carbides, and particularly WC, are by far the most widely used ceramic reinforcements in Nibased composites. WC actually presents a lower reactivity with Ni-based alloys than with ferrous alloys [54]. As a consequence, nickel is sometimes used to coat ceramic powder particles prior to their use as reinforcement, to slow down the dissolution of the ceramic particles in, e.g. ferrous matrix composites [3, 35, 60] and to lower the laser absorptivity of the ceramic to values similar to the laser absorptivity of metallic materials [5]. Due to the low reactivity of WC with Ni-based alloys, it is fairly easy to control the rate of dissolution of the WC particles in Ni-matrix composites by optimising the processing parameters (e.g. by decreasing the laser power) [5, 61]. Partial dissolution of WC in Ni alloys has been reported, leading to secondary precipitation of various types of tungsten carbides (secondary WC, W2C...) as well as of mixed carbides (M7C3, M23C6…) [5, 60, 62]. Ni–W intermetallics (Ni2W4C, Ni17W3…) were also observed in NiCrBSi/WC composites processed by LC at high laser power and on a substrate pre-heated up to 400°C [63]. MMCs reinforced with WC have been successfully produced in a wide range of Ni-based alloys: pure Ni [60, 64], NiCrSiFe [65], NiCrBSi alloys [62, 63], Ni35 (containing Cr, Si, B and Fe) [61] or Ni60 (containing also Co) [66].

Addition of chromium carbide Cr3C2 in alloy Inconel 625 gave good results in terms of wear resistance, due to the partial dissolution of Cr3C2 in the metallic matrix and reprecipitation of fine Cr-rich M7C3 carbides [8]. Titanium carbides (TiC) were also deemed very promising for use as reinforcement in Inconel 625 [67] and Inconel 690 [68]. Minimal dissolution of TiC was observed in Inconel 690 [68], leading to the formation of fine dendritic M23C6.

Other reinforcements suitable for use in Ni-based matrix composites include titanium diboride (TiB2) [26], and rare earth oxides like CeO2 [17] or Y2O3 nanoparticles [69]. Both types of oxide nanoparticles were found very efficient in refining the microstructure and in improving the corrosion resistance of the composites. Hydroxyapatite also proved interesting for biomedical applications, e.g. as second phase in nitinol (i.e. NiTi intermetallic alloy) matrix composites [70].
