Effects of Nitrogen Partial Pressure on the Microstructure and Mechanical Properties of High-Entropy Ti(C,N)-Based Gradient Cermets

Abstract

Titanium carbonitride (Ti(C,N))-based ceramics are widely utilized in mechanical machining, aerospace, and electronics, particularly in cutting tools and wear-resistant components. Two single-phase solid solution powders, non-high-entropy (Ti_0.83,W_0.07,Mo_0.04,Nb_0.03,Ta_0.04)(C_0.7,N_0.3) and high-entropy (Ti_0.6,W_0.1,Mo_0.1,Nb_0.1,Ta_0.1)(C_0.78,N_0.22), were synthesized via the carbothermal reduction–nitridation (CRN) method. Gradient-structured non-high-entropy (C-TiCN) and high-entropy (HE-TiCN) cermets were fabricated at 1450 °C by tailoring the nitrogen partial pressure in the range of 1–8 kPa. The effect of nitrogen partial pressure on the microstructure and mechanical properties of both materials was thoroughly analyzed. Both materials exhibited a three-layer gradient structure comprising a hard-phase-enriched surface layer, a binder-rich subsurface layer, and a chemically uniform core. Optimal performance was achieved at 4 kPa nitrogen partial pressure, at which both HE-TiCN and C-TiCN exhibited a desirable combination of surface hardness and fracture toughness. Compared with C-TiCN, HE-TiCN showed improvements in surface hardness and fracture toughness at subsurface and core regions (40 µm from the surface) by 4.9%, 11.2%, and 12.0%, respectively. The enhanced surface hardness of HE-TiCN is attributed to the significant lattice distortion and the synergistic effects associated with its high-entropy configuration. The improved toughness of the binder-rich layer is primarily ascribed to mechanisms such as crack deflection, crack branching, and the formation of tear ridges. These findings offer a promising strategy for developing gradient Ti(C,N)-based cermets with enhanced mechanical performance.

1. Introduction

Cermets are a class of composite materials that achieve optimized mechanical properties through the synergistic integration of ceramic and metallic phases. The ceramic phase imparts high hardness and thermal stability, while the metallic binder phase enhances fracture toughness and impact resistance [1,2]. Titanium carbonitride (Ti(C,N))-based cermets are advanced composite materials derived from WC-Co cemented carbides and TiC-based cermets. Typically fabricated through powder metallurgy techniques [3], these materials consist of ceramic hard phases (Ti(C,N)) and metallic binder phases, such as Co or Ni. These cermets exhibit a desirable combination of high hardness, wear resistance, oxidation resistance, and flexural strength, making them widely applicable in mechanical machining, aerospace, and electronics industries, with particularly outstanding performance in cutting tools and wear-resistant components [4,5,6]. Recent developments in Ti(C,N)-based cermets have primarily focused on enhancing their performance within Ti(C,N)-Co-Ni systems. This has involved the incorporation of additional ceramic phases or alloying elements to simultaneously improve hardness, wear resistance, and fracture toughness. For example, Zhu [7] et al. introduced Al₂O₃ into the Ti(C,N) matrix to fabricate Ti(C,N)/Al₂O₃ cermets, achieving a hardness of 18.44 ± 0.30 GPa, though at a cost of a relatively low fracture toughness of only 7.86 ± 0.10 MPa·m^1/2. Hu [8] et al. optimized the microstructure and mechanical behavior of layered Ti(C,N)-HfN/Ti(C,N)-WC ceramics by varying the layer thickness ratio. A ratio of 0.8 yielded a peak hardness of 16.93 ± 0.30 GPa, albeit with a decline in fracture toughness. Similarly, Gao [9] et al. reported that the addition of 2.5 at.% Re increased the Vickers hardness of Ti(C,N)-WC-HfN cermets to 19.25 ± 0.21 GPa, while simultaneously reducing both flexural strength and toughness. These findings underscore a persistent challenge in cermet development with regard to achieving the trade-off between hardness and fracture toughness. Simultaneously enhancing both properties remains a significant obstacle.

High-entropy alloys (HEAs) represent a novel alloy design concept initially proposed by Yeh [10] and Cantor [11]. Based on the configurational entropy (∆S_conf) of ideal solid solutions, early researchers categorized alloys as low-entropy (∆S_conf < 0.69R), medium-entropy (0.69R < ∆S_conf < 1.61R), and high-entropy (∆S_conf > 1.61R) systems [12]. Over the past decade, this concept has rapidly evolved from HEAs to high-entropy ceramics (HECs), extending to a broader range of composition types. With the expansion of the high-entropy concept into ceramics, extensive exploration has been conducted across various systems—including carbides, nitrides, and oxides—gradually establishing a comprehensive framework for the design, classification, and fabrication of high-entropy ceramics. Wang et al. [13] and Fang et al. [14] developed Ti(C,N)-based ceramics utilizing CoCrFeNiCu and Al₃CoCrFeNi high-entropy alloys as binders, respectively. Compared with conventional Co-Ni binders, these materials exhibited superior fracture toughness and hardness. Zhou et al. [15] successfully synthesized the first class of high-entropy solid solution carbides via carbothermal reduction and conducted an in-depth analysis of the synergistic interactions among constituent elements and the influence of lattice distortion on mechanical performance. Their findings confirm that the mechanisms of solid solution strengthening and sluggish diffusion in high-entropy ceramics significantly outperform those in conventional ceramic systems. Furthermore, Divilov et al. [2] proposed a novel theoretical framework describing high-entropy ceramic formation based on enthalpy–entropy synergy. They emphasized that the configurational entropy calculation must incorporate both cationic and anionic sublattices to accurately predict thermodynamic stability, thereby refining the theoretical methodology for HEC design.

Despite their outstanding hardness and thermal stability, high-entropy ceramics remain intrinsically brittle systems. In the absence of effective crack-control mechanisms, they frequently suffer from insufficient fracture toughness. To overcome this limitation, gradient microstructures in Ti(C,N)-based cermets have been explored as a potential solution. Zhao [16] et al. fabricated functionally graded Ti(C,N) cermets using low-pressure carburization, achieving a surface hardness of 1753 HV and a flexural strength of 2092 MPa. Liu [17] et al. employed nitrogen partial pressure sintering (130 Pa at 1450 °C) to produce ultrafine-grained Ti(C,N) cermets. A TiN-gradient surface layer and Ni-enriched layer were formed after 40 min of treatment, resulting in a hardness of 1635 HV and a flexural strength of 2350 MPa. However, excessive nitrogen partial pressure may lead to undesirable defects such as delamination and porosity. Lv [18] et al. reported abnormal grain growth and binder phase depletion in (Ti,Mo)(C,N) cermets when the nitrogen pressure exceeded 0.8 MPa, leading to increased porosity levels of 3–5%.

In previous work [19,20], high-entropy carbonitrides (HECNs) with the composition (Ti_0.6,W_0.1,Nb_0.1,Ta_0.1,Mo_0.1)(C_0.78,N_0.22)-Co-Ni were successfully synthesized via the carbothermal nitridation method, and the effects of Co/Ni binder content and sintering conditions on the performance of HECN-Co-Ni cermets were examined. To date, however, no studies have explored the role of nitrogen partial pressure in the formation of gradient structures within HECN-based cermets. Theoretically, the sluggish diffusion effect of high-entropy systems may suppress rapid atomic migration and reduce pore formation. Furthermore, the synergistic interactions among multiple principal elements—the so-called “cocktail effect”—may alleviate the inherent conflict between hardness and toughness. Solid solution strengthening induced by lattice distortion, coupled with stress buffering from gradient structures, has the potential to significantly enhance the mechanical performance of Ti(C,N)-based cermets [21].

This study aims to fabricate high-entropy gradient Ti(C,N)-based cermets (HE-TiCN) using (Ti_0.6,W_0.1,Nb_0.1,Ta_0.1,Mo_0.1)(C_0.78,N_0.22)-Co-Ni compositions and to systematically examine the effects of nitrogen partial pressure on the phase composition, microstructure, and mechanical properties. Comparative analysis with conventional, non-high-entropy Ti(C,N)-based cermets (C-TiCN) with the composition of (Ti_0.83,W_0.07,Mo_0.04,Nb_0.03,Ta_0.04)(C_0.78,N_0.22) will elucidate the structure–property relationships and establish both theoretical and practical foundations for the development of next-generation cermet materials.

2. Materials and Methods

To clearly illustrate the synthesis strategy and characterization pathway for high-entropy ceramic composites, the overall experimental procedure is summarized in the flowchart shown in Figure 1.

2.1. Synthesis of High-Entropy and Non-High-Entropy Hard Phase Precursors

The commonly accepted definition of high-entropy alloys involves multi-principal element systems comprising more than four major elements, each with a concentration between 5% and 35% [22]. In this study, the same definition was applied to high-entropy ceramics. The metal sources used included TiO₂ (purity ≥ 99.8%, particle size ≤ 0.1 µm), WO₃ (purity ≥ 99.9%, particle size ≤ 0.2 µm), MoO₃, Nb₂O₅, and Ta₂O₅ (all with purity ≥ 99.9% and a particle size of ≤1 µm). Carbon black (purity ≥ 99.9%, particle size ≤ 0.1 µm) was used as both a reducing agent and a carbon source. All raw materials were procured from Beijing Hongyu New Materials Technology Co., Ltd. (Beijing, China) or Shanghai Yantian New Materials Technology Co., Ltd. (Shanghai, China) Two sets of precursor mixtures were prepared based on the molar ratios listed in

Table 1. Composition of hard-phase powder materials.

Number	Elemental Ratio	Raw Materials (wt.%)
		C	TiO2	WO3	Nb2O5	Ta2O5	MoO3
HECN	(Ti0.6,W0.1,Mo0.1,Nb0.1,Ta0.1)(C0.78,N0.22)	23.42	30.36	14.69	8.42	14.00	9.12
CCN	(Ti0.83,W0.07,Mo0.04,Nb0.03,Ta0.04)(C0.78,N0.22)	25	50	12	4	4	5

.

The precursor powders were synthesized via a carbothermal reduction–nitridation (CRN) process. Specifically, the mixtures were heat-treated at 1450 °C for 2 h under a 2 kPa nitrogen atmosphere (purity ≥ 99.9%). After cooling, the resulting solid solution powders were labeled as HECN (high-entropy) and CCN (non-high-entropy), respectively. Both powders were confirmed to be single-phase solid solutions.

2.2. Ball Milling and Mixing

The synthesized HECN or CCN powders were blended with Co and Ni at a mass ratio of 85% HECN (or CCN)—7.5% Co—7.5% Ni. N-hexane was used as the milling medium. Wet planetary ball milling was performed in a tungsten carbide ball–stainless steel jar setup with a ball-to-powder ratio of 20:1. The milling was conducted at 250 rpm for 24 h to ensure uniform distribution and thorough mixing between the metallic binders and the ceramic hard phases.

Upon completion, the slurry was dried in a vacuum oven at 60 °C for 5 h. The dried powder was then sieved through a 125 µm mesh to obtain uniformly sized, well-dispersed composite powder. This sieved powder was loaded into a mold and compacted using uniaxial pressing at a pressure of 150 MPa.

2.3. Gradient Structure Sintering and Nitrogen Partial Pressure Control

The cold-pressed compacts were placed in a high-temperature vacuum sintering furnace to achieve densification and gradient structure formation. Sintering was conducted at 1450 °C with a holding time of 1 h. During the sintering process, high-purity nitrogen gas (N₂, purity ≥ 99.9%) was introduced at controlled partial pressures of 1, 2, 4, 6, and 8 kPa to regulate gradient evolution. This process yielded HE-TiCN and C-TiCN gradient cermets. After sintering, nitrogen flow was maintained during furnace cooling to room temperature, after which the specimens were retrieved for subsequent structural and mechanical characterization.

2.4. Structural and Mechanical Characterization Methods

Phase composition was analyzed using X-ray diffraction (XRD, Smartlab 3 KW, Rigaku Corporation, Tokyo, Japan) with a Cu target (Kα radiation, λ = 1.54 Å). The scanning range was 10–90°, with a step size of 0.02° and a scan rate of 1°/min. Microstructural morphology and elemental distributions were examined using a field-emission scanning electron microscope (FESEM, Sigma 500, Carl Zeiss AG, Oberkochen, Germany) equipped with energy-dispersive X-ray spectroscopy (EDS, X-MaxN, Oxford Instruments, Abingdon, UK), operated at an accelerating voltage of 15 kV and a working distance of 8 mm.

Microhardness was measured using a Vickers microhardness tester (HM-122, Akashi, Japan) under a load of 294 N and a dwell time of 10 s. For each specimen, twelve evenly distributed measurement points were selected on the surface or cross-section. After excluding invalid data caused by surface defects or unclear indentations, the average value of five valid measurements was recorded as the hardness. Fracture toughness (K_IC) was calculated using the following equation [23], and again, the average of five valid measurements was reported:

$$K_{IC}=0.15\times \sqrt{{\displaystyle \frac{{Hv}_{30}\times P}{4l}}}$$

where P is the applied load (N) and l is the length of the cracks emanating from the indentation corners (mm).

3. Results

3.1. Phase Composition

Figure 2 presents the X-ray diffraction (XRD) patterns of HE-TiCN and C-TiCN samples sintered under different nitrogen partial pressures. For both materials, only face-centered cubic (FCC) solid solution phases, denoted as (Ti,W,Mo,Nb,Ta)(C,N) phase, and Co–Ni binder phases were detected. Increasing the nitrogen partial pressure caused the diffraction peaks of the hard phase to shift toward higher angles; this reflected lattice contraction behavior. Notably, the intensity of the Co-Ni diffraction peaks did not decrease monotonically. Instead, a slight increase was observed as the nitrogen partial pressure rose from 1 kPa to 4 kPa, followed by a gradual decrease and eventual disappearance in the 4–8 kPa range. This trend may be attributed to the synergistic effect of nitrogen diffusion on the surface distribution and crystallinity of the binder phase. At lower nitrogen pressures, the crystallinity of the Co-Ni phase improved or underwent short-range enrichment, enhancing its diffraction response. At higher pressures, the binder phase tended to migrate inward, reducing its surface concentration and thereby diminishing diffraction intensity. Specifically, in HE-TiCN, the Co-Ni peaks at the surface nearly vanished at 6 kPa, whereas in C-TiCN, this occurred at 8 kPa, indicating that the high-entropy configuration promoted improved structural stability of the binder phase.

To verify the lattice distortion behavior, the 2θ values of the (200) diffraction peaks were extracted, and the corresponding lattice constants were calculated (as shown in

Table 2. The 2θ values of the (200) diffraction peak and corresponding lattice constants.

Pressure/kPa	2θ (200)/° (HE)	2θ (200)/° (C)	a/Å (HE)	a/Å (C)
1	41.36	41.26	4.362	4.373
2	41.5	41.3	4.348	4.369
4	41.58	41.52	4.341	4.346
6	41.62	41.7	4.336	4.328
8	41.7	41.73	4.328	4.325

). The results indicate that as the nitrogen partial pressure increased from 1 kPa to 8 kPa, the lattice constant of HE-TiCN gradually decreased from 4.362 Å to 4.328 Å, while that of C-TiCN decreased from 4.373 Å to 4.325 Å. Both materials exhibited typical lattice contraction behavior, which suggests that more nitrogen atoms were substituted for carbon atoms, thereby reducing the lattice constant. Comparatively, HE-TiCN showed a more gradual and stable lattice contraction, indicating that the high-entropy configuration facilitates the formation of a stable solid solution in the hard phase and suppresses localized phase instability—providing microstructural support for the balanced mechanical performance.

3.2. Microstructure

Figure 3 displays the cross-sectional microstructures and corresponding elemental distributions of HE-TiCN samples. As shown in Figure 3b,d,f, the surface layer was enriched in Ti and N, whereas Co and Ni were predominantly concentrated in the subsurface region. This distribution resulted in a well-defined three-layer gradient structure comprising (i) a surface layer rich in Ti and N, forming the FCC hard-phase; (ii) a subsurface layer enriched with Co and Ni; and (iii) a chemically homogeneous core.

As the nitrogen pressure increased from 1 kPa to 8 kPa, both the Ti,N-rich and Co,Ni-rich layers increased in thickness, although at different rates. At 1 kPa, a gradual compositional gradient was observed, with the hard-phase and binder layers measuring 7 µm and 6 µm, respectively. At 4 kPa, a more pronounced gradient developed, accompanied by a 57.1% increase in the hard-phase layer thickness, which increased to 11 µm, while the binder layer exhibited only a modest increase to 7 µm. The microstructure exhibited improved densification under moderate nitrogen pressures, characterized by well-defined phase boundaries and negligible porosity. However, at a higher nitrogen partial pressure of 8 kPa, although the hard-phase layer continued to grow and reached a thickness of 13 µm, the binder phase showed little additional expansion. Notably, localized porosity was observed at the ceramic–metal interface, which may have compromised the interfacial integrity and negatively impacted the mechanical performance.

Figure 4 illustrates the microstructure and elemental distribution of C-TiCN. Although a similar gradient structure was observed, the layer thicknesses were consistently lower than those of HE-TiCN. At 1 kPa, both the surface and subsurface layers measured approximately 4 µm. At 4 kPa, the hard-phase layer increased to 5 µm with only minor changes in the binder layer. At 8 kPa, the respective thicknesses reached 7 µm and 8 µm.

Core microstructures of both materials sintered at 8 kPa are shown in Figure 5. The core region of HE-TiCN exhibited a dense microstructure with a low porosity (<0.5 pores/100 μm²), consisting of fine, core-free hard-phase grains that were uniformly embedded in the binder. In contrast, the C-TiCN core contained numerous spherical pores (2–4 µm in diameter; 10 ± 3 pores/100 μm²) and distinct black core–gray shell features.

These differences can be attributed to the distinct diffusion behavior. In HE-TiCN, the increased configurational entropy induced substantial lattice distortion, increasing the activation energy for atomic diffusion [24,25]. This sluggish diffusion suppressed pore formation and enhanced densification [26,27]. Furthermore, the multicomponent elements in the high-entropy system reacted with nitrogen to compensate for vacancy flux, maintaining stoichiometric balance and minimizing pore nucleation. Additionally, the increased solubility of refractory elements in the Co-Ni binder lowered the nitrogen retention, thereby mitigating internal gas pressure buildup. By contrast, C-TiCN exhibited faster nitrogen diffusion due to lower entropy. As nitrogen diffuses outward under elevated pressure, internal vacancy accumulation leads to metastable pore formation [28]. During densification, residual nitrogen becomes entrapped, exceeding the local mechanical strength and resulting in gas-filled voids that degrade structural performance.

3.3. Mechanical Properties

Figure 6 presents the hardness and fracture toughness profiles of HE-TiCN and C-TiCN samples sintered under a nitrogen partial pressure of 4 kPa, measured from the surface to the core. As shown in Figure 6a,b, both materials exhibit a non-monotonic hardness distribution, characterized by an initial decrease followed by a subsequent increase with increasing depth. This trend reflects the underlying gradient architecture: the surface region, enriched in FCC hard phases, exhibits high hardness; the intermediate binder-rich layer, containing elevated concentrations of Co and Ni, shows reduced hardness; and the core maintains relatively stable hardness due to its uniform phase composition.

For HE-TiCN, the minimum hardness of 1750.7 HV was observed at a depth of 16 µm, corresponding to the binder-enriched region, representing a reduction of 239.7 HV compared to the surface. For C-TiCN, the minimum hardness was 1651.3 HV at 8 µm depth, representing a reduction of 156.7 HV from the surface.

Figure 6c,d display the evolution of fracture toughness with depth. In both materials, fracture toughness initially increased, reaching a maximum in the binder-rich region, followed by a decline and stabilization in the core. For HE-TiCN, the minimum fracture toughness was 9.4 MPa·m^1/2 at the surface, increasing to a peak of 11.9 MPa·m^1/2 at 16 µm depth, and stabilizing at 10.45 MPa·m^1/2 at 40 µm (core). In contrast, C-TiCN exhibited a peak toughness of 10.7 MPa·m^1/2 at 8 µm depth, with corresponding surface and core values of 8.4 and 9.33 MPa·m^1/2, respectively.

Figure 7 compares the peak hardness and fracture toughness values across different regions of both materials. Under 4 kPa nitrogen pressure, both HE-TiCN and C-TiCN achieved the desired “hard surface and tough core” configuration. Relative to C-TiCN, HE-TiCN demonstrated a 4.9% enhancement in surface hardness, alongside improvements of 11.2% and 12.0% in binder layer and core fracture toughness, respectively.

4. Discussion

4.1. Phase Transformation and Gradient Structure Formation

As shown in Figure 2, increasing the nitrogen partial pressure enhanced the intensity of Ti-N phases at the surface and led to the disappearance of the Co-Ni phase in both HE-TiCN and C-TiCN. This phenomenon was attributed to the strong chemical affinity between Ti and N. Under increased nitrogen chemical potential, Ti atoms diffuse rapidly through the Co-Ni binder toward the surface, where they form a dense FCC hard-phase layer. Simultaneously, the Co-Ni phase migrates toward the subsurface, generating a compositional gradient and eventually vanishing from the outermost surface region.

In addition, the limited solubility of nitrogen in Ni inhibits its inward diffusion [29], resulting in nitrogen enrichment near the surface under low nitrogen pressures. Compared with C-TiCN, HE-TiCN exhibits an earlier disappearance of the Co-Ni phase at a lower nitrogen pressure (4 kPa), which is attributed to the synergistic effects of multi-principal elements in the high-entropy system. These effects enhance the interaction between refractory elements (W, Ta, Mo, Nb) and nitrogen, accelerating the formation of multicomponent hard phases. As a result, the Co-Ni binder is more readily expelled to the subsurface [30,31], contributing to the formation of a thicker hard phase-enriched surface layer in HE-TiCN.

4.2. Mechanical Property Analysis

The mechanical performance of HE-TiCN and C-TiCN is closely linked to their gradient microstructures. The surface regions, enriched in hard phases, exhibit high hardness but reduced toughness due to their low binder content. Conversely, the subsurface layers, enriched in the Co-Ni binder, demonstrate enhanced fracture toughness. The gradual transition from a hard phase-dominated surface to a binder-rich interior enables both materials to simultaneously achieve a hard exterior and a tough core.

As shown in Figure 6 and Figure 7, both materials attain their maximum surface hardness at a nitrogen partial pressure of 4 kPa. HE-TiCN demonstrates superior surface hardness relative to C-TiCN, which is ascribed to its finer core–shell grain structure, higher densification, and pronounced lattice distortion associated with the high-entropy solid solution. Furthermore, as observed in Figure 3d and Figure 4d, the surface of HE-TiCN contains significantly higher concentrations of Ti and N than that of C-TiCN, reflecting the enhanced elemental interactions in high-entropy systems. This elemental synergy facilitates the migration of the Co-Ni phase to the subsurface, increasing the surface fraction of hard phases and thereby improving hardness.

Figure 8 compares the fracture surface morphologies and crack propagation behavior in the binder-rich layers of both materials under 4 kPa nitrogen pressure. In HE-TiCN, fracture occurs primarily via transgranular modes in coarse grains and intergranular modes in finer grains, accompanied by well-developed tear ridges along grain boundaries (Figure 8a). By contrast, C-TiCN predominantly exhibits intergranular fracture in fine grains, with limited transgranular features and less distinct tear ridges (Figure 8b). Figure 3d and Figure 4d also reveal a thicker binder-rich region in HE-TiCN, further supporting the conclusion that multi-principal-element interactions promote Co-Ni phase migration into the subsurface. This migration increases the average free path of the binder, enhancing its ability to undergo plastic deformation while fracturing. The resulting tear ridges dissipate stress and absorb energy, thereby showing improved fracture resistance [32].

As illustrated in Figure 8c,d, crack propagation in HE-TiCN involves greater branching and higher deflection angles, particularly at hard-phase grain boundaries. These characteristics increase the effective crack path and enhance resistance to crack growth, further contributing to the toughness of the binder-rich region [33]. The high fracture toughness observed in the binder-rich region of HE-TiCN is primarily attributed to synergistic toughening mechanisms such as crack deflection, crack branching, and the formation of tear ridges. These mechanisms effectively enhance crack resistance by increasing the crack propagation path, dispersing stress concentrations, and promoting energy dissipation. According to the findings of Ritchie et al. [34], such mechanisms are highly effective in improving the toughness of brittle and quasi-brittle materials, which fully supports the superior fracture behavior exhibited by the high-entropy system in this study.

Moreover, the superior microstructural stability of HE-TiCN is closely associated with its high configurational entropy. On one hand, the sluggish diffusion effect inherent to high-entropy structures suppresses excessive atomic migration during sintering, thereby reducing the risk of vacancy supersaturation and pore nucleation induced by nitrogen concentration gradients. On the other hand, the synergistic interactions among multiple principal elements promote uniform formation of the hard-phase skeleton, enhancing both the continuity and thermal stability of the gradient interfaces. As shown in Figure 5, under the same nitrogen pressure, the core microstructure of HE-TiCN is significantly superior to that of C-TiCN, exhibiting fewer pores, finer grains, and more coherent interfacial bonding. These findings confirm the beneficial effects of the high-entropy configuration in suppressing porosity and stabilizing the microstructure, which in turn account for the higher hardness and fracture toughness observed in HE-TiCN (Figure 7).

Compared with conventional Ti(C,N)-based cermets reported in the literature, the HE-TiCN and C-TiCN materials fabricated in this study achieved a more favorable synergy between hardness and fracture toughness. For instance, Hu et al. [8] reported a maximum hardness of 16.93 GPa in laminated Ti(C,N)-HfN/Ti(C,N)-WC ceramics; however, the toughness decreased as hardness increased. Zhao et al. [10] fabricated gradient-structured Ti(C,N) ceramics with a surface hardness of 1753 HV and a flexural strength of 2092 MPa via low-pressure carburization, but no significant improvement in fracture toughness was observed. In contrast, the HE-TiCN developed in this work achieved a surface hardness of 1989.7 HV while maintaining a maximum fracture toughness of 11.9 MPa·m^1/2. This demonstrates that a combination of high-entropy solid solution strengthening and gradient structuring effectively mitigates the common trade-off between hardness and toughness in conventional ceramic systems.

5. Conclusions

This study examined the effects of nitrogen partial pressure on the microstructure and mechanical properties of HE-TiCN and C-TiCN-based gradient cermets. The results demonstrated that nitrogen pressure critically influences phase composition, microstructural gradients, and mechanical properties. Both HE-TiCN and C-TiCN exhibited optimal mechanical performance when sintered at 1450 °C for 1 h under a nitrogen partial pressure of 4 kPa. The principal findings are summarized as follows:

(1) Both cermets developed a distinct three-layer gradient structure comprising a hard-phase-rich surface layer, a binder-enriched subsurface layer, and a chemically uniform core. With increasing nitrogen pressure, the thicknesses of the hard-phase and binder-phase layers increased. At nitrogen pressures below 8 kPa, HE-TiCN consistently exhibited thicker surface and subsurface layers compared to C-TiCN.

(2) At a nitrogen pressure of 4 kPa, both materials achieved a desirable combination of high surface hardness and improved interior toughness. The surface hardness of HE-TiCN reached 1989.7 HV. The fracture toughness values in the binder-rich region and at a depth of 40 µm from the surface were 11.9 MPa·m^1/2 and 10.45 MPa·m^1/2, respectively, representing increases of 4.9%, 11.2%, and 12.0% over those of C-TiCN.

(3) The enhanced surface hardness of HE-TiCN is primarily attributed to substantial lattice distortion and synergistic interactions among multiple principal elements within the high-entropy hard phases. The improved fracture toughness in the binder-rich region is associated with a combination of crack deflection, crack branching, and the formation of well-developed tear ridges that effectively dissipate energy during fracture.