Research Progress on Microstructure, Mechanical Properties, and Strengthening Mechanisms of In Situ-Synthesized Ceramic-Reinforced Titanium Matrix Composite Coatings via Laser Cladding

Abstract

The laser cladding (LC) of titanium matrix composite coatings (TMCCs) on titanium components not only effectively enhances the wear resistance, fatigue resistance, corrosion resistance, and biocompatibility of titanium and its alloys, but also circumvents the incompatibility and low bonding strength issues associated with other metallic composite coatings. While the incorporation of ceramic particles is a critical strategy for improving the coating performance, the limited interfacial bonding strength between ceramic particles and the matrix has historically constrained its advancement. To further elevate its performance and meet the demands of components operating in harsh environments, researchers worldwide have employed LC to synthesize in situ hard ceramic reinforcements such as TiC, TiB, TiN, and others within TMCCs on titanium substrates. This approach successfully addresses the aforementioned challenges, achieving coatings that combine a high interfacial bonding strength with superior mechanical properties. This paper provides a comprehensive review of the processing techniques, phase composition, microstructure, and mechanical properties of in situ-synthesized ceramic-reinforced TMCCs via LC on titanium components, with a focused summary of their strengthening mechanisms. Furthermore, it critically discusses the challenges and future prospects for advancing this technology.

1. Introduction

Titanium and its alloys have been widely recognized for their low density, high specific strength, excellent corrosion resistance, and outstanding biocompatibility, leading to extensive applications across biomedical, aerospace, automotive, marine engineering, and petrochemical sectors [1,2]. Nevertheless, these materials exhibit two fundamental limitations: (1) their low hardness results in inadequate wear resistance and fatigue resistance, frequently causing component damage and failure under prolonged cyclic loading conditions; (2) their corrosion resistance becomes significantly compromised when deployed in complex service environments such as physiological fluids, atmospheric exposure, marine settings, and chemical media [3,4,5,6]. To overcome these constraints, surface modification technologies have emerged as a strategically advantageous alternative to bulk modification approaches, offering enhanced efficiency and cost-effectiveness for localized component reinforcement [7].

Laser cladding (LC) has emerged as a novel and highly efficient surface modification technology [2,8,9]. In recent years, its exceptional forming quality and precision have established it as a critical technique for the surface modification of titanium-based materials. Typical cladding materials encompass ceramic (ranging from single-phase to multi-phase systems), heterogeneous alloys (enhanced or not), and homogeneous alloys (enhanced or not). When applying LC for surface modification, two pivotal factors demand meticulous attention: the interfacial bonding strength between the substrate and cladding material and the thermal expansion compatibility of these components [10,11]. The bonding strength fundamentally determines the coating’s resistance to delamination under service conditions, where strong adhesion is imperative for achieving reliable modification outcomes. Concurrently, mismatched thermal expansion coefficients between the coating and substrate generate residual stresses. Excessive stress accumulation typically manifests as microcrack nucleation, followed by progressive crack propagation, and eventual coating failure. Furthermore, beyond interfacial compatibility, the intrinsic mechanical properties of the coating itself critically govern the operational lifespan of the modified components. This necessitates the strategic selection of cladding materials exhibiting a superior hardness, wear resistance, and fatigue performance compared to the substrate.

For the LC surface modification of titanium-based components, enhanced homogeneous titanium matrix composites (TMCs) have been identified as the most promising cladding materials [5,6,8,9,11,12]. This is to minimize the thermal expansion coefficient difference between the substrate and the cladding layer, achieve strong metallurgical bonding at the interface, and ensure the mechanical performance of the coating. The reinforcement phases in titanium matrix composites coatings (TMCCs) are predominantly ceramic, with TMCCs classified into ex situ-synthesized and in situ-synthesized categories based on ceramic formation mechanisms [13]. Compared to their ex situ counterparts, in situ-synthesized ceramic-reinforced TMCCs capitalize on titanium’s strong chemical affinity to generate ceramic phases through controlled interfacial reactions during cladding. This process not only strengthens the interfacial bonding between ceramic phases and the titanium matrix, but also promotes a uniform particle distribution. Such effects significantly amplify the reinforcement efficacy of ceramic particles, ultimately improving the critical mechanical performance metrics including the hardness, wear resistance, and fatigue strength [13,14].

Current research predominantly emphasizes the application of heterogeneous alloys or foreign materials for cladding titanium-based components, largely neglecting the inherent compatibility advantages of homogeneous alloy systems. This oversight underscores the critical need for in-depth investigations into the phase, microstructure, mechanical properties, and strengthening mechanisms of in situ-synthesized ceramic-reinforced TMCCs via LC. Such a fundamental understanding is poised to provide groundbreaking insights for the rational design and development of advanced coatings.

This paper systematically reviews the processing methodologies, phase, microstructure, and mechanical properties of in situ-synthesized ceramic-reinforced TMCCs via LC, with particular emphasis on elucidating the underlying strengthening mechanisms. The discussion culminates in a critical analysis of prevailing challenges and future research directions in this field.

2. Titanium Matrix Composite Coatings

Titanium matrix composite coatings (TMCCs) refer to composite coatings formed by incorporating or in situ synthesizing one or more reinforcement phases into titanium or its alloys [15,16]. Ceramic are the most prevalent reinforcing compounds for these coatings, including TiC, TiB, TiN, Ti₅Si₃, TiB₂, B₄C, Si₃N₄, SiC, and so on. Additionally, carbon-based materials (e.g., carbon nanoparticles, carbon nanotubes, and graphene) and bioactive ceramics such as hydroxyapatite (HA) and calcium phosphate (CaP) have been employed as effective reinforcements [2,5,6,8,12]. For the forming methods of TMCCs, they include physical vapor deposition, chemical vapor deposition, thermal spraying, LC, and so on [17].

For TMCCs with different ceramics, the main phase composition of the coatings, in addition to in situ- or ex situ-synthesized ceramic phases, primarily consists of various allotropic phases of titanium [18,19]. The main forms of titanium include the α phase, β phase, and metastable phases retained during the transformation between the α and β phases, including the secondary α phase, β′ phase, α′ phase, α″ phase, α₂ phase, and ω phase. The mixing and regular arrangement of these phases form the microstructure of TMCCs. The process parameters and reinforcement phases influence the phase transformation during the melting and solidification process, thereby affecting the microstructure of the material and ultimately determining the fundamental mechanical properties of the material [20,21].

The phase composition of TMCCs can be precisely determined through phase analysis, while the microstructure is imaged using optical microscopy or scanning electron microscopy, reflecting the surface characteristics of the sample. Since it is an image and does not involve the structure and composition, the description of the microstructure often carries a certain degree of subjectivity. Therefore, researchers may have some differences in describing specific microstructures. In general, the description of the microstructure of TMCCs is usually based on the morphology and content of the α phase. Common descriptions include an equiaxed structure, bimodal structure, basketweave structure, lamellar structure, Widmanstätten structure, trimodal structure, and mixed structure [22,23,24,25].

3. Laser Cladding Technology

Laser cladding (LC) is an advanced surface modification technique that utilizes high-energy-density laser beams to simultaneously melt both the cladding material and substrate surface, forming a metallurgical bonded coating. Based on the material feeding mechanisms, this process can be categorized into coaxial powder and preplaced powder feeding methods [2,8,10,12], as schematically illustrated in Figure 1. In recent years, LC has rapidly evolved into a cutting-edge technology within the surface engineering domain. The near-adiabatic and rapid heating/cooling characteristics of LC enable the formation of coatings with refined grain structures, a high density, and a uniform composition, allowing precise control over the product performance [26]. To achieve high-performance laser-cladded coatings, three critical factors must be systematically controlled.

(1)

Laser source and substrate–cladding material pairs

The first considerations for LC involve laser source selection and the fundamental characteristics of substrate–cladding material pairs. The laser wavelength [27] and pulse frequency [28] critically influence energy absorption and reflection during processing, directly determining the coating’s microstructure and properties. These parameters require optimization based on material-specific attributes. The foremost material selection criterion is achieving thermal expansion coefficient compatibility between the substrate and cladding material to ensure interfacial bonding strength [29]. For LC of the titanium-based substrate, common cladding materials include ceramic systems (e.g., TiC, TiB, and HA), heterogeneous alloys (e.g., nickel-based [30,31], iron-based [32,33], cobalt-based [34], and magnesium-based [35]), and TMCs (e.g., Ti-TiC and Ti-TiC-TiB). Notably, high-entropy alloys have recently gained research attention in surface modification [36].

(2)

Process parameters

Under established conditions of the laser source and material characteristics, the selection of the process parameters plays a decisive role in determining the coating performance. The primary process parameters affecting the LC of metal matrix composites include the laser power, scanning speed, overlap rate, scanning strategy, and so on. These parameters collectively determine a critical aspect of LC: the energy density. Composite parameters such as the volumetric energy density, linear energy density, and mass energy density are widely used to optimize LC processes [37,38]. Taking the volumetric energy density as an example, it represents the energy absorbed per unit volume of powder material. Equation (1) provides the expression for the laser volumetric energy density [37,38]:

$$E_V={\displaystyle \frac{P}{vhd}}$$

(1)

where $P$ is the laser power (W), $v$ is the scanning speed (mm/s), $h$ is the cladding track width (mm), and $d$ is the layer thickness (mm). The volumetric energy density governs the energy absorption of the molten pool, thereby critically influencing its fluidity, which is characterized by viscosity. The selection of the viscosity involves a trade-off. On the one hand, the viscosity should be sufficiently low to ensure adequate fluidity for effective spreading onto the substrate or previously deposited material. On the other hand, the viscosity must be sufficiently high to prevent defects such as porosity. As shown in Equation (1), increasing the laser power raises the volumetric energy density, while increasing the scanning speed reduces it. Therefore, matching the process parameters is necessary to achieve an appropriate volumetric energy density, thereby obtaining the desired molten pool viscosity and ensuring both the cladding efficiency and high coating density [39,40].

(3)

The content of reinforcement phases

For reinforced alloy coatings with predetermined reinforcement phases, the content of these phases constitutes a critical factor influencing the coating performance [41,42]. Macroscopically, the content of hard reinforcement phases affects the molten pool viscosity, thereby governing the formation of defects such as porosity. Microscopically, it determines the dominant strengthening mechanisms and microstructural characteristics. Specifically, the contents of the reinforcement phases regulate the volume fraction of the secondary-phase particles, which governs the dispersion strengthening effects. Additionally, these particles influence the phase transformation kinetics and grain nucleation/growth processes, ultimately dictating the coating properties.

In recent years, diverse metal materials including Ni-based, Co-based, Fe-based, and Ti-based alloys, as well as high-entropy alloys, have been extensively employed as cladding materials in LC technology [36,43,44,45]. However, TMCs have not received commensurate research attention due to their inherent limitations in their wear resistance and processability. In LC applications where the substrate is also titanium-based, the inherent compatibility advantages of TMCCs fail to be prominently realized. Recent advancements in TMCs have provided critical insights for material selection in the LC of titanium substrates. TMCCs fabricated via laser cladding on titanium substrates have garnered increasing attention and research, particularly those with in situ-synthesized ceramic reinforcements, which exhibit a superior metallurgical bonding performance.

4. In Situ-Synthesized Ceramic-Reinforced TMCCs

As shown in Figure 2, in situ-synthesized ceramic-reinforced TMCCs can be classified into three categories based on the ceramic composition: single-phase ceramic coatings (e.g., TiC-, TiB-, and TiN-reinforced TMCCs), dual-phase ceramic coatings (e.g., TiC/TiB-, TiB/TiN-, Ti₅Si₃/TiN-, and Ti₅Si₃/TiC-reinforced TMCCs), and multi-phase ceramic coatings (e.g., TiC/TiB/TiN- and TiC/TiN/Ti₅Si₃-reinforced TMCCs). Current research on applying LC to fabricate ceramic-reinforced TMCCs on titanium substrates remains limited. In contrast, studies utilizing laser additive manufacturing technologies (including selective laser melting (SLM), directed energy deposition (DED), laser-based powder fusion (LBPF), and so on) to produce high-performance ceramic-reinforced TMCs have achieved substantial depth. Given the fundamental similarities in processing mechanisms and physical principles between these two technologies, incorporating representative additive-manufactured TMCs into the current discussion provides critical theoretical support for comprehensively understanding the coating performance characteristics and strengthening mechanisms.

4.1. Universal Strengthening Mechanisms

For ceramic-reinforced TMCCs/TMCs, certain strengthening mechanisms exhibit universal applicability due to the inherent high strength and uniform distribution of the reinforcement phases. The following sections first elucidate universal strengthening mechanisms, with distinctive mechanisms discussed subsequently.

4.1.1. Solid Solution Strengthening

Solid solution strengthening [46,47,48,49,50,51,52,53,54,55] arises from solute atoms dissolving in the metal matrix to form interstitial solid solutions or substitutional solid solutions. The solute atoms in these configurations interact with dislocations through elastic, chemical, and geometric interactions, with elastic interactions being the most significant. When dislocations exist in a crystal, severe lattice distortion occurs around the dislocation line, causing atoms in the distorted region to deviate from their equilibrium positions and generate elastic strain and stress fields. When solute atoms segregate around dislocation lines, various atmospheres form. Solute atoms clustering around edge dislocations form Cottrell atmospheres [56,57], while interstitial atoms (e.g., C, N) interacting with screw dislocations form Snoek atmospheres [58]. Additionally, chemical interactions between solute atoms and dislocations lead to the formation of Suzuki atmospheres [59]. These atmospheres pin or drag dislocations, hindering their movement and thereby enhancing the material strength and wear resistance, though often at the expense of ductility and toughness. When solute atoms do not segregate around dislocations, but instead distribute randomly in the matrix as individual atoms or clusters, they interact geometrically with dislocations, leading to homogeneous solid solution strengthening, which is related to the formation of anti-phase boundaries within the crystal.

4.1.2. Second-Phase Strengthening

Second-phase strengthening can be divided into precipitation strengthening [60,61,62,63,64,65] and dispersion strengthening [66,67,68,69,70]. Precipitation strengthening occurs when solute atoms precipitate from a supersaturated solid solution to form second-phase particles, which maintain coherent or semi-coherent interfaces with the matrix. The strengthening mechanism involves dislocation cut-through particles. Dispersion strengthening typically refers to externally introduced hard, refractory second-phase particles distributed within the matrix, forming incoherent interfaces. Dispersion strengthening operates through Orowan strengthening. During Orowan strengthening, dislocations cannot directly cut through second-phase particles, but instead bypass them under external stress, leaving dislocation loops around the particles. This bending increases the lattice distortion energy in the affected region, raising resistance to dislocation motion and enhancing the slip resistance.

4.1.3. Grain Refinement Strengthening

Grain refinement strengthening [71,72,73] improves metal strength by reducing the grain size. In polycrystalline metals, grain boundaries are typically high-angle boundaries. When adjacent grains with different orientations undergo plastic deformation, dislocation sources in grains with larger Schmid factors activate first, leading to slip and multiplication on specific crystallographic planes. Dislocations moving toward grain boundaries are blocked, preventing the direct propagation of plastic deformation into neighboring grains and causing a dislocation pile-up within the deforming grains. Under external stress, the stress field generated by a dislocation pile-up at grain boundaries can activate dislocation sources in adjacent grains. Notably, smaller grain sizes often enhance material plasticity due to the increased participation of grain boundaries in slip processes.

4.1.4. Load Transfer Strengthening

Load transfer strengthening [70,74,75,76,77,78,79] applies to materials containing hard particles, primarily enhancing plasticity. When the material is under a tensile load, shear stress is transferred from the matrix to the reinforcement through the interface. The shear stress parallel to the loading direction at the interface balances the normal stress perpendicular to the interface. Since reinforcements typically have a higher strength than the matrix, this results in strengthening.

For the above strengthening mechanisms, their relative contributions cannot be precisely quantified. However, some researchers have proposed strengthening equations to estimate the effects of specific mechanisms. These equations identify key variables influencing strengthening, allowing researchers to selectively control these variables to tailor material properties.

Table 1. Strengthening equations of different strengthening mechanisms.

Strengthening Mechanism		Formation Reasons	Equation		Reference
Solid Solution Strengthening		Interaction of solute atoms with dislocations.	$\mathsf{\Delta }\sigma ={\displaystyle \frac{1}{S_F}}{\left({\displaystyle \frac{F_m^4{c}^2\omega }{4G{b}^9}}\right)}^{{\displaystyle \frac{1}{3}}}$		[46,55]
			$\mathsf{\Delta }\sigma =A{C}_0^{{\displaystyle \frac{2}{3}}}$		[47,48,49,50]
			$\mathsf{\Delta }\sigma =M{\displaystyle \frac{3^{\frac{3}{4}}}{2}}{\left({\displaystyle \frac{1-\nu }{1+\nu }}\right)}^{{\displaystyle \frac{3}{2}}}G{\left|\epsilon \right|}^{{\displaystyle \frac{3}{2}}}{c}^{{\displaystyle \frac{1}{2}}}$		[51,52,53,54]
Second-phase strengthening	Precipitation Strengthening	Dislocations cut through second-phase particles precipitated from supersaturated solid solutions.	$\mathsf{\Delta }\sigma =\alpha MG{\epsilon }^{{\displaystyle \frac{3}{2}}}{\left({\displaystyle \frac{fr}{b}}\right)}^{{\displaystyle \frac{1}{2}}}$		[60,61,62,63]
			$\mathsf{\Delta }\sigma =\beta M{\displaystyle \frac{{\gamma }_{apb}}{2b}}{\left({\displaystyle \frac{3\pi f}{8}}\right)}^{{\displaystyle \frac{1}{2}}}$		[63,64,65]
			$\mathsf{\Delta }\sigma =\eta M{\left(\Delta G\right)}^{{\displaystyle \frac{3}{2}}}{\left({\displaystyle \frac{2f}{G}}\right)}^{{\displaystyle \frac{1}{2}}}{\left({\displaystyle \frac{r}{b}}\right)}^{{\displaystyle \frac{3}{2}}m-1}$		[63,64]
	Dispersion Strengthening	Dislocations bypass dispersed second-phase particles.	$\mathsf{\Delta }\sigma =\alpha {\displaystyle \frac{Gb}{\lambda }}\mathit{ln}{\displaystyle \frac{r}{b}}$		[66,67,68,69,70]
Grain Refinement Strengthening		Grain refinement and increased number of grain boundaries.	$\mathsf{\Delta }\sigma =K\left({\displaystyle \frac{1}{\sqrt{d_1}}}-{\displaystyle \frac{1}{\sqrt{d_0}}}\right)$		[71,72,73]
Load Transfer Strengthening		Hard reinforced phase particles for effective load sharing.	$\mathsf{\Delta }\sigma =V{\sigma }_m{\displaystyle \frac{\left(1+t\right)A}{4l}}$	$\mathsf{\Delta }{\sigma }_{particle}=kV{\sigma }_m$	[74,75,76]
				$\mathsf{\Delta }{\sigma }_{fibre}=kV{\sigma }_m{\displaystyle \frac{l}{d}}{c}_0$	[70,77,78,79]
Dislocation Strengthening		Mismatch in thermal expansion coefficients and elastic modulus.	$\mathsf{\Delta }\sigma =\alpha MGb{\rho }^{{\displaystyle \frac{1}{2}}}$		[70,80,81,82]

summarizes some of these strengthening equations.

4.2. TiC-Reinforced TMCCs

TiC, with its thermal expansion coefficient close to that of titanium and excellent compatibility, is one of the most widely used and intensively studied ceramic reinforcements [83,84,85]. With advancements in in situ synthesis techniques, various carbon sources (e.g., CNTs, carbon nanotubes, graphene, carbon powder, and TiC powder) have been utilized to generate in situ TiC [86,87]. Notably, even in ex situ TiC-reinforced coatings, in situ TiC is frequently detected and plays a significant role in enhancing the material performance [88,89,90,91,92].

Table 2. The process parameters, phase compositions, mechanical properties, and strengthening mechanisms of in situ-synthesized TiC-reinforced TMCCs/TMCs.

Method	Material	Laser Power (W)	Scanning Speed (mm/s)	Phase	Hardness	Loss/ Wear Rate	Coefficient of Friction	Yield Strength (MPa)	Ultimate Tensile Strength (MPa)	Maximum Strain or Elongation (%)	Strengthening Mechanism	Ref.
LC	CP-Ti + 5 wt% CNT	700	5	(α-Ti) + TiC	382 HV0.5	0.50 mm3	0.406				GRS; DPS; LTS	[93,94]
	CP-Ti + 10 wt% CNT				382 HV0.5	0.33 mm3	0.313
	CP-Ti + 15 wt% CNT				800 HV0.5	0.25 mm3	0.398
	CP-Ti + 20 wt% CNT				1125 HV0.5	0.14 mm3	0.263
LC	TiC + 5 wt% CNT	2000	5–5.83	Ti + TiC + VC + Al3Ti + Ti3AlC2	500–2700 HV0.5		0.467				DDS	[95]
DED	TC4 + 10 wt% TiCp	425	5	(α-Ti) + (β-Ti) + TiC	420.8 HV0.2	0.00786 g	0.375		816.8	3.32	GRS; SSS; LTS; CTE	[96]
	TC4 + 20 wt% TiCp				454.6 HV0.2	0.00759 g	0.367		783.5	3.06
	TC4 + 30 wt% TiCp	450			490.5 HV0.2	0.00689 g	0.483		713.5	2.58
	TC4 + 40 wt% TiCp				552.8 HV0.2	0.00600 g	0.551		563.1	2.05
	TC4 + 50 wt% TiCp				730.2 HV0.2	0.00461 g	0.608		515.5	1.83
DED	CP-Ti + 0.26 wt% C	200	3	(α-Ti) + TiC	222 HV		0.45	420	480	30	GRS; SSS; LTS; OWS; CTE	[97]
	CP-Ti + 0.43 wt% C				263 HV		0.4	490	510	32
	CP-Ti + 0.60 wt% C				273 HV		0.26	840	940	19
	CP-Ti + 1.20 wt% C				285 HV		0.36	610	710	24
	CP-Ti + 1.60 wt% C				312 HV		0.44	650	760	16
DED	CP-Ti + 1 wt% TiC	700	10	(α-Ti) + (β-Ti) + TiC	230 HV0.2	1.70 mm3/Nm	0.51		560	1.31	GRS; SSS; LTS	[98]
	CP-Ti + 2 wt% TiC				255 HV0.2	1.18 mm3/Nm	0.50		590	1.18
	CP-Ti + 3 wt% TiC				265 HV0.2	2.45 mm3/Nm	0.51		616	0.79
	CP-Ti + 5 wt% TiC				250 HV0.2	3.43 mm3/Nm	0.64		725	0.62
	CP-Ti + 10 wt% TiC				270 HV0.2	5.67 mm3/Nm	0.62		365	0.42
	CP-Ti + 15 wt% TiC				300 HV0.2	3.02 mm3/Nm	0.63		355	0.25
DED	TC4 + 2.93 wt% TiC	1800	6	(α-Ti) + (β-Ti) + TiC				954	1050	5.76	GRS; OWS	[68]
SLM	TC11 + 0.2 wt% GNP	280	1200	(α′-Ti) + (β-Ti) + GNP + TiC				1110	1384	8.1	DLS(CET); LTS; GRS; OWS	[99]
SLM	TC4 + 5.0 vol% TiC	220	600	(α-Ti) + (α′-Ti) + (β-Ti) + TiC				1257.40	1365.83	1.27	LTS; GRS	[100]
			900					1284.29	1424.16	1.47
			1200					1364.42	1538.98	2.92
LPBF	TC4 + 9% volCH4	180	200	(α-Ti) + (α′-Ti) + (β-Ti) + TiC + TiAl	415.76 HV	2825.43 μm2	7.93	1176	1884	23.28	GRS; DPS, PPS, LTS	[101]
	TC4 + 19% volCH4		200		437.09 HV	2463.51 μm2	6.34	1251	1919	21.69
	TC4 + 29% volCH4		200		431.41 HV	2824.08 μm2	8.16	1245	1718	7.71
	TC4 + 19% volCH4		125		415.98 HV	2928.55 μm2	7.82	1227	1709	10.3
			150		421.26 HV	2728.36 μm2	5.27	1236	1834	14.86
			300		432.43 HV	2690.47 μm2	7.82	1221	1843	20.63
		135	150		421.35 HV	2875.79 μm2	6.14	1237	1847	17.34
		99	110		415.07 HV	3030.50 μm2	8.15	1181	1649	10.78

Note: The data in the tables in this paper are approximate unless exact values are given. The titanium alloys in the tables in this paper are uniformly designated using Chinese national standard GB/T 3620.1-2016 [102] (e.g., TC4 corresponds to Ti6Al4V). Strengthening mechanisms’ abbreviations: SSS: solid solution strengthening; PPS: precipitation strengthening; DPS: dispersion strengthening; GRS: grain refinement strengthening; LTS: load transfer strengthening; DDS: dendritic strengthening; OWS: Orowan strengthening; DLS: dislocation strengthening; CTE: coefficient of thermal expansion mismatch strengthening.

summarizes the process parameters, phase compositions, mechanical properties, and strengthening mechanisms of representative in situ-synthesized TiC-reinforced TMCCs/TMCs.

The phase composition of laser-cladded TiC-reinforced TMCCs is influenced by multiple factors beyond the reinforcement phase and processing parameters. In addition, it is also related to the type of the component’s material and the titanium-based material. However, in practical production, pure α- or β-phase titanium alloys are rarely obtained, with two-phase coexistence being the norm [97,98,103]. This principle applies equally to other ceramic-reinforced TMCCs. While the interfacial bonding strength remains a critical consideration, this study focuses on TMCCs for titanium-based components. Therefore, when analyzing the phase composition and microstructure, particular attention is given to homogeneous non-fusion regions. To gain an in-depth understanding of the Ti-C system, Figure 3 presents the Ti-C-phase diagram (phase diagrams of other systems are also provided accordingly).

4.2.1. Phase and Microstructure

The main phase composition of in situ-synthesized TiC-reinforced TMCCs via LC includes α-Ti, β-Ti, and TiC (Figure 4a). Notably, TiC is a non-stoichiometric compound. Figure 5 is a schematic diagram of the phase transformation based on the Ti-C-phase diagram (Figure 3a). When the carbon concentration is in the hypoeutectic region (Line a1), β-Ti solidifies first from the titanium melt. As the temperature decreases to the eutectic transformation temperature (in practical solidification processes, these temperatures are influenced by alloying elements, and so are the temperatures discussed elsewhere in the text), the titanium melt undergoes a eutectic reaction, forming eutectic TiC particles and eutectic β-Ti. When the carbon concentration is in the hypereutectic region (Line a2), due to the higher carbon concentration, TiC precipitates as the primary phase and acts as a heterogeneous nucleation site for β-Ti, promoting its formation. When the temperature drops below the β-Ti transformation temperature, β-Ti transforms into α-Ti, with α-Ti dendrites distributed within β-Ti grains, forming a Widmanstätten structure. The incomplete transformation of β-Ti to α-Ti is primarily related to two factors [98]. First, the rapid cooling characteristics of laser technology leave insufficient time for a complete phase transformation [104,105]. Second, many titanium alloys contain β-stabilizing elements such as Mo and V, which expand the β-Ti-phase region. These explanations also generalize to other TMCCs with different reinforcements [106].

Figure 3. Phase diagrams of (a) Ti-C, (b) Ti-B, (c) Ti-N and (d) Ti-Si, redrawn from Refs. [107], [108], [109], and [110], respectively. Minor adjustments to scales and font styles have been made for consistency.

Figure 3. Phase diagrams of (a) Ti-C, (b) Ti-B, (c) Ti-N and (d) Ti-Si, redrawn from Refs. [107], [108], [109], and [110], respectively. Minor adjustments to scales and font styles have been made for consistency.

Figure 4. (a) XRD patterns of TiC-reinforced TMCs [96]. (b) SEM imaging of TMCs’ microstructures: (b1) 20% TiC_p/Ti6Al4V, (b2) 50% TiC_p/Ti6Al4V (DPT: dendritic primary TiC; GPT: granular primary TiC; UMT: unmelted TiC) [96]. (c) Densification map (black cycle: ≤50 μm; blue cycle: 50–100 μm; yellow cycle: ≥100 μm) [100]. (d) Microstructures of different laser powers TiC-TMCs: (d1) 800 W, (d2) 1200 W [104]. (e) The relationship between the microhardness and TiC_p content [96]. (f) The relationship between the mass loss or wear rate and TiC content [98]. (g) Wear mechanism diagram of TiC/TC17 composites [111].

Figure 4. (a) XRD patterns of TiC-reinforced TMCs [96]. (b) SEM imaging of TMCs’ microstructures: (b1) 20% TiC_p/Ti6Al4V, (b2) 50% TiC_p/Ti6Al4V (DPT: dendritic primary TiC; GPT: granular primary TiC; UMT: unmelted TiC) [96]. (c) Densification map (black cycle: ≤50 μm; blue cycle: 50–100 μm; yellow cycle: ≥100 μm) [100]. (d) Microstructures of different laser powers TiC-TMCs: (d1) 800 W, (d2) 1200 W [104]. (e) The relationship between the microhardness and TiC_p content [96]. (f) The relationship between the mass loss or wear rate and TiC content [98]. (g) Wear mechanism diagram of TiC/TC17 composites [111].

Figure 5. Schematic diagram of the phase transition of TiC-reinforced TMCCs.

Figure 5. Schematic diagram of the phase transition of TiC-reinforced TMCCs.

The phase transformation process of the coating is related to the carbon source content. As the carbon source content increases, the size of TiC dendrites and the titanium matrix increases due to the enhanced carbon activity in the molten pool. Secondary or multiple dendrites may also form, while the content of unmelted carbon sources (such as TiC and CNTs) in the molten pool increases (Figure 4b) [68,93,96,97,98]. The growth of the TiC dendrite size is also related to the energy absorption of the system. On the one hand, TiC has a higher laser absorption coefficient than the matrix, and its increased content raises the thermal input of the system [96,97,101]. On the other hand, the in situ reaction is exothermic, which further promotes the temperature rise in the molten pool. This provides a longer growth time for TiC dendrites and the titanium matrix [13,14,97].

The influence of process parameters on the phase transformation of the coating is evaluated through their impact on the thermal input, typically converted into the laser energy density for assessment. Compared to an excessively low laser energy density, a moderate laser energy density results in a higher molten pool temperature, promoting the complete reaction between Ti and C. This leads to a more accurate stoichiometric ratio of TiC, smaller XRD peak shifts, finer grains, smaller dendrite spacing, and a more uniform distribution of TiC grains in the α-Ti matrix (Figure 4d) [37,38,100,104,106]. This is because a high energy density not only enhances the fluidity and homogeneity of the molten pool, reducing unmelted particles and defects, but also induces rapid cooling that suppresses dendrite coarsening and refines the microstructure. However, it should be noted that an excessively high energy density is also undesirable (Figure 4c). First, when the energy density is too high, it can exceed the stability threshold of the molten pool, forming an unstable keyhole. During rapid solidification, the collapse of the keyhole may trap gas inside the material, forming pores. Second, an excessively high energy density can cause localized overheating, leading to the evaporation of the titanium matrix or carbon, resulting in compositional deviations [112,113,114]. Third, an excessively high energy density causes the rapid cooling of the molten pool, creating significant thermal gradients and residual stresses. Residual stresses may induce microcracks or macroscopic deformation, and an uneven stress distribution can also enhance the directional dependence of the mechanical properties. Fourth, an excessively high energy density may also result in poorer surface accuracy and uneconomical energy consumption.

4.2.2. Mechanical Properties and Strengthening Mechanisms

From Table 2, it can be seen that the TiC-reinforced TMCCs obtained by LC exhibit high hardness, excellent wear resistance, and good plasticity and toughness. However, the advantages of the process parameters, TiC content, and other factors vary under different conditions or for different properties, leading to nonlinear and even non-monotonic changes in performance over a wide range of contents or parameters (Figure 4e,f). This is closely related to the relative contributions of various strengthening mechanisms.

As the TiC content rises, the hardness continuously increases. This is primarily because the added TiC ceramic phase is a hard phase with a significantly higher hardness than that of titanium. The TiC particles impede the movement of dislocations, and this obstruction stems from multiple aspects.

(1) The in situ-synthesized TiC, as second-phase particles, is dispersed in the matrix, giving rise to dispersion strengthening [68,93,94,97,99]. When dislocations move, they need to bypass the second-phase particles, and the distortion of the dislocation lines increases the resistance to dislocation movement. When the content of the in situ-synthesized TiC is sufficiently high, the TiC may also precipitate as a supersaturated phase, thereby achieving precipitation strengthening [101]. When dislocations move, they need to cut through the TiC particles, thus exerting a strengthening effect.

(2) Due to various interactions between solute atoms like C and dislocations, different types of atmospheres are formed, which pin or drag the dislocations [96,97,98]. However, due to the small solubility of C in titanium, the effect of solid solution strengthening is weak.

(3) The grain refinement strengthening provides a large number of dense grain boundaries, and the hindering effect of the grain boundaries on dislocations is enhanced, contributing to the increase in hardness. There are multiple reasons for the grain refinement. From the perspective of the processing technology, the high energy input and high-frequency thermal cycles in laser processing can induce dynamic recrystallization, functioning as a heat treatment process and refining the grains [68,93,95,97,98]. From the perspective of the nucleation and growth process of grains, the uniformly distributed synthesized TiC can serve as the nucleus for heterogeneous nucleation, remarkably increasing the nucleation rate [96,97,104,106,111]. Additionally, the added solute elements can, to a certain extent, expand the solidification temperature range of the titanium alloy, thus creating a larger composition supercooling at the solidification front. This supercooling condition promotes the formation of a large number of nuclei rather than the epitaxial growth of a single grain, ultimately resulting in the formation of equiaxed grains [96,97,100,115].

The yield strength, tensile strength, and elongation of a material are crucial parameters for characterizing the ability of a material to resist plastic deformation. Numerous studies have shown that when the TiC content is low, both the yield strength and tensile strength of the material can be enhanced to a certain extent. The improvement in the material’s ability to withstand tensile loads is not only related to the hindering effects of the aforementioned strengthening mechanisms on dislocations. More notably, the following two strengthening mechanisms deserve attention.

(1) Grain refinement strengthening. The finer the grains are, the more grain boundaries there will be. Grain boundaries can confine plastic deformation within a certain range, making the plastic deformation more uniform. Thus, refining the grains can improve the material’s plastic deformation ability [96,97]. Meanwhile, grain boundaries also act as obstacles to crack propagation. Therefore, grain refinement can improve the toughness of the material [97]. However, a large amount of experimental data indicate that the elongation after the fracture of the material decreases after adding the ceramic phase, suggesting that the toughness and ductility of the material decline. This is mainly related to several other types of strengthening brought about by the addition of the hard ceramic phase. The brittle and hard ceramic phase significantly increases the brittleness of the material, which cannot be compensated for by the limited degree of grain refinement [68,96,97,98,99,100].

(2) Load transfer strengthening. When the composite material is subjected to a tensile load, the load will be transferred from the matrix to the reinforcement with higher strength, thereby dispersing the stress and reducing the local stress concentration in the matrix [68,96,97,98,100,106,111].

With a further increase in the TiC content, the tensile properties of the material decrease, which is also related to the aforementioned strengthening mechanisms. Through the analysis of the phase composition and microstructure of the material in the previous text, it is not difficult to find that as the content of the reinforcement increases, the grain size first becomes finer and then coarser, and finally, the microstructure also changes from an equiaxed structure to a Widmanstätten structure. When the in situ reaction is carried out by adding TiC particles, due to the increase in the content, the number of TiC particles that cannot be completely melted in the final material increases, and the size of the in situ-synthesized TiC particles also increases further [96]. There is a certain degree of segregation of various TiC phases in the matrix. Even if the in situ TiC phase is obtained through other methods, when a higher content of the TiC phase is required, there will also be a certain degree of size increase and uneven distribution [97]. On the one hand, such a result weakens the strengthening effect of grain refinement, further reducing the tensile properties, plasticity, and toughness of the material. On the other hand, the strengthening effect of load transfer strengthening is largely restricted by the interfacial bonding strength between the reinforcement and the matrix. The interfacial bonding between a large number of ex situ TiC phases and the matrix is much worse than that between in situ TiC phases. In addition, the increase in the size of TiC particles will reduce the interfacial area, thereby decreasing the strengthening effect. At the same time, the segregation of the TiC phase is also unfavorable for the transfer of the load.

The wear resistance of a material is generally characterized by the wear amount in a wear test. Under the same conditions, the greater the mass or volume loss due to wear, the poorer the wear resistance [116]. Regarding the TiC-reinforced TMCCs/TMCs, there is no consistent opinion on the influence of the TiC phase content. Extensive research has shown that the changes in the wear loss and friction coefficient are not synchronized, indicating that the wear resistance of materials is influenced by the coupling of multiple factors. The friction coefficient, along with other factors, determines the wear resistance of materials, primarily in two aspects: the hardness of the material and the interfacial bonding strength between the reinforcement and the matrix.

When the friction coefficient of the material is high, if the material also exhibits high hardness, the material will also show excellent wear resistance due to its high compressive strength. When the friction coefficient is low, theoretically, it can reduce the wear rate of the material. However, if the material has insufficient hardness or poor toughness, the material may still fail rapidly due to plastic deformation or brittle spalling. In more studies, the improvement of the friction properties of the material often occurs when the content of the reinforcement is low [97,98,100]. This is because TiC is uniformly distributed in the form of fine eutectic phases, which reduces the wear amount through grain refinement strengthening, load transfer strengthening, and other hardness strengthening mechanisms. At the same time, the size of the detached TiC particles is small, forming a lubricating layer, and the friction coefficient of the material decreases. When the content of the reinforcement is high, the unmelted, coarse TiC dendrites cause three-body wear. At the same time, the brittle phase leads to the loss of plasticity of the matrix, and the wear loss increases again, while the rough surface makes the friction coefficient rise (Figure 4g) [101,111,115]. It is worth considering that as the content of the reinforcement increases, the hardness of the material continuously increases, but the wear resistance of the material fails to continuously improve. This is because more unmelted initial TiC phases or large-sized in situ TiC phases are synthesized. The interfacial bonding strength between the unmelted phases and the matrix is poor, and they are likely to spall off as brittle phases, thus deteriorating the friction environment and increasing the wear. At the same time, the resulting composition segregation and the increase in grain size will also reduce the strengthening effects of several strengthening mechanisms.

4.3. TiB-Reinforced TMCCs

In the design of TMCCs, TiB exhibits remarkable advantages due to its unique chemical stability and interfacial properties [117,118,119]. Compared with TiC, the difference in the coefficient of thermal expansion between TiB and the titanium matrix is smaller [120]. By combining its high elastic modulus and excellent interfacial compatibility, it can not only effectively improve the strength of the material, but also perform outstandingly in enhancing the corrosion resistance and high-temperature properties of the material [121,122]. Commonly used materials for in situ synthesizing TiB-reinforced TMCCs are TiB, TiB₂, boron powder, etc.

Table 3. The process parameters, phase compositions, mechanical properties, and strengthening mechanisms of in situ-synthesized TiB-reinforced TMCCs/TMCs.

Method	Material	Laser Power (W)	Scanning Speed (mm/s)	Phase	Hardness	Loss/ Wear Rate	Coefficient of Friction	Yield Strength (MPa)	Ultimate Tensile Strength (MPa)	Maximum Strain or Elongation (%)	Strengthening Mechanism	Ref.
LC	TC4 + 5 wt% TiB2	2400	8–10	(α-Ti) + (β-Ti) + TiB2 + TiB	500.3 HV0.2	0.2949 g/min					LTS; GRS; PPS; SSS	[123]
	TC4 + 15 wt% TiB2				555.3 HV0.2	0.1146 g/min
	TC4 + 25 wt% TiB2				635.9 HV0.2	0.1078 g/min
	TC4 + 35 wt% TiB2				747.1 HV0.2	0.0216 g/min
LC	TiB2	2950	6	(α-Ti) + (β-Ti) + TiB2 + TiB	667 HV0.3						IFS	[124]
LC	CP-Ti + 30 vol% TiB2	2600	7	(α-Ti) + TiB2 + TiB	650 HV0.2	0.17 mm3	0.46				DDS; DPS; GRS	[125]
LB	Amorphous B	1000	5	(α-Ti) + (β-Ti) + TiB2 + TiB + TiB25 + TiO2	650 HV0.2						SPS; DLS; SS	[103]
		2500			1500 HV0.2		0.393
		3000			950 HV0.2	2.216 mm3/Nm	0.413
DLD	Ti6242 + 2 wt% TiB2	2500	5	(α-Ti) + (β-Ti) + TiB				1024.9	1100	6.2	LTS; GRS; IFS; SS	[72]
	Ti6242 + 5 wt% TiB2							1285.5	1376	3.7
DED	TC4 + 1.25 vol% TiB	1500	10	(α-Ti) + (β-Ti) + TiB				1010	1075	7.1	GRS; LTS; DLS; OWS; CTE; SS	[73]
	TC4 + 2.50 vol% TiB							1095	1153	6.2
	TC4 + 5.00 vol% TiB							1085	1174	3.5
SLM	TC4 + 0.01 wt% TiB2	160	800	(α-Ti) + (β-Ti) + TiB	351.8 HV						GRS; DPS; IFS; DLS; SS	[126]
	TC4 + 1.0 wt% TiB2				372.6 HV
	TC4 + 3.4 wt% TiB2				410.1 HV
	TC4 + 4.8 wt% TiB2				430.8 HV
SLM	TC4 + 5 wt% TiB2	350	972	(α-Ti) + (β-Ti) + TiB2 + TiB	580 HV	45.5 μm	0.44				GRS; SPS; SS	[127]
			530		573 HV	57.3 μm	0.51
			364		565 HV	78.5 μm	0.64

Note: IFS: interface strengthening; SPS: second-phase strengthening; SS: structure strengthening.

summarizes the process parameters, phase compositions, mechanical properties, and strengthening mechanisms of representative in situ-synthesized TiB-reinforced TMCCs/TMCs.

4.3.1. Phase and Microstructure

The main phase composition of in situ-synthesized TiB-reinforced TMCCs via LC includes α-Ti, β-Ti, and TiB (Figure 6a). Figure 7 illustrates the phase transformation process based on the Ti-B phase diagram (Figure 3b). In the hypoeutectic region (Line b1), β-Ti solidifies as the primary phase from the titanium melt with B atoms segregating at the β-Ti interface, followed by a eutectic reaction at lower temperatures that produces eutectic TiB whiskers (incongruent compound) and β-Ti. For the hypereutectic region (Line b2), the elevated boron concentration results in TiB whiskers forming as the primary phase, which simultaneously serve as heterogeneous nucleation sites for β-Ti formation. Upon further cooling below the β-Ti transformation temperature, dendritic or lamellar α-Ti precipitates from the β-Ti.

The in situ reaction for TiB formation is highly exothermic, which increases the temperature of the residual liquid phase between dendrites and exhibits a growth preference along the [010] direction, leading to the formation of a large number of needle or whisker TiB between dendrites (Figure 6e) [73]. As the β-Ti dendrites continue to grow, TiB is pushed and aggregated, forming a cellular network structure (Figure 6c,d). At the bottom of the molten pool, the high temperature gradient promotes the vertical alignment of TiB along the build direction, forming a columnar network. In the middle and upper parts of the molten pool, the reduced temperature gradient and enhanced compositional undercooling lead to the nucleation of equiaxed β grains, which encapsulate TiB, forming a continuous network skeleton [73,128]. The continuity of this structure improves with the increasing TiB content. In the columnar network region, TiB whiskers are distributed along the interior or grain boundaries of β grains and exhibit a strong <010> texture [72,73,129]. In the equiaxed region, TiB is randomly distributed at the grain boundaries.

With excessive increases in the boron source content, the increased B activity in the molten pool causes TiB whiskers to coarsen into needle or plate forms, and the size of the titanium matrix also increases (Figure 6b) [73,123]. The content and size of unmelted boron sources (such as TiB and TiB₂) in the molten pool increase. Additionally, the increase in the B content promotes the formation of the TiB₂ (congruent compound) phase [119,123,124,127].

When the laser energy density is low, the top of the molten pool consists of nearly equiaxed grains, while the bottom consists of columnar dendritic grains. Grain boundaries contain a large number of TiB whiskers, exhibiting significant agglomeration and forming a network structure (Figure 6c). When the laser energy density is moderately increased, a larger equiaxed grain zone can be obtained, with larger TiB whiskers arranged more uniformly along the grain boundaries, forming a network structure with larger unit sizes (Figure 6d) [103,127,128]. When the laser energy density is excessively increased, TiB transforms into needle or even plate forms, while the grain size of the titanium matrix increases, causing the network structure to break apart [128].

4.3.2. Mechanical Properties and Strengthening Mechanisms

As can be seen from Table 3, the TiB ceramic-reinforced TMCCs obtained through LC, similar to the TiC ceramic-reinforced TMCCs, exhibit improved hardness, strength, wear resistance, and so on. The strengthening mechanisms related to dislocations and grain refinement also bear a high degree of similarity to those of the TiC ceramic-reinforced TMCCs. However, considering the differences in the physical and chemical properties and morphologies of the two reinforcing materials, there are also certain differences in their strengthening effects and mechanisms.

(1)

Different physical properties

TiC belongs to the cubic crystal system, with extremely high hardness. However, it is highly brittle and there is a certain difference in the coefficient of thermal expansion between it and the titanium matrix, which is likely to cause interfacial thermal stress [130]. Nevertheless, the reinforced composite materials usually show very high hardness. In contrast, TiB belongs to the hexagonal crystal system, with slightly lower hardness, but better toughness. Its coefficient of thermal expansion is closer to that of the titanium matrix, resulting in smaller residual stress [131]. The smaller difference in the coefficient of thermal expansion is also beneficial for the material to exhibit better high-temperature mechanical properties.

(2)

Different morphology

The in situ-synthesized TiC mainly exists in the form of dispersed particles. Therefore, in most studies, Orowan strengthening (dispersion strengthening) is considered the dominant strengthening mechanism [68,97,98,99,105]. However, excessive addition is likely to trigger brittle cracks, leading to a sharp decline in ductility. In contrast, TiB is mainly distributed in the form of whiskers or short fibers, and the network structure it induces in the material effectively improves the material’s properties. Many studies have pointed out the existence of this strengthening effect [72,73,103,126,127,128,132]. Liu et al. [73] found that the additional structural strengthening plays the most significant role among various strengthening mechanisms. Figure 6f is the schematic diagram of the mechanism of structural strengthening. In a hard/soft two-phase heterogeneous structure, geometrically necessary dislocations (GNDs) slip from the Frank-Read source inside the soft-phase grains to the grain boundaries (hard-phase regions), forming a dislocation pile-up and generating long-range back stress, forcing the soft phase to undergo strain gradient hardening. The in situ-synthesized hard TiB phase significantly enhances the hindering effect of the grain boundaries on GNDs. Especially in the later stage of deformation, the dislocation pile-up density and strain gradient at the hard phase interface increase significantly, thereby forming stronger back stress within the soft phase (such as the titanium matrix) and achieving the strengthening of the soft phase. In fact, strain gradient strengthening caused by GNDs also exists in TMCCs/TMCs reinforced with other phases [99,133,134].

(3)

Different chemical affinity

The implications of the difference in the chemical affinity between the two are manifold. First, compared to C, B is a better β-phase stabilizer, which can further promote the transition from columnar grains to equiaxed grains, achieving more significant grain refinement [135,136,137]. Second, TiC tends to react with the titanium matrix, particularly with the aluminum present in common titanium alloys, to form brittle interfacial phases such as Ti₃AlC, which weakens the interfacial bonding strength and limits the load transfer efficiency [95,101]. In contrast, the interface between TiB and the matrix is clean with fewer reaction products, forming a stronger interfacial bond with α-Ti/β-Ti. Third, a B₂O₃ or composite oxide layer may form on the surface of TiB, which, in synergy with the TiO₂ from the matrix, can also slow down the high-temperature oxidation rate [138,139,140]. The initial formation of these oxide films can inhibit the invasion of oxygen, thereby reducing the further oxidation of TiB and enhancing the strengthening effect of TiB.

4.4. TiN-Reinforced TMCCs

TiN, as an ideal high-temperature reinforcement phase, exhibits an exceptional performance in wear-resistant coatings due to its high hardness, superior thermal stability, and low friction coefficient [141,142,143]. Compared to TiC, TiN demonstrates higher fracture toughness and better wettability with metal matrices, making it suitable for dynamic load-bearing wear applications [144,145]. However, challenges arise in laser-nitrided in situ TiN coatings, including insufficient nitrogen diffusion under rapid cooling rates, leading to pore formation and interfacial stresses from thermal expansion mismatch [146,147,148]. To address these limitations, researchers increasingly adopt multiphase synergistic reinforcement strategies combining TiN with other phases. Nevertheless, understanding single-phase TiN strengthening mechanisms remains critical. Table 4 summarizes the process parameters, phase compositions, mechanical properties, and strengthening mechanisms of representative in situ-synthesized TiN-reinforced TMCCs/TMCs.

4.4.1. Phase and Microstructure

The main phase composition of in situ-synthesized TiN-reinforced TMCCs via LC includes α-Ti, β-Ti, and TiN (Figure 8a). The localized high temperature induced by the laser promotes the decomposition of N₂ molecules into active nitrogen atoms (or ions), which enter the molten titanium matrix through diffusion and convection, forming a supersaturated solid solution [149,150,151]. Figure 9 illustrates the phase transformation process based on the Ti-N phase diagram (Figure 3c). When the atomic percentage of N is less than 15.2% (taking Line c1 and Line c2 as examples), α-Ti precipitates first upon cooling to the liquidus temperature. As the temperature further decreases to the peritectic temperature (2020 °C), a peritectic reaction occurs, forming β-Ti. If the reaction time is sufficient, β-Ti will transform into α-Ti upon continued cooling until all β-Ti in the system is consumed. Simultaneously, due to the decreasing temperature and reduced solid solubility of N, Ti₂N precipitates from α-Ti. When the N composition lies in the region defined by Line c3 or Line c4, the higher nitrogen content causes TiN (a congruent compound but non-stoichiometric) to precipitate preferentially upon cooling to the liquidus temperature. This occurs primarily at the gas–liquid interface or near the edge of the melt pool where the nitrogen concentration is high. As the temperature decreases to the peritectic temperature (2350 °C), a peritectic reaction occurs, forming α-Ti, with TiN acting as the heterogeneous nucleation core. As the peritectic reaction proceeds, the remaining liquid phase along Line c3 is completely consumed upon cooling to the solidus temperature, transforming entirely into α-Ti. Ti₂N subsequently precipitates with further cooling. For compositions along Line c4, the liquid phase is fully consumed during the peritectic reaction, entering the α-Ti and TiN two-phase region. Upon further cooling to 1050 °C, a peritectoid reaction occurs, precipitating Ti₂N. However, in TiN-reinforced TMCCs/TMCs, Ti₂N is seldom detected. This is primarily attributed to rapid laser cooling inhibiting sufficient time for Ti₂N precipitation. The presence of β-Ti, conversely, is associated with the numerous β-phase-stabilizing elements present in titanium alloys.

Table 4. The process parameters, phase compositions, mechanical properties, and strengthening mechanisms of in situ-synthesized TiN-reinforced TMCCs/TMCs.

Method	Material	Laser Energy/Power	Scanning Speed (mm/s)	Phase	Hardness	Loss/ Wear Rate	Coefficient of Friction	Yield Strength (MPa)	Ultimate Tensile Strength (MPa)	Maximum Strain or Elongation (%)	Strengthening Mechanism	Ref.
LN	N2	100 mJ	0.6	(α-Ti) + TiN	286.62 HV0.1	4.01 × 10−4 mm3/Nm	0.62				SSS; SPS	[152]
		200 mJ			322.83 HV0.1	2.89 × 10−4 mm3/Nm	0.63
		250 mJ			352.19 HV0.1	2.00 × 10−4 mm3/Nm	0.66
		300 mJ			435.55 HV0.1	1.44 × 10−4 mm3/Nm	0.64
		400 mJ			560.90 HV0.1	8.01 × 10−4 mm3/Nm	0.62
LN	N2	25 W	40	Ti + TiN	1277 HV						SPS; GRS	[151]
LN	100 vol% N2	175 W	100	(α-Ti) + TiN + TiN0.26	7.9 GPa						SSS; SPS; GRS	[153]
	100 vol% N2		400		5.6 GPa
	100 vol% N2		1600		5.3 GPa
	75 vol% N2		400		3.7 GPa
	50 vol% N2		400		3.6 GPa
	20 vol% N2		400		3.4 GPa
	0 vol% N2		400		3.3 GPa
LN	N2	650 W	6	(α-Ti) + (α′-Ti) + (β-Ti) + TiN	320 HV						SSS; SPS; GRS; DLS	[154,155]
LN	TC4 + 1 min N2	180 W	200	(α-Ti) + (α′-Ti) + TiN + AlN	395 HV						GRS; DPS; PPS; IFS	[150]
	TC4 + 2 min N2				440 HV
	TC4 + 5 min N2				511 HV
LPBF	CP-Ti + 0 vol% N2	110–170 W	400–1200	(α-Ti) + (α′-Ti) + (β-Ti) + TiN	235 HV0.2	1.51 × 10−3 mm3/Nm	0.39				SSS; SPS; GRS	[156]
	CP-Ti + 5 vol% N2				286 HV0.2	7.1 × 10−4 mm3/Nm	0.37
	CP-Ti + 10 vol% N2				345 HV0.2	6.5 × 10−4 mm3/Nm	0.35
	CP-Ti + 15 vol% N2				389 HV0.2	4.5 × 10−4 mm3/Nm	0.32
LPBF	0% N2	160 W	1000	(α-Ti) + (α′-Ti) + (β-Ti) + TiN	239 HV0.2			599.19	696.48	33.6	SSS; SPS; GRS; IFS	[55]
	5% N2		500		290 HV0.2			886.17	958.83	17.27
	10% N2				330 HV0.2			907.64	1006.6	2.37
	15% N2				360 HV0.2			1032.29	1132.39	1.03
DED	TC4 + 0% N2	380 W	14.16	(α-Ti) + (β-Ti) + TiN	350 HV0.3			840			SSS; GRS; DPS; PPS; DLS	[157]
	TC4 + 2% N2				540 HV0.3			1050
	TC4 + 10% N2				550 HV0.3			1600		15

Table 4. The process parameters, phase compositions, mechanical properties, and strengthening mechanisms of in situ-synthesized TiN-reinforced TMCCs/TMCs.

Method	Material	Laser Energy/Power	Scanning Speed (mm/s)	Phase	Hardness	Loss/ Wear Rate	Coefficient of Friction	Yield Strength (MPa)	Ultimate Tensile Strength (MPa)	Maximum Strain or Elongation (%)	Strengthening Mechanism	Ref.
LN	N2	100 mJ	0.6	(α-Ti) + TiN	286.62 HV0.1	4.01 × 10−4 mm3/Nm	0.62				SSS; SPS	[152]
		200 mJ			322.83 HV0.1	2.89 × 10−4 mm3/Nm	0.63
		250 mJ			352.19 HV0.1	2.00 × 10−4 mm3/Nm	0.66
		300 mJ			435.55 HV0.1	1.44 × 10−4 mm3/Nm	0.64
		400 mJ			560.90 HV0.1	8.01 × 10−4 mm3/Nm	0.62
LN	N2	25 W	40	Ti + TiN	1277 HV						SPS; GRS	[151]
LN	100 vol% N2	175 W	100	(α-Ti) + TiN + TiN0.26	7.9 GPa						SSS; SPS; GRS	[153]
	100 vol% N2		400		5.6 GPa
	100 vol% N2		1600		5.3 GPa
	75 vol% N2		400		3.7 GPa
	50 vol% N2		400		3.6 GPa
	20 vol% N2		400		3.4 GPa
	0 vol% N2		400		3.3 GPa
LN	N2	650 W	6	(α-Ti) + (α′-Ti) + (β-Ti) + TiN	320 HV						SSS; SPS; GRS; DLS	[154,155]
LN	TC4 + 1 min N2	180 W	200	(α-Ti) + (α′-Ti) + TiN + AlN	395 HV						GRS; DPS; PPS; IFS	[150]
	TC4 + 2 min N2				440 HV
	TC4 + 5 min N2				511 HV
LPBF	CP-Ti + 0 vol% N2	110–170 W	400–1200	(α-Ti) + (α′-Ti) + (β-Ti) + TiN	235 HV0.2	1.51 × 10−3 mm3/Nm	0.39				SSS; SPS; GRS	[156]
	CP-Ti + 5 vol% N2				286 HV0.2	7.1 × 10−4 mm3/Nm	0.37
	CP-Ti + 10 vol% N2				345 HV0.2	6.5 × 10−4 mm3/Nm	0.35
	CP-Ti + 15 vol% N2				389 HV0.2	4.5 × 10−4 mm3/Nm	0.32
LPBF	0% N2	160 W	1000	(α-Ti) + (α′-Ti) + (β-Ti) + TiN	239 HV0.2			599.19	696.48	33.6	SSS; SPS; GRS; IFS	[55]
	5% N2		500		290 HV0.2			886.17	958.83	17.27
	10% N2				330 HV0.2			907.64	1006.6	2.37
	15% N2				360 HV0.2			1032.29	1132.39	1.03
DED	TC4 + 0% N2	380 W	14.16	(α-Ti) + (β-Ti) + TiN	350 HV0.3			840			SSS; GRS; DPS; PPS; DLS	[157]
	TC4 + 2% N2				540 HV0.3			1050
	TC4 + 10% N2				550 HV0.3			1600		15

An excessive nitrogen concentration and supersaturation can lead to the formation of coarse TiN particles and pores (Figure 8b) [150,153,156]. Additionally, under the rapid cooling characteristics of laser processing, the excessive solid solution of nitrogen can lower the β→α phase transformation temperature, potentially resulting in residual β-Ti or even the formation of non-equilibrium α′ martensite [55,150,154,155,156].

When the laser energy density is too low, the solubility and diffusion ability of nitrogen are limited, resulting in fewer TiN nucleation sites. When the laser energy density is too high, some TiN particles may coarsen or remelt due to prolonged exposure to high temperatures, reducing the effective volume fraction of TiN. At the same time, the proportion of residual β-Ti in the matrix may increase (Figure 8c) [55,152,153].

4.4.2. Mechanical Properties and Strengthening Mechanisms

From Table 4, it can be seen that the TiN phase can effectively improve the strength, wear resistance, and other properties of the coating, but these properties are significantly influenced by the nitrogen concentration and laser process parameters (Figure 8d,e). These excellent mechanical properties are attributed to multiple strengthening mechanisms. Clearly, the strengthening mechanism of TiN is similar to that of TiC, which is also a particulate reinforcement, so the similar parts will not be repeated here.

In the strengthening mechanisms of TiC and TiB, solid solution strengthening generally plays a relatively weak role, mainly because the content of solute atoms is low and does not form a sufficient amount of solid solution. In situ-synthesized TiN-reinforced TMCCs often use nitriding processes. When enough nitrogen gas enters the molten pool, it first decomposes into active nitrogen atoms or ions, which then diffuse and convect into the molten titanium matrix, forming a supersaturated solid solution [55,149,150,151,152,153,154,155,156,157]. It is worth noting that the specific solid solution situation depends on the molten pool temperature and N concentration. This solid solution interacts with dislocations in the material in various ways, creating atmospheres that pin or drag dislocations. The strengthening of in situ-synthesized TiN TMCCs/TMCs mainly comes from dislocation strengthening dominated by solid solution strengthening, grain refinement strengthening, and load transfer strengthening (Figure 8f) [55].

At the same time, TiN-reinforced TMCCs often exhibit excellent corrosion resistance [153,155]. TiN, similar to TiB and TiB₂, promotes the formation of a highly stable passive film in the early stages, thereby enhancing the material’s corrosion resistance. Zhao et al. [158] conducted comparative experiments on TiN/Ti composites and pure Ti fabricated by SLM to evaluate their corrosion resistance. They found that TiN acts as a micro-cathode, accelerating the anodic reaction, while the depletion of Ti helps maintain the passive state. Therefore, the corrosion current density of TiN/Ti composites decreases faster with the corrosion potential compared to pure Ti. As the TiO₂ passive film and insoluble products accumulate, they isolate the matrix from the solution, inhibiting the further dissolution of the Ti matrix and ultimately slowing down the corrosion process to some extent. It should be noted that the corrosion resistance of different reinforcement materials varies in different corrosive environments.

4.5. TiC/TiB-Reinforced TMCCs

The successful performance enhancement achieved by in situ-synthesized single-phase ceramic has spurred interest in dual-phase or multi-phase synergistic reinforcement strategies. Among these, the TiC + TiB system has garnered significant attention. A common in situ reaction system utilizes B₄C as the boron and carbon source, reacting with titanium to form dual reinforcements through high-temperature reactions (Ti + B₄C → TiB + TiC) [15]. This approach avoids interfacial contamination and uneven distribution inherent to ex situ reinforcements while enabling precise control over phase ratios by adjusting the B₄C content, establishing a foundation for microstructure–property relationship studies [159,160,161].

Table 5. The process parameters, phase compositions, mechanical properties, and strengthening mechanisms of in situ-synthesized TiC/TiB-reinforced TMCCs/TMCs.

Method	Material	Laser Power (W)	Scanning Speed (mm/s)	Phase	Hardness	Loss/ Wear Rate	Coefficient of Friction	Yield Strength (MPa)	Ultimate Tensile Strength (MPa)	Maximum Strain or Elongation (%)	Strengthening Mechanism	Ref.
DED	CP-Ti + 12.61 wt% B4C	1500	2	(α-Ti) + TiB + TiC + TiB2 + B4C	981.2 HV0.5	3.73 × 10−5 mm3/Nm	0.1685				SPS; GRS; DLS; SGS	[162]
LMD	TC4 + 20 wt% B4C	1000	16.67	(α-Ti) + (β-Ti) + TiB + TiC + TiB2	470 HV0.3						GRS; DPS; DLS; SGS	[163]
		1200			548 HV0.3
		1400			461 HV0.3
LC	CP-Ti + 30 wt% B4C	450	3	(α-Ti) + (β-Ti) + TiB + TiC + AlTi3 + TiVC2	758.0 HV0.2	2.31 × 10−1 mm3	0.32				SPS; GRS; SGS	[164]
					836.4 HV0.2	1.44 × 10−1 mm3	0.28
					896.1 HV0.2	1.02 × 10−1 mm3	0.16
					785.9 HV0.2	1.61 × 10−1 mm3	0.48
DED	CP-Ti + 10 wt% B4C	350	4.23	Ti + TiB + TiC + TiB2 + B4C + B25C	300 HV0.2	4.4 × 10−4 mm3/Nm					GRS; DPS	[165]
	CP-Ti + 15 wt% B4C				500 HV0.2	3.9 × 10−4 mm3/Nm
	CP-Ti + 20 wt% B4C				1000 HV0.2	1.2 × 10−4 mm3/Nm
SLM	CP-Ti + 0 wt% B4C	250	1500	(α-Ti) + TiB + TiC	205.07 HV0.1			499.05	582.39	26.37	GRS; DPS; SGS	[166]
	CP-Ti + 1 wt% B4C				275.06 HV0.1			761.77	945.99	4.75
	CP-Ti + 2 wt% B4C				320.18 HV0.1			744.66	834.14	1.19
	CP-Ti + 3 wt% B4C				343.08 HV0.1				627.96	0.56
	CP-Ti + 5 wt% B4C				424.83 HV0.1
LPBF	CP-Ti + 0 wt% TiB/TiC	160	1200	(α-Ti) + (β-Ti) + TiB + TiC	173.1 HV			489.1	540.7	20.9	GRS; SSS; DPS; LTS; SGS	[69]
	CP-Ti + 2.5 wt% TiB/TiC				228.4 HV			668.1	803.1	11.0
	CP-Ti + 5 wt% TiB/TiC				275.9 HV			777.1	940.3	6.1
	CP-Ti + 7.5 wt% TiB/TiC				298.5 HV			828.8	999.2	3.7
	CP-Ti + 10 wt% TiB/TiC				309.9 HV			869.7	1059.2	2.9
DED	TC4 + 0%B4C	3200	12	(α-Ti) + (β-Ti) + TiB + TiC	375.7 HV	0.00044 mm3/Nm	0.42	920	1020	4.61	GRS; DPS; SSS; LTS	[134]
	TC4 + 16.7%B4C				440.2 HV	0.00038 mm3/Nm	0.36	1140	1200	2.27
	TC4 + 25%B4C				455.2 HV	0.00036 mm3/Nm	0.33	1160	1230	2.09
	TC4 + 14.3%B4C + 42.9%C				478.1 HV	0.00041 mm3/Nm	0.37	1100	1160	1.89
DED	TC4 + 0 wt% B4C	3300	3.33	(α-Ti) + (β-Ti) + TiB + TiC	344.0 HV0.5				983.3	8.2	LTS; OWS; GRS; SGS	[167]
	TC4 + 5 wt% B4C				414.0 HV0.5				1126.1	4.2
LPBF	TC4 + 0 wt% B4C	200	1000	(α-Ti) + (α′-Ti) + (β-Ti) + TiB + TiC	380 HV						GRS; LTS; OWS; DLS; SSS; SGS	[70]
	TC4 + 0.2 wt% B4C	250	1000		395 HV
DED	Ti65 + 0.1 wt% B4C	2000	4	(α-Ti) + (β-Ti) + TiB + TiC				1024	1135	11.40	GRS; LTS; SSS	[168]
	Ti65 + 0.2 wt% B4C							1048	1142	10.19
	Ti65 + 0.3 wt% B4C							1085	1201	5.54
	Ti65 + 0.5 wt% B4C							1138	1155	1.28

Note: SGS: synergistic strengthening.

summarizes the process parameters, phase compositions, mechanical properties, and strengthening mechanisms of representative in situ-synthesized TiC/TiB-reinforced TMCCs/TMCs.

4.5.1. Phase and Microstructure

The main phase composition of in situ-synthesized TiC/TiB-reinforced TMCCs via LC includes α-Ti, β-Ti, TiC, and TiB (Figure 10a). During the solidification process using B₄C as the boron and carbon source, although the Gibbs free energy of formation for TiB is higher than that of TiC, TiB preferentially precipitates over TiC due to the higher B concentration (C/B ratio: 1/4) and chemical potential at the interface between the titanium melt and B₄C powder [134,166,169]. However, this is not true in all cases, and is also related to the concentration of the B or C. Figure 11 illustrates the phase transformation process based on the Ti-C and Ti-B phase diagram (Figure 3a,b). In the hypoeutectic region (Line a1 and Line b1), β-Ti solidifies as the primary phase. As the temperature decreases to the eutectic transformation temperature, TiB whiskers and TiC particles precipitate. In the hypereutectic region (Line a2 and Line b2), both TiC and TiB can serve as the primary precipitate phase, depending on the C/B ratio. Taking the case of a sufficiently high B content as an example, TiB whiskers form as the primary phase, followed by heterogeneous nucleation and the growth of TiC on the surface of TiB. Simultaneously, both TiB and TiC precipitates act as nucleation sites for β-Ti, promoting its formation. When the temperature drops below the β-Ti transformation temperature, β-Ti transforms into α-Ti.

As the B₄C content increases, the thermal convection pattern within the molten pool undergoes significant changes, leading to the evolution of the microstructure from lamellar α-Ti to dendritic structures, and subsequently, to a mixed structure of cellular and dendritic grains (Figure 10d) [69,134,165,166,168]. When the content is excessively high, this mixed structure coarsens and disintegrates. This microstructural evolution is primarily attributed to changes in thermal gradients and cooling rates within the molten pool [69,134]. Specifically, as the B₄C content increases, localized “thermal resistance zones” form in the melt pool. On the one hand, this significantly amplifies thermal gradients near B₄C particles, generating an inhomogeneous temperature distribution that induces differential growth rates of crystals along distinct orientations, thereby destabilizing crystalline growth. On the other hand, the addition of B₄C reduces the melt pool’s cooling rate, extending the solidification duration and facilitating sustained grain growth, ultimately resulting in the coarsening and disintegration of the hybrid microstructure. At a lower B₄C content, the thermal convection in the molten pool is relatively uniform, with TiB whiskers and TiC particles distributed in a dendritic form, forming a well-defined network structure [70,166,168]. As the B₄C content increases, Marangoni convection drives the TiB whiskers and TiC particles to form clustered aggregates through mechanical interlocking, further restricting the growth of matrix grains and refining the α-Ti grains (Figure 10b,c) [69,134,162,165]. However, the excessive addition of B₄C increases the molten pool viscosity and induces a localized thermal stress concentration, resulting in defects such as pores and cracks, which reduce the material’s density [69,134].

4.5.2. Mechanical Properties and Strengthening Mechanisms

Through Table 5, it can be seen that the synergistic reinforcement of TiC and TiB significantly enhances the mechanical properties of TMCCs/TMCs. This not only stems from the individual strengthening effects of multiple reinforcements, but also arises from the unique structures they form.

On the one hand, the high aspect ratio of TiB whiskers makes them show an excellent load-bearing capacity during the load transfer process, which is exactly the advantage shown when TiB is used as a single reinforcement, although few researchers have proposed the structural strengthening of the material during the synergistic effect, but it is essentially due to the special geometry of TiB [69,70,134,167,168,169,170,171,172].

Meanwhile, the high hardness of TiC particles and the high strength of TiB whiskers work together to make the composites exhibit excellent deformation resistance when subjected to external loads, which is largely attributed to the dispersion strengthening [69,70,134,162,163,164,165,166,167,168,169,170,171,172,173].

What is more, it is worth proposing that, compared with a single reinforcement, the cluster aggregates formed by the mechanical interlocking of TiB whiskers and TiC particles further refine the grain size of the matrix, which enhances the effect of grain refinement strengthening and synchronously improves the strength and plastic toughness of the material (Figure 10c) [69,70,134,166].

In terms of tribological properties, the reticular distribution formed by TiB whiskers and TiC particles significantly improved the wear resistance of the composites by suppressing abrasive, oxidative wear and load transfer strengthening [134,164,165,173].

Regarding the strengthening mechanism, Han et al. [166] performed molecular dynamics simulations in order to investigate the strengthening mechanism of the materials, and interesting mechanisms were obtained. Figure 10e shows the tensile force applied along the z-axis direction, and it was found that the internal stresses were concentrated at the grain boundaries and pore tips, and the stresses on the grain boundaries were reduced after the introduction of TiC. This is due to the diffusion of C atoms from the TiC particles to the Ti matrix during deformation, resulting in the formation of vacancies in the particles. This movement of C atoms may be due to the formation of complex low-energy interstitial defects and the movement of these defects during the violent deformation process. The formation of these vacancies partially releases the accumulated stresses in the TiC particles, leading to stress reduction in the GB region.

4.6. TiB/TiN-Reinforced TMCCs

In situ-synthesized TiB/TiN composite reinforcements significantly enhance the comprehensive performance of TMCCs/TMCs through unique synergistic effects. TiB exhibits excellent thermodynamic compatibility and interfacial bonding strength with the titanium matrix, effectively improving the tensile strength, hardness, and creep resistance [117,118,119,120]. Meanwhile, TiN enhances the wear resistance and high-temperature stability due to its high melting point, superior fracture toughness, and biocompatibility [174,175,176]. Compared to directly adding both reinforcements, in situ synthesis generates TiB and TiN directly within the matrix via chemical reactions, eliminating interfacial contamination and particle agglomeration, thereby optimizing the reinforcement distribution and interfacial quality [177]. Research commonly employs BN/Ti powder mixtures as feedstock. Unfortunately, the preparation of TiB/TiN-reinforced TMCCs on titanium alloys has not been sufficiently and extensively studied.

Table 6. The process parameters, phase compositions, mechanical properties, and strengthening mechanisms of in situ-synthesized TiB/TiN-reinforced TMCCs/TMCs.

Method	Material	Laser Power (W)	Scanning Speed (mm/s)	Phase	Hardness	Loss/ Wear Rate	Coefficient of Friction	Yield Strength (MPa)	Ultimate Tensile Strength (MPa)	Maximum Strain or Elongation (%)	Strengthening Mechanism	Ref.
LENS	TC4 + 5 wt% BN	400	10	(α-Ti) + (β-Ti) + TiB + TiN + BN	604 HV1.0	4.26 × 10−5 mm3/Nm	0.47				DPS; SS; GRS; IFS	[178,179]
		300			570 HV1.0	8.50 × 10−5 mm3/Nm	0.48
		400	20		568 HV1.0	7.50 × 10−5 mm3/Nm	0.46
		300			543 HV1.0	1.51 × 10−4 mm3/Nm	0.44
	TC4 + 15 wt% BN	400	10		877 HV1.0	1.90 × 10−6 mm3/Nm	0.49
		300			765 HV1.0	4.50 × 10−6 mm3/Nm	0.45
		400	20		816 HV1.0	6.20 × 10−6 mm3/Nm	0.49
		300			733 HV1.0	4.80 × 10−6 mm3/Nm	0.46
LENS	CP-Ti + 0 wt% BN	427	12.7	(α-Ti) + (α′-Ti) + TiB + TiN + TiN0.176 + BN	256.9 HV0.2						SSS; SPS; GRS	[180]
	CP-Ti + 3 wt% BN	475	10.6		538.1 HV0.2
	CP-Ti + 6 wt% BN	500	10.6		584.7 HV0.2
LC	TC4 + 2 wt% BN	1600	16.7	(α-Ti) + (β-Ti) + TiB + TiN	650 HV0.2		0.55				SPS; SSS; PPS; GRS; SGS	[181]
	TC4 + 4 wt% BN				850 HV0.2		0.65
	TC4 + 6 wt% BN				920 HV0.2		0.75
LC	Ti60 + 25 wt% BN	1500	150	(α-Ti) + (β-Ti) + TiB + TiN + BN	700 HV0.3	6.5 × 10−5 mm2	0.55				SSS; SPS; SGS	[182]
			200		650 HV0.3	7.5 × 10−5 mm2	0.52
			250		610 HV0.3	7.9 × 10−5 mm2	0.54
		2000	150		1100 HV0.3	4.4 × 10−5 mm2	0.61
			200		950 HV0.3	6.0 × 10−5 mm2	0.55
			250		1000 HV0.3	6.1 × 10−5 mm2	0.53
		2500	150		1200 HV0.3	3.5 × 10−5 mm2	0.68
			200		1100 HV0.3	4.1 × 10−5 mm2	0.68
			250		1000 HV0.3	6.4 × 10−5 mm2	0.66
DED	TC4 + 2 wt% BN	1600	10	(α-Ti) + (β-Ti) + TiB + TiN	530 HV0.2	1.61 × 10−4 mm3/Nm	0.44	888.51		4.45	SPS; SSS; DLS; GRS; SGS	[183]
		2400			480 HV0.2	1.06 × 10−4 mm3/Nm	0.41	892.16		5.81
LPBF	CP-Ti + 0 vol% BN			(α-Ti) + TiB + TiN	203.19 HV				519.8		GRS; LTS; SSS; SPS; SGS	[184]
	CP-Ti + 1.5 vol% BN				369.45 HV				508.7
LPBF	CP-Ti + 0.5 wt% BN	250	1500	(α-Ti) + TiB + TiN + BN				1020	1040	3.2	GRS; DPS; SSS	[185]

summarizes the process parameters, phase compositions, mechanical properties, and strengthening mechanisms of representative in situ-synthesized TiB/TiN-reinforced TMCCs/TMCs.

4.6.1. Phase and Microstructure

The main phase composition of in situ-synthesized TiB/TiN-reinforced TMCCs via LC includes α-Ti, β-Ti, TiB, and TiN (Figure 12a). When BN is used as the boron and nitrogen source, BN decomposes into B and N under the high-temperature molten pool induced by laser irradiation and undergoes in situ reactions with titanium [186]. The reaction pathway is primarily influenced by the distribution uniformity of B and N and the dynamics of the molten pool. Figure 13 illustrates the phase transformation process based on the Ti-B and Ti-N phase diagram (Figure 3b,c). When the N concentration lies along Line c1 and the B concentration lies along Line b2, α-Ti precipitates as the primary phase. Subsequently, as the temperature decreases to the peritectic temperature, a peritectic reaction occurs, forming β-Ti. From the perspective of the Ti-B system, β-Ti may also precipitate directly from the liquid phase, followed by a eutectic reaction forming β-Ti and TiB. When the N concentration falls within the Line c3 or Line c4 region and the B concentration is in the hypereutectic region (Line b2), TiN precipitates as the primary phase while simultaneously acting as a heterogeneous nucleation site for α-Ti formation. As the melt pool temperature gradually decreases, the peritectic temperature is reached first, triggering the peritectic reaction where α-Ti grows attached to TiN surfaces. For the N concentration along Line c3, the residual liquid phase remains. With the increasing B concentration, as the temperature further decreases to the liquidus, TiB nucleates and grows. Upon reaching the eutectic reaction temperature, β-Ti and TiB precipitate from the remaining liquid. For the N concentration along Line c4, the peritectic reaction would typically consume all liquid, but the elevated B concentration promotes TiB precipitation. With further cooling, the system undergoes a eutectic reaction, forming β-Ti. In all above scenarios, β-Ti transforms into α-Ti upon continued cooling. Concurrently, due to the decreasing solid solubility of N in α-Ti with a further temperature reduction, Ti₂N precipitates within the system. It should be stressed that most in situ-synthesized TiB/TiN-reinforced TMCCs/TMCs fail to yield conclusions fully consistent with the above descriptions. This is primarily attributed to the influence of other alloying elements in the system and processing parameters on the melt pool solidification behavior, which aligns with the earlier descriptions for single-reinforcement systems.

At a lower BN content, the molten pool develops interlocked structures of TiB and TiN, similar to those observed in TiC/TiB-reinforced TMCCs/TMCs [178,181,185]. When the BN content reaches a critical level, the broader diffusion range of nitrogen facilitates the formation of TiN dendritic cores, while boron, due to its lower diffusion rate, forms a TiB network around TiN under intense Marangoni convection in the molten pool (Figure 12b). Figure 12d illustrates this formation mechanism.

At a low laser energy density, an insufficient molten pool temperature results in residual unmelted BN particles, creating localized defects [178,182,183]. Conversely, an excessive energy density induces keyhole effects due to a high thermal input, increasing the porosity. Optimizing the BN content and energy density achieves a near-full density and forms a well-defined network structure (Figure 12c).

4.6.2. Mechanical Properties and Strengthening Mechanisms

From Table 6, it can be seen that in situ-synthesized TiN and TiB significantly enhance the mechanical properties of TMCs. Through the analysis of numerous researchers, it is evident that the TiB/TiN reinforcement system and the TiC/TiB reinforcement system share significant similarities in the geometric shapes of the reinforcements and the resulting microstructures. This is because, in both systems, the TiB rod or whisker precipitates earlier than TiC or TiN, serving as heterogeneous nucleation sites for a TiC or TiN particle. Additionally, the characteristics of the laser process result in strong Marangoni convection within the molten pool, promoting the formation of cellular structures.

However, there are differences. On the one hand, the elemental ratios in B₄C and BN used for in situ synthesis are different, which affects the distribution of reinforcements in the matrix. On the other hand, N has a much higher solubility than C, which significantly enhances the solid-solution-strengthening effect in the TiB/TiN reinforcement system [187]. Therefore, the high strength of TiB/TiN-reinforced TMCCs/TMCs primarily stems from the following mechanisms.

(1)

Grain refinement strengthening. Multi-scale causes, as described earlier, lead to the grain size refinement of the matrix and reinforcing phases, effectively hindering the dislocation motion [178,179,180,181,182,183,184,185,188,189].

(2)

Solid solution strengthening. Nitrogen and boron atoms (especially nitrogen atoms) dissolve into the Ti matrix, interacting with dislocations and forming atmospheres that pin dislocations [180,181,182,183,184,185].

(3)

Load transfer strengthening. TiB and TiN (especially TiB with a large aspect ratio) effectively improve the load transfer efficiency [184].

(4)

Dispersion strengthening. The uniform distribution of hard TiB and TiN particles triggers the Orowan bypass mechanism [178,179,180,181,182,183,184,185,188,189].

(5)

Structure strengthening. The interconnected network structure formed by TiB and TiN effectively enhances the material’s deformation capability [178,179].

4.7. Ti₅Si₃/TiN-Reinforced TMCCs

Ti₅Si₃, as an important intermetallic compound, exhibits significant reinforcement effects in TMCCs/TMCs [79,188]. Its high melting point and excellent thermal stability enable it to maintain mechanical properties under high-temperature environments, making it suitable for high-temperature structural materials [189]. Additionally, Ti₅Si₃ possesses high hardness, effectively enhancing the wear resistance and deformation resistance of composites. Furthermore, the strong interfacial bonding between Ti₅Si₃ and the titanium matrix minimizes crack initiation at interfaces, thereby improving the overall toughness and fatigue performance of the composites [190]. However, research on Ti₅Si₃-reinforced TMCCs/TMCs remains relatively limited, primarily due to the complexity of its formation mechanism: the low diffusion rate of silicon in the titanium matrix and the need for the precise control of the reaction temperature and cooling rate to avoid the formation of coarse or uneven reinforcement phases [191]. Moreover, the inherent brittleness of Ti₅Si₃ may lead to reduced toughness if the phase is unevenly distributed or excessively large. Recent studies have focused on obtaining silicon sources from other compounds, with the in situ synthesis of Ti₅Si₃ gaining attention. Using Si₃N₄, the simultaneous generation of Ti₅Si₃ and TiN can be achieved, enabling synergistic reinforcement in TMCCs/TMCs. Figure 14c depicts this process. But up to now, the study of Ti₅Si₃/TiN-reinforced TMCCs has not been sufficiently in-depth, especially on titanium alloys. Table 7 summarizes the process parameters, phase compositions, mechanical properties, and strengthening mechanisms of Ti₅Si₃/TiN-reinforced TMCCs/TMCs.

Table 7. The process parameters, phase compositions, mechanical properties, and strengthening mechanisms of Ti₅Si₃/TiN-reinforced TMCCs/TMCs.

Method	Material	Laser Power (W)	Scanning Speed (mm/s)	Phase	Hardness	Loss/ Wear Rate	Flexural Strength (MPa)	Flexural Strain (%)	Strengthening Mechanism	Ref.
LC	CP-Ti + 12 wt% Si3N4	1400	3	(α-Ti) + TiN + Ti5Si3 + Ti3Al + Ti2N	721.4 HV0.5	0.50 mm3			SSS; DPS; IFS; GRS	[192]
	CP-Ti + 10 wt% Al + 12 wt% Si3N4				757.6 HV0.5	0.37 mm3
	CP-Ti + 20 wt% Al + 12 wt% Si3N4				781.2 HV0.5	0.29 mm3
SPS	CP-Ti + 5 wt% Si3N4			(α-Ti) + TiN + Ti5Si3	632 HV		880	0.0148	SPS	[193]
	CP-Ti + 10 wt% Si3N4				920 HV		840	0.0168
	CP-Ti + 15 wt% Si3N4				1126 HV		710	0.0168
LPBF	TC4 + 5 wt% Si3N4	96	400	(α-Ti) + (β-Ti) + TiN + Ti5Si3	860 KHN				DPS; SSS; GRS; SGS	[194]
			600		793 KHN
			900		801 KHN
			1200
SLM	CP-Ti + 24.55 wt% Si3N4	900	100	Ti + TiN + Ti5Si3					DPS; GRS; IFS	[195]
SPS	TC4 + 5 wt% Si3N4			Ti + Si3N4	600 HV				SPS; GRS; IFS	[196]
	TC4 + 10 wt% Si3N4				680 HV
	TC4 + 15 wt% Si3N4				590 HV
SPS	CP-Ti + 1 wt% Si3N4			(α-Ti) + (α-Si3N4) + (β-Si3N4) + Ti5Si3	430 HV1	0.00030 mm3/Nm	200	0.012	SSS; SPS	[197]
	CP-Ti + 1.5 wt% Si3N4				480 HV1	0.00026 mm3/Nm
	CP-Ti + 2 wt% Si3N4				610 HV1	0.00020 mm3/Nm
	CP-Ti + 2.5 wt% Si3N4				650 HV1	0.00016 mm3/Nm	650	0.017
	CP-Ti + 5 wt% Si3N4				800 HV1	0.00010 mm3/Nm	1280	0.023

Table 7. The process parameters, phase compositions, mechanical properties, and strengthening mechanisms of Ti₅Si₃/TiN-reinforced TMCCs/TMCs.

Method	Material	Laser Power (W)	Scanning Speed (mm/s)	Phase	Hardness	Loss/ Wear Rate	Flexural Strength (MPa)	Flexural Strain (%)	Strengthening Mechanism	Ref.
LC	CP-Ti + 12 wt% Si3N4	1400	3	(α-Ti) + TiN + Ti5Si3 + Ti3Al + Ti2N	721.4 HV0.5	0.50 mm3			SSS; DPS; IFS; GRS	[192]
	CP-Ti + 10 wt% Al + 12 wt% Si3N4				757.6 HV0.5	0.37 mm3
	CP-Ti + 20 wt% Al + 12 wt% Si3N4				781.2 HV0.5	0.29 mm3
SPS	CP-Ti + 5 wt% Si3N4			(α-Ti) + TiN + Ti5Si3	632 HV		880	0.0148	SPS	[193]
	CP-Ti + 10 wt% Si3N4				920 HV		840	0.0168
	CP-Ti + 15 wt% Si3N4				1126 HV		710	0.0168
LPBF	TC4 + 5 wt% Si3N4	96	400	(α-Ti) + (β-Ti) + TiN + Ti5Si3	860 KHN				DPS; SSS; GRS; SGS	[194]
			600		793 KHN
			900		801 KHN
			1200
SLM	CP-Ti + 24.55 wt% Si3N4	900	100	Ti + TiN + Ti5Si3					DPS; GRS; IFS	[195]
SPS	TC4 + 5 wt% Si3N4			Ti + Si3N4	600 HV				SPS; GRS; IFS	[196]
	TC4 + 10 wt% Si3N4				680 HV
	TC4 + 15 wt% Si3N4				590 HV
SPS	CP-Ti + 1 wt% Si3N4			(α-Ti) + (α-Si3N4) + (β-Si3N4) + Ti5Si3	430 HV1	0.00030 mm3/Nm	200	0.012	SSS; SPS	[197]
	CP-Ti + 1.5 wt% Si3N4				480 HV1	0.00026 mm3/Nm
	CP-Ti + 2 wt% Si3N4				610 HV1	0.00020 mm3/Nm
	CP-Ti + 2.5 wt% Si3N4				650 HV1	0.00016 mm3/Nm	650	0.017
	CP-Ti + 5 wt% Si3N4				800 HV1	0.00010 mm3/Nm	1280	0.023

Figure 14. (a) XRD patterns of Ti₅Si₃/TiN-reinforced TMCs [195]. (b) SEM images of the TMCs microstructure with different laser energy densities [194]. (c) Schematic representation of in situ-synthesized Ti₅Si₃/TiN-reinforced TMCs via LBPF [194].

Figure 14. (a) XRD patterns of Ti₅Si₃/TiN-reinforced TMCs [195]. (b) SEM images of the TMCs microstructure with different laser energy densities [194]. (c) Schematic representation of in situ-synthesized Ti₅Si₃/TiN-reinforced TMCs via LBPF [194].

4.7.1. Phase and Microstructure

The main phase composition of in situ-synthesized Ti₅Si₃/TiN-reinforced TMCCs via LC includes α-Ti, β-Ti, Ti₅Si₃, and TiN (Figure 14a). However, some studies have found that the reaction may be accompanied by the formation of other interface reactants, or it is difficult to form Ti₅Si₃ [192,197]. When Si₃N₄ is used as the silicon and nitrogen source, the in situ reaction between titanium and Si₃N₄ follows a thermodynamically driven mechanism. The high temperature in the molten pool promotes the decomposition of Si₃N₄ into Si and N, which subsequently react with liquid titanium [194]. Figure 15 illustrates the phase transformation process based on the Ti-N and T-Si phase diagram (Figure 3c,d). When the Si concentration lies within the hypoeutectic region (Line d1) and the N concentration lies within the Line c1 or Line c2 region, α-Ti precipitates as the primary phase. Subsequently, a peritectic reaction occurs upon cooling to the peritectic temperature, forming β-Ti. The remaining liquid phase then undergoes a eutectic reaction, precipitating Ti₅Si₃ and β-Ti. When the Si concentration lies within the Line d2 region and the N content lies within the Line c3 or Line c4 region, TiN precipitates as the primary phase. Simultaneously, the precipitated TiN acts as a nucleation core for Ti₅Si₃, promoting its nucleation and growth. As the temperature decreases to the peritectic temperature, a peritectic reaction occurs, forming α-Ti. With further cooling, the system undergoes a eutectic reaction, generating β-Ti and Ti₅Si₃. Subsequently, as the temperature continues to decrease, β-Ti transforms into α-Ti. Notably, both Si and N exhibit high solubility in the titanium matrix, resulting in the partial dissolution of Si and N atoms as solid solutions within the matrix (the solid solubility of Si is less than that of N) [198]. Due to the decreasing temperature reducing the solid solubility of Si and N, Ti₃Si and Ti₂N precipitate.

The distribution and size of the reinforcements are significantly influenced by the laser energy density (Figure 14b) [194]. At a low laser energy density, an insufficient molten pool temperature and higher cooling rates lead to incomplete reactions, causing the reinforcements to concentrate at the molten pool boundaries and form severe segregation. Additionally, too low an energy density results in poor fluidity and the incomplete melting of Si₃N₄ powder, leading to defects such as interlayer porosity and surface roughness. Conversely, at an appropriately high energy density, the molten pool temperature increases, and the cooling rate decreases, enhancing the melt fluidity. Marangoni convection promotes the uniform distribution of the reinforcing phases, yielding a higher quantity of TiN and Ti₅Si₃ with slightly larger sizes. Furthermore, the solid solution of Si and N in the titanium matrix becomes more pronounced at a lower energy density, as the slower cooling rates provide a sufficient time for solute atoms to dissolve.

4.7.2. Mechanical Properties and Strengthening Mechanisms

From Table 7, it can be seen that in situ-synthesized Ti₅Si₃ and TiN significantly enhance the hardness and wear resistance of TMCCs/TMCs, while Ti₅Si₃ and TiN synthesized by other techniques also contributed to the compressive strength [192,193,194,197]. Although this research is not yet sufficiently in-depth, the success of ex situ techniques and the effective attempts of in situ techniques both imply the application prospects of in situ-synthesized Ti₅Si₃/TiN-reinforced TMCCs/TMCs. This strengthening effect is primarily based on the strengthening mechanisms generated by the two reinforcements individually. It is worth noting the following.

(1)

Due to the high solubility of Si and N in titanium, the solid solution of Si and N in the Ti matrix causes lattice distortion, making solid solution strengthening an important part of the strengthening mechanism [194,196,199].

(2)

The TiN and Ti₅Si₃ particles, as hard phases, are uniformly dispersed, and their high distribution uniformity effectively improves the dispersion strengthening effect [194,195,196,197,198,199].

(3)

The in situ-synthesized TiN and Ti₅Si₃ exhibits strong interfacial bonding with the matrix, reducing crack initiation [194,197,198].

(4)

Grain refinement caused by various factors effectively increases the number of grain boundaries [194,196,197,198].

In addition to hardness, such composites also show potential in their wear resistance, tensile properties, and high-temperature performance. The high wear resistance of TiN and the high-temperature oxidation resistance of Ti₅Si₃ can synergistically improve the tribological performance, while the fine reinforcements can inhibit the crack propagation and enhance toughness. At the same time, the load transfer strengthening of the reinforcements may also significantly improve the tensile properties of the material. Unfortunately, LC or laser additive manufacturing technology has not been widely studied for Ti₅Si₃/TiN-reinforced TMCCs/TMCs, with plasma sintering technology being more commonly used [193,196,197,199].

4.8. Ti₅Si₃/TiC-Reinforced TMCCs

Since both Ti₅Si₃ and TiC-reinforced TMCCs/TMCs have demonstrated considerable promise, the in situ-synthesized Ti₅Si₃/TiC-reinforced TMCCs/TMCs has also been studied to a certain extent. To prepare Ti₅Si₃ and TiC at the same time, the addition of SiC powder is usually thought of [200,201]. However, it is well-known that SiC powder exhibits inherent stability and is generally incorporated as an ex situ ceramic reinforcement in TMCCs/TMCs [202,203,204]. Researchers have discovered that interfacial reactions still occur between SiC and the Ti matrix. The Gibbs free energy changes of these reactions vary with the temperature but remain negative overall, with TiC, Ti₅Si₃, Ti₃Si, and Ti₃SiC₂ identified as the most common interfacial reaction products. These findings provide a theoretical basis for utilizing SiC to synthesize in situ Ti₅Si₃/TiC-reinforced TMCCs.

Table 8. The process parameters, phase compositions, mechanical properties, and strengthening mechanisms of in situ-synthesized Ti₅Si₃/TiC-reinforced TMCCs/TMCs.

Method	Material	Laser Power (W)	Scanning Speed (mm/s)	Phase	Hardness	Loss/ Wear Rate	Coefficient of Friction	Yield Strength (MPa)	Ultimate Tensile Strength (MPa)	Maximum Strain or Elongation (%)	Strengthening Mechanism	Ref.
LC	SiC	60	1	Ti + TiC + Ti5Si3 + TiO2 + SiO2 + SiC	950 HV	550 μm	0.05				GRS; SPS	[205]
			3		740 HV	400 μm	0.09
			5		600 HV	200 μm	0.13
DED	TC4 + 17.6 wt% SiC	500	5	(α-Ti) + (β-Ti) + TiC + Ti5Si3 + SiC	393.7 HV				1241.1	15.01	GRS; PPS; SSS; IFS	[206]
DED + LC					476.8 HV				1099.4	12.90
LC	TC4 + 4 wt% SiC	1000	10	(α-Ti) + (β-Ti) + TiC + SiC	648.5 HV0.2	3.18 × 10−4 mm3/Nm	0.585				SPS; GRS; SS	[207]
		1300			715.5 HV0.2	1.99 × 10−4 mm3/Nm	0.511
		1600			730.1 HV0.2	1.57 × 10−4 mm3/Nm	0.491
		1900			691.3 HV0.2	2.61 × 10−4 mm3/Nm	0.538
LC	CP-Ti + 10 vol% SiC	600	10	Ti + TiC + Ti5Si3	527.3 HV	7.5 μm	0.45				DPS; IFS; GRS;	[208]
	CP-Ti + 20 vol% SiC				932.2 HV	0.35 μm	0.10
SLM	TC4 + 0 wt% SiC	350	1000	(α-Ti) + (β-Ti) + TiC + Ti5Si3 + Ti3Si	332 HV						GRS; LTS; SSS; OWS; CTE; DLS; SGS	[133]
	TC4 + 1.1 wt% SiC				403 HV
	TC4 + 2.2 wt% SiC				429 HV
	TC4 + 3.3 wt% SiC				576 HV
DED	TC4 + 1 vol% SiC	1200	20	(α-Ti) + (β-Ti) + TiC + Ti5Si3 + SiC	450 HV0.3			1110	1300	2.1	GRS; SPS; IFS	[209]
	TC4 + 3 vol% SiC				490 HV0.3			1030	1310	0.5
	TC4 + 5 vol% SiC				500 HV0.3				1010	0.2
	TC4 + 7 vol% SiC				510 HV0.3
SLM	CP-Ti + 1 wt% SiC	120	1000	(α-Ti) + TiC + TiSi2 + Ti3Si2 + SiC	303 HV	3.19 × 10−11 mm3/Nm	0.635	814	955	4.4	PPS; DPS; GRS; DLS; IFS; LTS; SGS	[210]
		140			344 HV	3.07 × 10−11 mm3/Nm	0.615	939	1097	11.5
		160			289 HV	3.50 × 10−11 mm3/Nm	0.629	915	1062	12.9
SLM	CP-Ti + 0 wt% SiC	150	20	(α-Ti) + Ti5Si3	466 VHN	3 × 10−4 mm3/Nm					OWS; DLS; CTE; LTS; GRS; SSS; SS; SGS	[79]
	CP-Ti + 1 wt% SiC				588 VHN
	CP-Ti + 2.5 wt% SiC				617 VHN
	CP-Ti + 5 wt% SiC				706 VHN	9 × 10−5 mm3/Nm

summarizes the process parameters, phase compositions, mechanical properties, and strengthening mechanisms of representative in situ-synthesized Ti₅Si₃/TiC-reinforced TMCCs/TMCs.

4.8.1. Phase and Microstructure

The main phase composition of in situ-synthesized Ti₅Si₃/TiC-reinforced TMCCs via LC includes α-Ti, β-Ti, Ti₅Si₃, and TiC (Figure 16a). When SiC is used as the silicon and carbon source, LC elevates the molten pool temperature above the melting point of SiC, causing its decomposition into Si and C atoms [133]. Figure 17 illustrates the phase transformation process based on the Ti-C and T-Si phase diagram (Figure 3a,d). When Si is located in the hypoeutectic or hypereutectic region (Line d1 or Line d2) and C consistently lies in the hypereutectic region (Line a2), TiC becomes the primary precipitate phase. For the Si concentration along Line d1, silicon atoms primarily exist as solid solution atoms above the eutectic temperature, with the β-phase precipitating from the liquid. Upon cooling to the eutectic transformation temperature of the Ti-C system, a eutectic reaction occurs, forming the eutectic β-phase and eutectic TiC. Subsequently, a eutectic reaction in the Ti-Si system takes place, precipitating β-Ti and Ti₅Si₃, followed by the precipitation of Ti₃Si with further cooling. For the Si concentration along Line d2, the prior precipitation of TiC promotes Ti₅Si₃ precipitation from the liquid. Subsequent cooling triggers two distinct eutectic reactions: one forming eutectic β-Ti and eutectic TiC, and another forming eutectic β-Ti and Ti₅Si₃. With continued cooling, Ti₃Si precipitates, with its specific precipitation behavior governed by the Si concentration. Notably, during these solidification processes, β-Ti transforms into α-Ti upon cooling below its phase transformation temperature.

At a high SiC content, the diffraction peak intensities of TiC and Ti₅Si₃ significantly increase and the lattice constants of α-Ti decrease, indicating intensified lattice distortion caused by the solid solution of Si and C atoms. The influence of the reinforcing phase content on the microstructure is also reflected in texture weakening and recrystallization behavior [79,133,208,209]. At a low SiC content, α-Ti exhibits a strong (0001) basal texture, whereas at higher SiC contents, the texture strength markedly decreases and α-Ti tends to grow equiaxially [73,126,195,196]. This is attributed to the synergistic effects of TiC and Ti₅Si₃ (Figure 16d). TiC provides heterogeneous nucleation sites, disrupting the epitaxial growth of columnar grains. Ti₅Si₃ stores strain energy by hindering dislocation motion, thereby promoting recrystallization.

4.8.2. Mechanical Properties and Strengthening Mechanisms

From Table 8, it can be seen that in situ-synthesized Ti₅Si₃ and TiC enhance the mechanical properties of TMCCs/TMCs. The strengthening mechanisms of the Ti₅Si₃ and TiC system are similar to those of the Ti₅Si₃ and TiN system, which is mainly due to the fact that TiC and TiN have similar geometries and both preferentially precipitate out of the molten pool over Ti₅Si₃. The difference is that the solid solubility of C in titanium is very low, whereas the solid solubility of N in titanium is much higher than that of Si, which largely influences the impact of the different strengthening mechanisms. The strengthening mechanisms of Ti₅Si₃/TiC-reinforced TMCCs/TMCs can be summarized as follows:

(1)

The multi-faceted grain refinement effect of Ti₅Si₃ and TiC particles effectively reduces the grain size, leading to simultaneous improvements in the hardness, strength, and toughness of the material or coating [79,133,200,201,205,206,207,208,209,210,211].

(2)

Dissolved Si and C atoms induce lattice distortion, which impedes dislocation motion and contributes to solid solution strengthening. However, the strengthening effect is limited due to the low solubility of C in the titanium matrix [79,133,206].

(3)

The fine chain-like Ti₅Si₃ particles and dispersed TiC particles create strong second-phase strengthening by effectively obstructing dislocation movement and inhibiting crack propagation (Figure 16d) [79,133,200,201,205,206,207,208,209,210,211].

(4)

The hard Ti₅Si₃ particles and TiC particles are diffusely distributed and establish a strong interfacial bond with the matrix, which effectively improves the load transfer strengthening [79,133,206,208,209,210,211].

(5)

Strong interfacial bonding also minimizes particle pull-out during the service, while the fine particle size reduces third-body wear damage when particles do detach. This combination of factors effectively prevents wear-induced failure mechanisms, contributing to the coating’s excellent wear resistance [79,205,207,208,210].

4.9. Other TMCCs

In addition to the above-mentioned single and binary in situ-synthesized TMCCs, there are still some less widely studied, such as WC, Al₂O₃, TiC/TiN, Ti₅Si₃/TiB-reinforced TMCCs, etc. Moreover, the research on ternary and higher-order multicomponent synergistic reinforced TMCCs has also attracted extensive attention. Researchers primarily investigate multicomponent synergistic reinforced TMCCs/TMCs through the following three strategies, as shown in Figure 18.

The first strategy focuses on the multiphase coupling of traditional ceramic phases to achieve a superior performance. Traxel et al. [212] fabricated in situ-reinforced TMCs by introducing B₄C and BN (total 5 wt%) into TC4 via SLM. The composite achieved a relative density of 98.3% and microhardness of 587–811 HV_0.2/15, significantly higher than pure TC4. High-temperature oxidation resistance at 850 °C for 50 h revealed a 39% reduction in mass gain, with an oxide layer thickness reduced to 27.3 μm compared to 102.1 μm for pure TC4. Monisha et al. [200] developed multiphase TMCCs using Boropak (50% B₄C, 30% SiC, 20% KBF₄) on a titanium substrate via laser technology. The coating consisted of TiB₂, TiB, TiC, Ti₅Si₃, and unmelted B₄C and SiC particles uniformly distributed in the titanium matrix, forming dendritic and coral nanostructures. The coating exhibited an average microhardness of 800–2200 HV_0.2 and a wear rate of 1.623 × 10⁻⁴ mm³/Nm, which is six times lower than untreated titanium.

The second strategy involves adding rare earth elements to unary or binary systems to further enhance the performance. Feng et al. [213] in situ-synthesized Ti₃Al/TiB-reinforced TMCCs on a TC4 substrate via LC and investigated the effect of LaB₆ addition. The incorporation of 3.0 wt% LaB₆ led to the formation of nano-sized La₂O₃ particles within the coating, which effectively refined the TiB and Ti₃Al distribution and alleviated a residual stress concentration through enhanced heterogeneous nucleation. This microstructural optimization resulted in a significant improvement in mechanical and high-temperature properties. The coating exhibited a microhardness of 917 HV_0.5, which represented a 16% increase and showed a 20% reduction in the wear mass loss along with a friction coefficient of 0.23. Furthermore, the weight gain after oxidation at 700 °C for 60 h was only 25% of that of the TC4 substrate, and the oxide layer thickness was reduced by 22%. Liu et al. [214] developed CeO₂-reinforced TMCCs on the Ti811 substrate via LC. With 2 wt% CeO₂, the coating achieved the optimal performance: microhardness of 811.67 HV_0.5 (114% higher than the substrate), a stable friction coefficient below 0.45, and wear depth of 56 μm (44% lower than the substrate). CeO₂ enhanced the hardness and wear resistance through grain refinement, grain boundary strengthening, and optimized reinforcement phase distribution, while improving the interfacial bonding strength.

The third strategy leverages synergistic effects between ceramic phases and intermetallic compounds. Li et al. [215] fabricated Ti/TiBCN-reinforced TMCCs on the TC4 substrate via LC. At a 60 wt% TiBCN content, the coating comprised dendritic or rod TiBCN phases, blocky TiC, layered TiN, TiB₂, and minor Al₃Ti and TiAl, forming a multiscale reinforced structure. The microhardness reached 1596 HV (4.6 times higher than the substrate), with the corrosion current density reduced to 4.035 × 10⁻⁵ A/cm² (one order of magnitude lower than the substrate) and a wear mass loss of 1.22 g (9/50 of that of the substrate). The coating demonstrated enhanced surface hardness, corrosion resistance, and wear resistance through synergistic strengthening by in situ-synthesized TiC/TiN/TiB₂ ceramic phases and Al₃Ti/TiAl intermetallic compounds. Pu et al. [216] synthesized TiC/TiN/SiC-reinforced Ti₃Al intermetallic compound coatings on pure titanium via LC. When the reinforcement content is 40 vol%, the coating exhibits the optimal comprehensive performance. The microhardness reaches 1124 HV, representing a 63% increase compared to the pure Ti₃Al coating. Under a 5 N load, the friction coefficient remains stable in the range of 0.70–0.75 and its wear volume is significantly smaller than those of both pure Ti₃Al and pure SiC coatings.

It is noteworthy that excessive multicomponent additions do not necessarily yield synergistic strengthening. On the one hand, excessive material types complicate interfacial conditions in TMCCs, making unfavorable interfacial reactions difficult to control and a compromising interfacial bonding strength and stability. On the other hand, excessive components may induce agglomeration or segregation, leading to significant local performance variations. Additionally, the thermal expansion coefficient matching remains a fundamental principle in a mixed ceramic phase design. Excessive component types may generate high local thermal stresses due to mismatched thermal expansion coefficients, potentially initiating cracks.

5. Summary, Challenges, and Prospects

Through an analysis of the available literature on the in situ-synthesized ceramic-reinforced TMCCs/TMCs via laser technology, we have identified their significance and systematically elucidated their technical principles, phase evolution mechanisms, and strengthening mechanisms. The in situ-synthesized ceramic phases (such as TiC, TiB, and TiN, either singular or multiple) establish strong metallurgical bonds with the titanium matrix. Through strengthening mechanisms, including grain refinement strengthening, second-phase strengthening, solid solution strengthening, load transfer strengthening, synergistic strengthening, and so on, these composites demonstrate remarkable improvements in the hardness, strength, wear resistance, and corrosion resistance, while also exhibiting promising potential in biocompatibility, oxidation resistance, high-temperature performance, and electrochemical corrosion resistance. Although a certain degree of progress has been made in LC technology for the in situ synthesis of ceramic-reinforced TMCCs, further breakthroughs still require addressing the following core challenges and exploring more visionary research directions.

(1) The development of in situ-synthesized ceramic-reinforced TMCCs via LC lags behind that of laser additive manufacturing for in situ-synthesized ceramic-reinforced TMCs. Coating technology should draw insights from material forming techniques in terms of the processing methods, parameter optimization, compositional design, and performance enhancement.

(2) Some promising combinations of reinforcement phases have not received sufficient research attention and in-depth investigation, warranting further investigation. Additionally, exploring the synergistic effects of multicomponent ceramic phases with rare earth elements and intermetallic compounds could enrich the ceramic reinforcement systems.

(3) Currently, most studies focus on conventional material properties such as the hardness, tensile strength, and wear resistance under normal environmental conditions. There is a lack of research on the corrosion resistance, biocompatibility, high-temperature performance, and extreme environmental performance (such as complex biofluid conditions, extreme temperatures, high-pressure settings, and high-heat flux conditions).

(4) The optimization of process parameters and compositional design still rely heavily on trial-and-error methods and traditional parameter optimization approaches, making dynamic control challenging. Future efforts should integrate artificial intelligence (AI) and machine learning to construct real-time predictive models linking the process parameters, microstructure, and performance. For example, digital twin technology could be employed to simulate molten pool dynamics, phase transformations, and stress evolution, enabling the precise control of high-performance products.

(5) The evaluation of the weightage of different strengthening mechanisms remains incomplete. Researchers are encouraged to conduct further in-depth investigations integrating materials science and interface engineering to elucidate strengthening mechanisms and establish an evaluation theory, thereby guiding the design of high-performance materials.

(6) The field of coatings currently lacks sufficient interdisciplinary integration. To advance the high-performance material design and sustainable development, cross-disciplinary collaboration must be strengthened through initiatives such as applying the Materials Genome Initiative (MGI) to accelerate coatings material design, combining environmental science to develop green coatings compliant with international environmental regulations, and implementing geometric sensing technology for the real-time monitoring of orthopedic implants like artificial joints.