The Effect of CeO₂ Content on the Microstructure and Properties of TiC/WC/Co Composite Cladding Layers

Abstract

To address the issue that the insufficient surface hardness and wear resistance of ductile iron under harsh working conditions are likely to lead to early failure, using a cladding layer with dual hard phases is an effective method to improve the surface properties. However, the issue that a large amount of hard phases decompose under the action of a high-energy laser to generate brittle phases in the microstructure is quite troublesome. Therefore, by adding CeO₂ to the cladding layer, a TiC/WC/Co composite cladding layer containing CeO₂ is prepared on the substrate by means of a fiber laser. Through OM, SEM-EDS, XRD, and Rockwell hardness tests, the effects of the CeO₂ content on the microstructure, phase composition, and hardness of the coating were studied to determine the optimal addition amount. The results show that the secondary dendrite arm spacing (SDAS) of the γ-Co phase and the sizes of TiWC₂ and WC dendrites exhibit a non-monotonic trend of first decreasing and then increasing with the increase in the CeO₂ content, and the morphology of TiWC₂ evolves from a cross shape to a granular shape and then to a dendritic shape. When the CeO₂ content is 2 wt.%, the WC dendrites are completely inhibited, and the SDAS of γ-Co reaches the minimum value; when the content increases to 4 wt.%, WC dendrite coarsening occurs, and at the same time, the γ-Co dendrite packing density increases significantly, and the eutectic fraction decreases obviously. The hardness of the coating first increases and then decreases with the increase in the CeO₂ addition amount, and reaches a peak value of 91.4 HRC when the CeO₂ content is 4 wt.%, which is approximately 2.57 times the hardness of the substrate.

1. Introduction

Ductile iron (DI) has become a cornerstone material in heavy-duty engineering applications, including elevator traction wheel, gearbox housings, and hydraulic components, owing to its exceptional combination of tensile strength exceeding 600 MPa (ASTM A536), impact toughness of 12–15 J (Charpy V-notch), and cost-effectiveness [1]. Nevertheless, its limited surface hardness renders its surface vulnerable to abrasive wear and surface-initiated fatigue failure under cyclic loading conditions [2]. Surface degradation often precipitates catastrophic component failure through crack propagation. This inherent limitation has driven intensive research into surface engineering solutions, with particular emphasis on enhancing near-surface hardness [3].

Laser cladding has emerged as a paradigm-shifting surface modification technology, distinguished by its precision energy deposition and ultra-rapid solidification (cooling rates up to 106 K/s) capabilities [4]. The process involves the laser-induced formation of metastable phases and metallurgical bonds with the substrate surface, producing coatings with a small heat-affected zone, minimal deformation, rapid cooling, dense microstructure, low dilution rate, and a wide range of applicable materials [5,6].

Contemporary laser cladding technologies predominantly employ three self-fluxing alloy systems: ferrous-based (Fe-Cr-B-Si), nickel-based (Ni-Cr-B-Si-Mo), and cobalt-based (Co-Cr-W-C) formulations, each demonstrating distinct performance matrices [7]. Fe-based variants (e.g., Fe55 alloys), while economically advantageous with raw material costs 40%–60% lower than alternatives, exhibit inherent limitations including restricted self-fluxing capacity (wetting angle q > 110°), elevated crack susceptibility, and porosity rates exceeding 12% in as-clad conditions [8]. Ni-based systems (typified by IN625 derivatives) achieve good self-fluxing capacity, delivering enhanced oxidation resistance, heat resistance, and corrosion stability, but the wear resistance is relatively poor [9]. Co-based alloy systems (e.g., Stellite series) demonstrate exceptional high-temperature stability, maintaining 0.2% yield strength (s0.2) above 450 MPa at 800 °C, coupled with outstanding corrosion resistance and wear resistance [10]. Their superior molten pool wettability stems from favorable surface tension characteristics and a strategically engineered melting hierarchy, where the cobalt matrix (Tm » 1495 °C) liquefies prior to carbide phases (Tm > 1600 °C). Additionally, Co-based alloys have a compatible thermal expansion coefficient with ductile iron substrates.

The incorporation of hard ceramic phases has emerged as a pivotal strategy for enhancing coating hardness and wear resistance. Extensive studies have demonstrated that transition metal carbides, particularly titanium carbide (TiC) (Knoop hardness HK = 24–31 GPa) and tungsten carbide (WC) (HK = 18–22 GPa), exhibit exceptional interfacial compatibility with ferrous matrices, showing contact angles below 35° with iron, cobalt, and chromium at 1450 °C [11,12]. This superior wettability, coupled with their intrinsic hardness and thermal stability (decomposition temperature >2600 °C), renders them ideal reinforcing constituents for advanced wear- and corrosion-resistant coatings.

Recent advancements have focused on synergistic reinforcement through multi-phase ceramic systems. Zhu et al. [13] systematically investigated TiC/WC composite coatings on 316 L stainless steel, establishing a quantitative relationship between ceramic content (5–20 wt.%) and grain refinement. Their findings revealed that a 15% WC—5% TiC formulation yielded optimal performance, achieving a 3.2-fold increase in microhardness compared to the substrate. Liu et al. [14] developed a novel Ni-WC-TiC coating system on Cr₁₂MoV die steel, demonstrating that TiC addition (5–10 wt.%) effectively refined WC grains through the Zener pinning effect, reducing the average carbide size from 3.8 mm to 1.2 mm. The optimized coating exhibited a maximum hardness of 770.7 HV0.5 and wear resistance improvement by an order of magnitude (specific wear rate: 2.1 × 10⁻⁵ mm³/N m vs. 2.4 × 10⁻⁴ mm³/N m for substrate). Shi et al. [15] pioneered a nano–micro dual-scale reinforcement approach in (nano-WC + micro-TiC)/Ti6Al4V coatings. Their experimental matrix revealed that 1.5 wt.% nano-WC addition generated the most favorable Hall–Petch strengthening effect, yielding a microhardness of 459.99 HV (35.99% enhancement over substrate) while maintaining adequate fracture toughness (KIC = 8.2 MPa m^1/2).

Previous studies have demonstrated the efficacy of rare-earth oxides (REOs) as multifunctional additives in laser cladding composite coatings. Rare-earth elements possess unique electronic structures that can purify alloy solutions and refine grains. The appropriate addition of REOs can effectively reduce the dilution rate, refine the grains, lower porosity, and alleviate residual thermal stress, thereby reducing the crack sensitivity of the coating [16]. This significantly prevents defects and markedly enhances the surface hardness and wear resistance of the cladding layer.

Systematic investigations have elucidated the concentration-dependent effects of nano-CeO₂ additions in various coating systems. Ding et al. [17] engineered Co-based composite coatings on SPHC steel, establishing a non-monotonic relationship between CeO₂ content (0–2.5 wt.%) and mechanical properties. The optimal formulation (1.5 wt.% CeO₂) yielded a 2.8-fold hardness increase (from 350 HV to 980 HV) and 85% reduction in wear rate (from 4.7 × 10⁻⁴ to 7.2 × 10⁻⁵ mm³/N m), attributed to enhanced grain refinement (average grain size reduction from 12.5 mm to 3.2 mm) and homogeneous Ce-rich precipitate distribution. Li et al. [18] optimized processing parameters for Co-based coatings on Q235 steel, demonstrating that 2.0 kW laser power combined with 1.5 wt.% CeO₂ generated optimal tribological performance (coefficient of friction m= 0.32, specific wear rate k = 1.8 × 10⁻⁵ mm³/N m). Excessive CeO₂ (2.0 wt.%) induced particle agglomeration, increasing the wear rate by 40%. Sui et al. [18] prepared MMC cermet coatings on a TiAl/WC/CeO₂/TC21 alloy substrate. The results show that CeO₂ did not alter the phases of the composite coating, yet the shape of the TiC phase is closely related to the CeO₂ content. CeO₂ enhanced the fluidity of the molten pool, thereby further promoting the TiC/Ti₂AlC core-shell strengthening stage. With the increase in CeO₂ content, the optimized coating is conducive to a uniform microstructure distribution and fine grain size. Due to the strengthening effect of the hard phase and the dispersion strengthening effect of CeO₂, the microhardness of the composite coating is significantly higher than that of the substrate (nearly 1.6 times). Importantly, the addition of CeO₂ significantly improves the wear resistance of the composite coating. Sang et al. [19] introduced CeO₂ into the Fe-based laser cladding coatings on the surface of MgAl₂O₄/GCr15 steel. The results showed that with the increase in CeO₂ content, the metal grain size decreased first and then increased, while the hardness and wear resistance of the MgAl₂O₄/Fe-based coatings increased first and then decreased. When the CeO₂ content was 1.0 wt%, the hardness increased by 10% and the wear rate decreased by 40%. Wu et al. [20] demonstrated that 1% CeO₂ in WC/Ni60 coatings on Q235 steel led to the refinement of the grains in the coating, the formation of carbides was reduced, and the coating exhibited the best uniformity, showing the lowest friction coefficient, the lowest wear rate, and the best corrosion resistance. Zhang et al. [21] designed several innovative mixed powders of Ti6Al4V and NiCr-Cr₃C₂ with different CeO₂ contents (0, 1, 2, 3, and 4 wt.%), and prepared Ti₂C-reinforced CrTi4-based composite coatings on the surface of Ti6Al4V through laser cladding technology. The results show that the amount of CeO₂ added has a significant influence on the forming quality of the composite coatings. When the CeO₂ content is 2 wt.%, the cracks are completely eliminated; in addition, the lowest porosity is obtained when 3 wt.% of CeO₂ is added. Moreover, the average hardness of the composite coatings is 1.28 to 1.34 times higher than that of the Ti6Al4V substrate. The wear resistance of the composite coatings is significantly higher than that of the substrate. Both the composite coatings and the Ti6Al4V substrate exhibit a mixed wear mode of abrasive wear and adhesive wear.

Despite extensive research on single-phase reinforced coatings, the synergistic effects of CeO₂ addition in dual-hard-phase (TiC/WC) composite systems remain largely unexplored. This study addresses this knowledge gap by systematically investigating the influence of CeO₂ concentration (0–7.0 wt.%) on the microstructural evolution and mechanical properties of laser-clad TiC/WC/Co composite coatings on ductile iron substrates. Building upon established single-phase reinforcement mechanisms [20], we employ advanced characterization techniques to quantify CeO₂’s role in microstructural change, establish composition–property relationships for hardness optimization, and identify the critical CeO₂ concentration threshold for dual-phase reinforcement. This comprehensive approach aims to determine the optimal CeO₂ addition for maximizing coating performance while maintaining interfacial integrity with the ductile iron substrate.

2. Materials and Methods

2.1. Materials

The laser cladding experiment employed QT600-3 ductile iron as the substrate which measured 150 mm × 150 mm × 10 mm (length·width·thickness).

Table 1. Chemical composition of ductile iron (wt.%).

Element	C	Cu	Si	Ni	Mn	S	P	Fe
wt.%	3.33	0.32	2.71	-	0.42	0.012	0.036	Bal.

presents the chemical composition of the substrate material. QT600-3 was selected as the substrate material due to its excellent combination of strength (Rm³ 600 MPa) and ductility (elongation³ 3%), which is representative of cast iron components commonly used in industrial applications. The cobalt-based Co6 alloy powder (particle size range: 44–125 mm) served as the cladding matrix material, with its chemical constituents detailed in

Table 2. Chemical composition of cobalt-based alloy powder (wt.%).

Element	C	Cr	Si	W	Fe	Mo	Ni	Mn	B	Co
wt.%	1.21	29.56	1.26	5.09	2.62	0.38	2.59	0.22	-	Bal.

. The Co6 alloy system was chosen for its inherent self-fluxing characteristics and excellent wettability with ferrous substrates, while the selected carbide reinforcements (TiC/WC) provide synergistic strengthening through solid-solution hardening and grain refinement mechanisms. To enhance the coating performance, micron-scale ceramic reinforcements were incorporated through the addition of TiC and WC powders (99.0% purity, 40–60 mm particle size), along with rare-earth oxide cerium dioxide (CeO₂, 99.9% purity, 1 mm average particle size). The composite powder formulation maintained a volume ratio of 7:3:3 between the Co6 matrix alloy, TiC, and WC, respectively. Building upon the composite powder system, cerium dioxide (CeO₂) additions were introduced at varying mass fractions of 0 wt.%, 2 wt.%, 4 wt.%, and 7 wt.% to investigate rare-earth modification effects, as seen from

Table 3. Chemical composition of laser cladding powders (wt.%).

Sample No.	CeO2	Co6+TiC+WC (7:3:3)
S1	0	100
S2	2	98
S3	4	96
S4	7	93

.

Mechanical homogenization was achieved through planetary ball milling (300 rpm, ball-to-powder ratio 4:1) with an optimized blending duration of 10 min, followed by vacuum drying at 120 °C for 2 h (10⁻² Pa atmosphere) to remove absorbed moisture and minimize oxide formation. The resultant powder batches were systematically designated as S1–S4, corresponding to ascending CeO₂ concentrations. Figure 1 and Figure 2 represent the morphological evolution of the powder system through SEM-EDS analysis, where point A represents Co based powder, point B is TiC powder, point C is WC powder, and point D corresponds to CeO₂ powder, while Table 3 provides a quantitative breakdown of the precisely controlled ratios implemented in the laser cladding trials.

Prior to the cladding experiment, the substrate surface was meticulously prepared through a series of treatments. First, it was ground using 400-grit SiC paper. The substrate was then ultrasonically cleaned in analytical-grade ethanol (99.9% purity) for 10 min at a frequency of 40 kHz and power of 300 W to remove any contaminants. After cleaning, it was air-dried at 80 °C for 10 min to ensure the complete evaporation of residual solvents. Finally, the substrate was preheated to 300 ± 10 °C in a controlled environment to minimize thermal stress during the cladding process. Following preheating, the mixed powder was evenly spread onto its surface using a powder spreading method to ensure consistent powder distribution. The laser cladding experiment was conducted using a IPG YLS-6000 fiber laser system, featuring a wavelength of 1060 nm and beam quality factor M² < 1.1, which deposited a rectangular cladding layer measuring 15 mm·20 mm on the substrate surface. The schematic diagram of the laser cladding process is illustrated in Figure 3.

The laser power was maintained at 1.2 kW, corresponding to a power density of 17 kW/cm², while the beam diameter was set to 3 mm with a Gaussian intensity distribution. The scanning speed was controlled at 5 mm/s with a 50% overlap ratio between adjacent tracks, resulting in a consistent layer thickness of 1.2 mm. These parameters were carefully selected to achieve a balance between coating quality and processing efficiency. Throughout the process, shielding gas was supplied at a flow rate of 15 L/min to maintain an inert atmosphere.

2.2. Methods

The microstructure of the coating was analyzed using an Olympus metallographic microscope and a scanning electron microscope (SEM). Additionally, the phase composition of the coating surface was analyzed using X-ray diffraction (XRD).

Following the laser cladding process, the sample blocks were precisely sectioned into 11 mm’18 mm’10 mm specimens using a precision wire-cutting machine, with the cutting orientation maintained perpendicular to the cladding surface to ensure consistent cross-sectional analysis. The prepared specimens underwent a sequential surface treatment process: initial grinding with progressively finer SiC papers (up to 2000 grit), followed by diamond suspension polishing, and final etching using a mixed acid solution of HF, HNO₃, and H₂O for 30–35 s to reveal the microstructural features. After etching, the specimens were thoroughly rinsed with anhydrous ethanol and dried using a compressed air blower to prevent water marks or surface contamination.

Microstructural characterization was performed using an Olympus GX71 inverted metallographic microscope (manufactured by Olympus Corporation, Tokyo, Japan) equipped with digital image capture capabilities, complemented by detailed scanning electron microscopy (SEM) analysis using a Zeiss-Sigma field emission microscope operating (manufactured by Carl Zeiss AG, Oberkochen, Germany) at 15 kV. Phase identification was conducted through an X-ray diffraction (XRD) analysis using a Rigaku SmartLab diffractometer (manufactured by Rigaku Corporation, Tokyo, Japan) with Cu Ka radiation (l = 1.5406 Å), employing a continuous scan mode with a step size of 0.02° and a scanning rate of 2°/min over a 2q range of 20–90°.

2.3. Hardness Measurement

The microhardness profile across the coating cross-section was systematically evaluated using an HR-150A Rockwell hardness tester under a controlled load of 150 kg with a dwell time of 15 s. To ensure comprehensive characterization, five equidistant measurement points were strategically selected along a linear path extending from the substrate through the interfacial zone to the cladding surface. At each designated location, three independent measurements were performed to account for potential microstructural heterogeneity, and the arithmetic mean of these values was calculated to represent the local hardness. This rigorous measurement protocol was implemented to generate a reliable hardness distribution profile that accurately reflects the mechanical property gradient across the different regions of the coated system.

3. Results

3.1. XRD Analysis of Co-Based Composite Cladding Layers

The X-ray diffraction patterns presented in Figure 4 compare the phase compositions of composite coatings with 0 wt.% and 7 wt.% CeO₂ additions. Comprehensive phase analysis demonstrates that the fundamental phase constituents remain consistent across different CeO₂ concentrations, comprising γ-Co solid solution, Fe₃C, TiWC₂, Cr₂₃C₆, WC, and various Co-Cr-W-C composite carbides. Interestingly, despite systematic CeO₂ incorporation, no cerium-containing phases were detected within the sensitivity limits of XRD (typically > 1–2 wt.%), suggesting either the complete dissolution of Ce into the matrix or formation of nanoscale precipitates below the detection threshold.

The quantitative analysis of the diffraction patterns reveals significant compositional evolution with increasing CeO₂ content, and that beneficial alloying elements (Co, Cr, Ni, Si) exhibit concentration increases and detrimental elements, particularly Fe, show a marked reduction. The Fe₃C phase demonstrates progressive suppression, evidenced by the reduction in characteristic diffraction peak intensity at 7 wt.% CeO₂ addition, which contributes to improved coating performance. The absence of detectable Ce-containing phases, despite significant property enhancements, suggests that CeO₂ primarily functions through grain boundary segregation and catalytic effects during solidification rather than discrete phase formation.

3.2. Microstructural Characteristics of Composite Cladding Layers

The microstructural evolution of the TiC/WC/Co composite coatings is systematically presented in Figure 5 and Figure 6. Figure 4 specifically illustrates the bonding zone characteristics comparing specimens without CeO₂ addition and with 4 wt.% CeO₂ addition, while Figure 6 provides a comprehensive view of the microstructural morphology across different regions under a constant laser power of 1.2 kW. The micrographs in Figure 5 are organized to facilitate comparative analysis: (a,b), (c,d), (e,f), and (g,h) correspond to the middle and near-surface regions of coatings with CeO₂ contents of 0 wt.%, 2 wt.%, 4 wt.%, and 7 wt.%, respectively.

The detailed microstructural analysis reveals a distinct gradient evolution from the bonding zone to the near-surface region: planar dendrites → cellular dendrites → fine dendritic dendrites → equiaxed dendrites, as seen from Figure 5 and Figure 6. This morphological transition is fundamentally governed by the rapid solidification dynamics inherent to laser cladding, characterized by the solidification rates (R) and steep temperature gradients (G) (10⁵–10⁷ K/m) and the growth direction of the dendrites follows the direction of alloy latent heat dissipation. The solidification process initiates with heterogeneous nucleation at the substrate–coating interface, where the thermal conditions (G→∞, R→0) favor planar crystal formation [21]. As solidification progresses, the decreasing temperature gradient (G) and increasing solidification rate (R) lead to interface instability, promoting the transition to cellular and subsequently dendritic structures. The development of prominent primary and secondary dendrite arms reflects the expansion of the constitutional undercooling zone. In the near-surface region, thermal conditions (low G, high R) facilitate equiaxed grain formation through the homogenization of the molten pool temperature distribution, the expansion of the constitutional undercooling zone, and the free growth and proliferation of grains.

The microstructural evolution is quantitatively described by the G/R ratio, which serves as the primary determinant of solidification morphology, influencing crystal growth patterns, secondary dendrite arm spacing (SDAS), and grain size [22]. Complementary factors including Marangoni convection, thermal distribution within the molten pool, and alloy element segregation further modulate the final microstructure [23]. The resultant gradient structure significantly impacts the mechanical properties of the coating.

As evidenced in Figure 6b, in the absence of CeO₂ addition, the microstructure is characterized by substantially coarse TiC hard phases, with an average particle size ranging from 2 to 3 mm, which exhibit pronounced segregation along dendritic boundaries. The introduction of CeO₂ significantly alters the morphological evolution and spatial distribution of TiC particles in the composite coating. As the CeO₂ content increases from 2 wt.% to 4 wt.%, a progressive refinement of TiC particles is observed, with the average particle size decreasing from 1–2 mm while simultaneously achieving improved distribution homogeneity and increased particle number density. The optimal refinement effect occurs at 4 wt.% CeO₂ addition, where TiC particles reach their minimum average size of 0.5–1 mm and demonstrate uniform dispersion within the matrix, as clearly evidenced in Figure 6e.

However, this beneficial refinement effect exhibits a non-monotonic relationship with CeO₂ content. When the CeO₂ concentration increases to 7 wt.%, the system shows contrasting behavior characterized by three distinct phenomena: that the population of unmelted TiC particles reaches its maximum, the particle size distribution broadens significantly (0.5–2.5 mm), and the particles develop a disordered aggregation pattern along dendritic boundaries, as illustrated in Figure 6g. This transition suggests the existence of an optimal CeO₂ concentration window (2–4 wt.%) for microstructural refinement, beyond which excessive addition may lead to particle agglomeration and distribution heterogeneity.

The microstructural evolution demonstrates a clear dependence on CeO₂ content, exhibiting distinct morphological transformations. The baseline microstructure without CeO₂ addition (Figure 7a) reveals four distinct microstructural constituents with characteristic morphologies and compositions. Region A exhibits characteristic WC dendritic structures (by EDS analysis in Figure 8a), featuring coarse primary dendrites (5–8 mm arm spacing) with pronounced petal-like secondary branches. Adjacent to the WC dendrites, Region B displays a cross-shaped microstructure that, identified by EDS analysis in Figure 8b, contains Ti, W, and C elements in a near-equiatomic ratio, confirming the presence of TiWC₂ phase combining with XRD results in Figure 4. Region C presents a dendritic morphology rich in Co, Cr, and W elements (Figure 8c), consistent with γ-Co solid solution formation. The remaining Region D appears as featureless black areas in SEM images, with EDS quantification (Figure 8d) revealing characteristics of Fe-Co intermetallic compounds.

The introduction of 2 wt.% CeO₂, while not altering the fundamental phase composition, induces significant modifications in microstructural features, as shown in Figure 7b. The most notable changes include a substantial refinement of dendritic structures, with dendrite arm spacing reduced by 30%–50%, accompanied by improved spatial distribution homogeneity and a 120%–150% increase in dendrite density per unit area. Compared to the characteristic cross-shaped TiWC2 observed in the absence of CeO₂, the introduction of CeO₂ promotes a progressive refinement of dendritic structures, marked by a transition from granular to well-defined dendritic morphologies. Complete suppression of WC dendritic formation occurs. Concurrently, the secondary dendrite arm spacing (SDAS) of γ-Co dendrites decreases substantially from 2.58 mm to 1.53 mm, as presented in Figure 9, accompanied by improved spatial distribution uniformity.

Further increasing CeO₂ content to 4 wt.% induces partial microstructural coarsening. The Fe-Co intermetallic phase content diminishes due to reduced Fe availability, and the morphology of WC changes from coarse. TiWC₂ develops a strip-like morphology, and γ-Co dendrites exhibit intermediate coarsening with SDAS increasing to 1.87 mm, a value between those observed at 0 wt.% and 2 wt.% CeO₂.

The progressive depletion of Fe-Co intermetallic compounds (Region D) directly correlates with the enhanced precipitation of W and Cr elements, which preferentially participate in the formation of WC, TiWC₂, and γ-Co dendritic phases. However, when the CeO₂ content reaches 7 wt.%, the TiWC₂ dendrites lose their characteristic alignment and develop a randomized spatial distribution, while simultaneously the γ-Co dendritic structures undergo secondary refinement. This paradoxical behavior—combining disordered TiWC₂ arrangement with enhanced γ-Co refinement—is accompanied by a marked increase in the intermetallic phase content, as clearly evidenced in Figure 7d.

These microstructural changes, particularly the non-monotonic variation in secondary dendrite arm spacing as illustrated in Figure 8, suggest the existence of an optimal CeO₂ concentration window (2–4 wt.%) for achieving desirable microstructural characteristics in the cladding layers. The observed morphological evolution underscores the complex interplay between CeO₂ content and solidification behavior during cladding formation. Although the addition of CeO₂ can improve the fluidity of the molten metal and reduce the melting point and surface tension, excessive addition leads to the formation of more agglomerates in the coating, affecting the uniformity of the microstructure [24].

3.3. Hardness and Analysis

Figure 10a presents the hardness profiles across the cross-section of laser-clad coatings with varying rare-earth oxide contents. The hardness distribution follows a characteristic pattern for all the compositions: beginning with substrate-equivalent hardness values, the curves demonstrate a gradual increase through the bonding zone, exhibit an abrupt change in the bonding zone, and ultimately stabilize at maximum values in the coating surface layer. The detailed analysis of these profiles reveals that rare-earth oxide additions significantly influence coating hardness. The optimal CeO₂ concentration of 4 wt.% produces a peak hardness of 91.4 HRC, representing a 15% enhancement compared to the rare-earth-free coating (79.5 HRC) and a 157% improvement over the substrate material (35.6 HRC).

Figure 10b presents the relationship between CeO₂ content and the average hardness of the cladding layers. The data reveal a non-monotonic dependence, with the maximum hardness (87.34 HRC) achieved at 4 wt.% CeO₂ content, representing a 17% enhancement compared to the rare-earth-free coating (74.56 HRC). This optimal hardening effect correlates directly with the microstructural characteristics observed in Figure 6c, which demonstrates superior uniformity and density at this composition.

4. Discussion

The laser cladding process subjects the TiC, WC, CeO₂, and Co alloy powder mixture to intense laser energy, triggering rapid heating and melting. Of these constituents, CeO₂ exhibits unique thermal behavior due to its exceptionally small particle size (~1 mm) compared to other powders. This nanoscale dimension endows CeO₂ with substantially increased specific surface area and surface energy, while its relatively low melting point promotes early-stage decomposition during laser processing. Under the high-temperature conditions (exceeding 2000 °C) generated by the laser, CeO₂ preferentially dissociates into atomic Ce and O [25], preceding the melting of other components in the powder mixture. This decomposition occurs within an extremely brief temporal window (10–50 ms) of laser interaction, with the liberated Ce atoms subsequently participating in various metallurgical processes that significantly influence the final microstructure. The rapid thermal decomposition of CeO₂ and the active role of its byproducts fundamentally differentiate its behavior from other powder constituents during laser cladding.

CeO₂ → [Ce] + 2[O],

(1)

The absence of CeO₂ in the cladding powder results in significantly elevated melting temperatures during laser processing due to the exceptionally high melting point of TiC (3140 °C). This creates a molten pool with substantially higher overall temperatures, causing almost the complete melting of TiC and WC (melting point 2750 °C) under the intense laser irradiation, while some very small particles of TiC and WC are always left in the molten pool. Some of them promote the nucleation and growth of TiC and WC during the subsequent solidification after laser passes (as indicated by A and B in Figure 7a); others directly become TiC/WC particles in the microstructure, as presented in Figure 6a,b.

However, the addition of 2 wt.% CeO₂ fundamentally modifies the system’s thermodynamic behavior during laser cladding. Under laser irradiation, while most CeO₂ decomposes, a fraction remains intact due to shielding by surrounding TiC or WC particles that attenuate direct laser energy absorption. The preferential decomposition of CeO₂ at elevated temperatures releases highly reactive Ce atoms into the molten pool [26]. These Ce atoms exhibit selective segregation around TiC particles due to the presented mutual solubility characteristics with TiC, as shown in Figure 11. This localized Ce enrichment induces two significant effects: first, it substantially reduces the melting temperature of TiC through chemical interaction; second, it enhances TiC solubility in the molten matrix. These combined effects lead to a marked decrease in the overall molten pool temperature and a size decrease of TiC particles. WC particles completely melt also due to the earlier melting of TiC, despite lowering the molten pool temperature.

Both the decomposed Ce species and preserved CeO₂ particles contribute synergistically to microstructural refinement through distinct yet complementary mechanisms. The undissolved CeO₂ particles actively promote phase nucleation, while the liberated Ce modifies solidification kinetics and stabilizes grain boundaries throughout the process. Chen’s research [27] confirms that residual CeO₂ particles act as effective heterogeneous nucleation sites for TiC, significantly enhancing grain refinement. Concurrently, the surface-active nature of Ce reduces molten pool surface tension [28], lowering nucleation energy barriers and enabling the formation of smaller critical nuclei. Ce exerts equivalent effects on the nucleation processes of both γ-Co dendrites and TiC dendrites. As a surface-active element, Ce preferentially adsorbs onto the surfaces of γ-Co or TiC crystal nuclei, where it functions as a surface poison during crystal growth. This dual mechanism operates through two complementary pathways: First, Ce atoms uniformly promote nucleation across different crystalline phases by reducing interfacial energy barriers. Second, their selective surface adsorption modifies growth kinetics by altering atomic attachment energy at advancing solid–liquid interfaces [29]. The surface poisoning effect manifests through the partial blocking of active growth sites, leading to controlled dendrite refinement. This unique combination of enhanced nucleation and constrained growth enables Ce to simultaneously regulate the microstructure evolution of multiple phases in the cladding system. It is noteworthy that Ce atoms preferentially segregate to the surfaces of TiC particles, thereby inhibiting the attachment and growth of Ti and C atoms on these surfaces. This surface blocking effect prevents the TiC particles from serving as effective substrates for TiC crystal growth, resulting in their preservation as unmelted TiC particulate phases in the final microstructure.

The formation of WC grains is significantly influenced by the competitive phase evolution during solidification, driven by the structural and thermodynamic similarities between WC and TiC. Both the phases share a face-centered cubic crystal structure [30], with TiC exhibiting higher thermal stability and consequently nucleating preferentially during solidification. This preferential nucleation leads to the incorporation of W atoms into the TiC lattice, forming a (Ti,W)C solid solution that effectively depletes the available W content in the molten pool [31]. The microstructural evidence of TiWC₂ phase formation (Figure 3) further demonstrates the strong tendency for TiC to react with W and C atoms during solidification, creating an additional sink for W atoms [32]. These sequential reactions establish a compositional threshold that prevents the simultaneous accumulation of sufficient W and C concentrations required for WC nucleation. Consequently, despite the thermodynamic possibility of WC formation, the kinetic competition from (Ti,W)C and TiWC₂ phase development effectively suppresses WC grain formation under these conditions.

The addition of 4 wt.% CeO₂ leads to a distinct microstructural evolution, as revealed by elemental surface distribution analysis. Figure 11 clearly demonstrates significant Ce enrichment on TiC dendrites, indicating a strong chemical interaction between Ce and Ti atoms. This preferential segregation suggests the formation of stable Ce-Ti atomic clusters, reflecting their inherent chemical affinity. The phase diagram (Figure 12) further supports this observation, showing good mutual solubility between Ce and Ti at elevated temperatures, primarily attributed to their compatible crystal structures. This structural compatibility ensures high lattice matching, facilitating the incorporation of Ce atoms into Ti lattice sites. Moreover, the relatively large atomic radius of Ce plays a crucial role in this process. So, with the addition of CeO₂ increase from 2 wt.% to 4 wt.%, the Ce-Ti atomic clusters formed during solidification readily exceed the critical nucleus size for TiC nucleation, producing two significant effects: first, the nucleation incubation period for TiC is substantially reduced; second, the growth duration of existing nuclei is extended. These combined effects result in pronounced dendritic coarsening and a corresponding reduction in eutectic formation, ultimately altering the final microstructure of the cladding layer. Furthermore, the nucleation and growth of TiC dendrites consume a substantial amount of Ce atoms, consequently diminishing the grain-refining effect of Ce on γ-Co dendrites. Simultaneously, the latent heat released during TiC dendritic growth significantly reduces the temperature gradient within the molten pool. And the marked decrease in solute content in the melt elevates the liquidus temperature of the remaining alloy melt. These interrelated factors collectively enhance the constitutional undercooling at the solid–liquid interface front, thereby promoting the coarsening of γ-Co dendrites.

With the CeO₂ content increased to 4 wt.%, the particle size of TiC further decreases, while the molten pool temperature experiences additional reduction. Under these modified thermal conditions, the high-melting-point WC particles demonstrate incomplete dissolution in the molten pool at lower temperature than that of CeO₂ content of 2 wt.%, resulting in the preservation of numerous fine WC particulates. During solidification, the dissolved W and C atoms preferentially deposit onto these residual WC particles through direct epitaxial growth, leading to the reappearance of WC dendritic structures in the metallographic microstructure of the clad layer. Moreover, the addition of CeO₂ leads to a reduction in molten pool temperature, which consequently decreases substrate melting and significantly lowers the dilution rate of the cladding layer. As a result, the iron content within the microstructure is markedly reduced, accompanied by a corresponding decrease in Fe₃C in the eutectic structure.

When 7 wt.% CeO₂ is added, a significant number of Ce atoms are introduced into the molten pool. Due to their surface-active nature, these Ce atoms tend to accumulate on the surfaces of nucleating and growing crystals, forming a dense layer that induces surface poisoning. This layer hinders the adsorption of atoms necessary for crystal growth, effectively suppressing grain growth and leading to substantial grain refinement [33]. However, the inhibition of primary crystalline phase growth results in a notable increase in the eutectic structure, which enhances brittleness and reduces the toughness of the cladding layer. Furthermore, excess Ce elements tend to segregate at grain boundaries, promoting the formation of brittle phases and further deteriorating the material’s mechanical properties [34].

In summary, an appropriate CeO₂ content (e.g., 2 wt.%) optimizes the coating’s microstructure and mechanical properties. While higher CeO₂ additions (e.g., 4 wt.%) induce microstructural coarsening and the reappearance of WC dendrites, the increased volume of hard phases still enhances coating hardness significantly. At 7 wt.% CeO₂, pronounced grain refinement occurs; however, excessive Ce segregates at grain boundaries, forming brittle phases that degrade mechanical performance. These findings highlight the dual role of CeO₂ in regulating cladding layer properties, offering critical guidance for coating optimization.

5. Conclusions

In this study, by adding different contents of the rare-earth oxide CeO₂, the effects of CeO₂ on the microstructure and mechanical properties of the laser-cladded TiC/WC/Co composite coating were systematically investigated, and the following main conclusions were drawn:

• While the addition of CeO₂ preserved the original phase constituents (γ-Co, TiWC₂, WC, Cr₂₃C₆, Fe₃C, and Co-Cr-W-C) in the composite coating, it induced notable modifications in their morphological features, particularly in terms of grain refinement and spatial distribution;
• The secondary dendrite arm spacing (SDAS) of the γ-Co phase, the sizes of TiWC₂ dendrites and WC dendrites, as well as the fraction of iron-containing compounds all exhibit a non-monotonic trend, initially decreasing and then increasing with raising CeO₂ content. Notably, the morphology of TiWC₂ evolves from a cruciform shape to a granular shape, and finally to a dendritic shape. At 2 wt.% CeO₂, WC dendrites are entirely suppressed, and the SDAS of γ-Co reaches the minimum value. However, when the CeO₂ content increases to 4 wt.%, localized dendrite coarsening occurs, accompanied by a significant increase in dendrite packing density and a marked reduction in the eutectic fraction;
• During laser cladding, the introduced CeO₂ underwent thermal decomposition at elevated temperature, releasing Ce atoms that preferentially segregated around TiWC₂ phases and associated dendrites. These Ce atoms exerted a coordinated regulatory effect on multiple phases through two primary mechanisms: (1) grain refinement via heterogeneous nucleation and (2) the modification of molten pool surface tension characteristics. The Ce-mediated regulation simultaneously influenced the γ-Co matrix, TiWC₂ precipitates, and WC reinforcing phases, and even unmolten TiC particles dispersed in the coating;
• The coating hardness exhibited a distinct non-monotonic dependence on CeO₂ content, demonstrating an initial increase followed by subsequent decrease. The coating exhibited peak hardness performance at the optimal CeO₂ addition of 4 wt.%, with the hardness value reaching 91.37 HRC—representing a 1.57-fold enhancement relative to the substrate material.