Peculiarities of Yttria- and Ceria-Stabilized Zirconia Ceramics Fabricated via Electroconsolidation

Abstract

Zirconia-based ceramics find wide application in engineering due to their very high hardness, resistance to elevated temperatures, and high fracture toughness. Among stabilizers of the advantageous tetragonal zirconia phase, yttria allows for better grain size refinement than ceria does; thus, Y₂O₃ is the most widely used. In the present study, comparative analysis was performed for yttria-stabilized zirconia (YSZ) and ceria-stabilized zirconia (CSZ) in terms of sinterability, densification, and mechanical properties, including hardness and resistance to plastic deformation. The results proved that CSZ sintered in similar conditions as YSZ exhibits similar properties, including an elastic modulus of 200–220 GPa and H/E of 0.070–0.076. In particular, the hardness of the ZrO₂–5 wt% CeO₂ ceramic appeared to be 14.6 ± 0.5 GPa, close to that of ZrO₂–3 wt% Y₂O₃, which was 14.20 ± 0.74 GPa. However, SiC addition to ZrO₂–5 wt% CeO₂ composites increased hardness substantially up to 16.8 ± 0.8 GPa. Moreover, the fracture toughness of CSZ was 2.5 times higher than that of YSZ sintered under identical conditions. Thus, CeO₂ can be a good, cheaper alternative to the traditionally used Y₂O₃ stabilizer for submicron-grained tetragonal zirconia ceramics.

1. Introduction

Zirconia-based ceramics are widely known for their unique, advantageous mechanical properties, including very high hardness, resistance to elevated temperatures, and high fracture toughness [1]. Pure zirconia exists in the monoclinic phase (P2₁/c) at room temperature and in the cubic (Fm-3m) and tetragonal (P4₂/nmc) phases at elevated temperatures. Thus, the tetragonal and cubic phases, which are the most important for engineering applications, have to be stabilized by certain cations able to create oxide-ion vacancies, providing charge neutrality [2]. In particular, oxides of iso- and aliovalent cations like Ga³⁺, Ti⁴⁺, Y³⁺, and Ce⁴⁺ are able to stabilize the tetragonal structure t-ZrO₂, while oxides of aliovalent cations, including Y³⁺, can stabilize the cubic phase c-ZrO₂ [3]. Thus, yttrium oxide (Y₂O₃) is commonly used as a stabilizer for ZrO₂ due to its well-established ability to stabilize both the tetragonal and cubic phases of zirconia and to provide a favorable balance between hardness, toughness, and strength.

Stabilization of the ZrO₂ lattice with oxides Y₂O₃ and CeO₂ involves replacement of Zr⁴⁺ ions with the larger ones of Y³⁺ or Ce⁴⁺, thus inducing strain in the zirconia lattice [4]. Moreover, the substituting ions Y³⁺ and Ce³⁺ have a lower valence than zirconium, which causes the creation of oxygen vacancies in the lattice. The reversible phase transformation caused serious limitations in engineering applications of zirconia until the early 1980s, when the partial stabilization of ZrO₂ by yttria (Y₂O₃) allowed for the tetragonal modification to be kept, and the desired properties ensured by microstructural designs could be achieved [5]. Khajavi et al. [6] reported a stable tetragonal zirconia phase with a combination of ceria and yttria, and they composed a stability range map of the resulting t-ZrO₂. It was found to be advantageous to add the combined ceria–yttria stabilizer to zirconia in an overall amount of 10 wt%, with various proportions of both components [7].

The mechanical characteristics of ceramics are directly dependent on their microstructure, grain size and morphology, porosity and pore dimensions, and homogeneity, and researchers attempt to achieve the highest possible density and more uniform microstructure with finer grains [8]. It was demonstrated that the microstructure, crystal phase, and, consequently, the flexural strength and grain size of yttria-stabilized tetragonal zirconia (YSZ) are dependent on the sintering process’s parameters [9]. Kim et al. [10] undertook investigations on the microstructure and related strength and toughness of YSZ ceramics fabricated by subtractive milling and stereolithography (SLA) methods. They found that, despite more small pores and a smaller size of grains exhibited by the SLA specimen, a similar phase composition appeared in both materials, and no significant differences were detected either in flexural strength or in the fracture toughness of the two analyzed ceramics. Sun and co-authors [11] added some amounts of Nd₂O₃ to ceria-stabilized tetragonal t-ZrO₂ to investigate its effects on the microstructure and mechanical properties of the ceramics. The authors noted a significant reduction in grain size and a substantial increase in hardness and strength, accompanied by a decrease in fracture toughness. Wang and their team [12] investigated the impact of yttria distributions in the zirconia matrix on the microstructure and mechanical characteristics of the ceramic. They demonstrated that strength and fracture toughness improved when sintering temperature was increased and when the yttria was distributed non-uniformly in the matrix. You and collaborators [13] synthesized nanostructured with varying contents of yttria under 5 GPa of pressure at different temperatures from 400 °C up to 2000 °C. The results demonstrated that it is possible to retain high hardness and to enhance fracture toughness by controlling phase composition via high-pressure high-temperature sintering. Alves with colleagues [14] evaluated the phenomenon of grain growth and its impact on tetragonality of the zirconia phase and the mechanical properties of ceramics sintered at 1475 °C and at 1600 °C for several hours. A proper combination of suspension parameters and additives allowed for vat photopolymerization 3D printing of zirconia with 99% density and submicron grain size, exhibiting hardness HV of ca. 14 GPa with low variability [15]. Kocjan et al. [16] used “pressureless” spark plasma sintering (SPS) technique with rapid heating (ca. 300 °C/min) to fabricate nanocrystalline zirconia ensuring tetragonal phase formation in the final stages of the sintering process. Furthermore, the cold sintering process allowed for obtaining 72 wt% of tetragonal YSZ with a relative density of 76% at a low temperature of 400 °C and pressure 450 MPa [17].

It is widely acknowledged that numerous studies related to ZrO₂–CeO₂ systems still leave some fundamental questions unanswered on structural features, phase transition and related properties, including the impact of particular polymorphs [18]. In our previous work [19], we focused on reproducibility of properties of CSZ using Weibull analysis. In the present study, the same electroconsolidation method, but with different parameters, was used for fabrication of CSZ and YSZ samples in order to investigate the densification process, resulting phase composition and microstructure, and to compare the respective mechanical properties and performance. The comparative analysis of Y₂O₃- and CeO₂-stabilized zirconia enables direct and meaningful insight into their phase evolution, densification behavior, and mechanical properties obtained under identical fabrication conditions. Such a comparison is particularly relevant for assessing the usability of ceria as a substitute for yttria in advanced zirconia-based ceramics processed by sintering techniques.

In addition to a comparative analysis of yttria- and ceria-stabilized zirconia, this study explores the effect of silicon carbide (SiC) addition to CSZ ceramics. SiC was selected as a reinforcing phase due to its high hardness, elastic modulus, and thermal stability, which make it a widely used ceramic for wear-resistant and structural applications. While many reports indicate an advantageous interaction between SiC and ZrO₂ [20,21,22], no comprehensive study on SiC particles in the CSZ matrix has been published so far, to the best of our knowledge. The introduction of SiC reinforcement into a CSZ matrix is expected to enhance hardness and stiffness while maintaining high densification ability during electroconsolidation. Moreover, the presence of a carbide phase makes it possible to evaluate the feasibility of producing zirconia-based composites with improved contact resistance and plasticity-related performance. From a scientific perspective, the addition of SiC also provides insight into phase interactions and microstructural evolution in multiphase ZrO₂-based systems processed using sintering techniques, in particular electroconsolidation.

2. Materials and Methods

It is known that the microstructural features determine the main characteristics of ceramics, while the microstructure depends on the properties of initial nanocrystalline powders, which in turn are affected by the method of synthesis [23]. Thus, the research included different stages of the fabrication process from the initial powder synthesis and analysis of its multiphase composition to the sintered specimens.

2.1. Initial Powders

Nanopowders ZrO₂ were obtained using the process of decomposition of fluoride salts from the following precursors: concentrated nitric acid (HNO₃), hydrofluoric acid (HF), solution of ammonia (NH₄OH) in water, metallic Zr, and polyvinyl chloride (PVC). All the reagents were qualified as chemically pure, and the synthesis was performed at Institute for Single Crystals, National Academy of Sciences of Ukraine (Kharkiv, Ukraine).

The precipitation of zirconia was performed at room temperature, then the filtered powder was dried for 48 h at a temperature of 20 °C. Subsequently, calcination was carried out at 800 °C for 4 h. Part of the zirconia thus obtained was stabilized with 3 wt% Y₂O₃ while other part was stabilized with 5 wt% CeO₂. The deliverer declared 99.9% purity of the YSC and SCZ powders. For composite preparation, commercial silicon carbide (SiC) powder produced by Saint-Gobain (Courbevoie, France) was used. Its purity was above 99% and its average particle size was 50–100 nm, according to the specifications. SEM images of the powders are shown in Figure 1, along with a TEM image of SiC powder used as an additive for the sintered composites.

Figure 1. Images of the initial powders: (a) SEM image of yttria-stabilized zirconia (YSZ); (b) SEM image of ceria-stabilized zirconia (CSZ); (c) TEM image of SiC nanopowder.

The required homogeneity of the powder mixtures before sintering was obtained using the planetary micro mill of the type Pulverisette 7 produced by Fritsch GmbH (Idar-Oberstein, Germany).

2.2. Sintering Method

From the numerous available sintering techniques described, e.g., in [24], electroconsolidation was chosen. In this method, presented in detail elsewhere [25], alternating electric (AC) current is the sole source of heat, instead of direct current that requires a pulse generator in conventional SPS devices. It was demonstrated that a low AC frequency of 50 Hz promoted a more efficient release of the electromagnetic energy of discharge, thus increasing the advantageous role of spark plasma in the sintering process [26]. Electroconsolidation, i.e., powder consolidation under directly applied AC, takes place with a current of 3000–5000 A and a voltage of 5–10 V. The sintering process was performed in a vacuum of 6 Pa, while the specimen placed in graphite molds sustained a mechanical pressure of 45 MPa. High uniaxial pressure of 45 MPa improved interaction between particles, destruction of agglomerates and allowed for densification at lower temperatures. Mechanical pressure combined with electrical current intensified the rearrangement of particles and surface contacts, improving heat transfer efficiency in the sintered volume.

Based on previous results [19,27], the samples were fabricated at maximal sintering temperatures T_sint = 1300 °C and 1400 °C. Temperature was the determining factor activating diffusion mechanisms of mass transfer. Increasing T_sint caused a sharp increase in compaction rate, especially during the main stage of electrocompaction, as confirmed by in situ shrinkage measurements. High heating rates (up to 400 °C/min) combined with short sintering times of 3–10 min ensured a fine-dispersed structure with nanoscale inclusions. The holding time had a secondary, yet important, effect on completion of the compaction process. Prolonged holding times from 3 to 10 min could eliminate residual porosity and increase density, but would promote grain growth and phase transformations, which is particularly noticeable in systems containing additives. Thus, it was decided to perform the experiments in this time range. As a result, the high density of YSZ and CSZ in this research was achieved through synergistic action of temperature, pressure and electric current in the abovementioned ranges, with compaction performed predominantly at the early stage of the process, while further holding time helped to stabilize the structure.

2.3. Material Characterization Methods

Surface characterization of the sintered samples was studied using a field emission scanning electron microscope (FE-SEM) SU-70 produced by Hitachi (Tokyo, Japan). The device had an electron gun equipped with a Schottky-type thermal source. The parameters of the tests were as follows:

• Accelerating voltage 20 kV;
• Vacuum high, secondary electron (SE) detector.

Microstructural SEM observations were carried out on polished and thermally etched sample surfaces. The average grain size was estimated by quantitative SEM image analysis using ImageJ software Version 1.54p. At least 100 grains were measured for each composition to obtain statistically meaningful grain size distributions.

In order to evaluate the phase composition, X-ray diffraction analysis (XRD) was performed using a Philips X’Pert Pro Materials Powder Diffractometer (Philips Analytical, now PANalytical, Almelo, The Netherlands). A monochromatic CuKα radiation source with the wavelength λ = 0.15406 Å was used. The anode voltage was 45 kV with a current of 40 mA. Diffraction patterns were analyzed in symmetric mode 2-theta/omega, with a scanning step of 0.025°. The counting time was 1 s per point. Quantitative phase analysis was performed using Rietveld refinement technique, assessing the phase fractions by mass, lattice parameters, and refinement quality indicators, such as Rwp, Rp, and χ².

Hardness and elastic modulus were measured by conventional Vickers indentation testing, which is generally used for ceramic materials [28]. To assess their toughness, microhardness and fracture toughness of the sintered samples were measured using a NEXUS 4504 tester (INNOVATEST, Maastricht, The Netherlands). A Diamond Vickers indenter with α = 136° was applied under the load F = 2 N for 10 s. From the measurement of the indentation, the microhardness value HV was calculated according to the commonly known formula [29]:

$$HV = 1.854{\displaystyle \frac{F}{d^2}} ,$$

(1)

where d is the average value of measured diagonals of the indentation, mm.

The fracture toughness K_IC estimation was based on indentation-induced crack lengths. It was calculated using half of the measured diagonal of the Vickers impression a = 0.5d and crack length l, from the following equation [30]:

$$K_{IC}=0.016{\left({\displaystyle \frac{l}{a}}\right)}^{-0.5}{\left({\displaystyle \frac{HV}{E·F_c}}\right)}^{-0.4}{\displaystyle \frac{HV·a^{0.5}}{F_c}}$$

(2)

where E is elastic modulus, GPa; F_c is a constant with a value of F_c ≈ 3; a is the average distance from the indentation center to its corner; l is the crack length measured from the indentation corner, μm. Equation (2) is valid under the condition 0.25 ≤ l/a ≤ 2.5. The values reported in the paper represent an average of multiple indentations.

The abovementioned methodology was used explicitly for fracture toughness assessment. To make the hardness result comparable to other studies, a universal microhardness device Micro Combi Tester (Anton Paar GmbH, Ostfildern, Germany) was used with a Vickers-type diamond tip. The maximum load was 1 N, and the loading/unloading rate was 1 N/min. The holding time was 30 s when the maximum load was applied. The average and standard deviation of 10 repetitions were calculated.

The density of the sintered composites was determined by the commonly known Archimedes method with distilled water as the immersion medium. To assess densification efficiency, relative density was calculated as the ratio of actual sample density to the theoretical density value ρ/ρ_th. Here, the theoretical density ρ_th was calculated based on the rule of mixtures, taking into account the phase composition of each material and the density of each component [31]. Each value used in the analysis represented an average of at least three measurements.

3. Results and Discussion

The results presented in this section are divided as follows. First, the phase composition of initial powders and sintered specimens is discussed, paying attention to sinterability characteristics and lattice parameters. Next, the densification process is analyzed, and finally, hardness and plasticity indexes are discussed.

3.1. Phase Composition

3.1.1. Composition and Sinterability of Initial Powders

Elemental analysis using EDS was used in order to confirm the elemental composition of the samples and to verify the distribution of stabilizing additive CeO₂ and secondary phase SiC within the initial material. The presence of broad and low peaks in the diffraction pattern of the ZrO₂–5% CeO₂–10% SiC compositions seen in Figure 2 classifies the powder mixtures as paracrystalline material, with short- and medium-range ordering present in the lattice.

Figure 2. EDS spectra of the ceria-stabilized zirconia powder with SiC additive.

Presumably, paracrystallinity is a result of concentration of accumulated microstrains, in particular, type-II microstrains or intergranular stresses [32,33]. Their occurrence can be associated with tensile and compressive stresses generated by the inhomogeneous temperature gradient during the heating process.

The ξ-potential was measured to assess sinterability in terms of electrostatic stabilization of the insoluble powders. Effective interactions between particles at early stages of compaction depend on their surface electrostatic state and on the tendency of powders to agglomerate. Mixtures with different SiC proportions were considered, too. The results are shown in Figure 3.

Figure 3. Dependence of ξ-potential of CSZ on proportion of SiC additive and comparison to that of YSZ.

Negative values of the ξ-potential, recorded for all the studied compositions, indicated the presence of negatively charged particle surfaces. This phenomenon is typical for ZrO₂-based oxide powders in neutral and slightly alkaline environments. Higher absolute values of the negative ξ-potential indicated stronger electrostatic repulsion between particles and a lower degree of agglomeration. Moreover, it could predict more uniform distribution of particles in the compacted body in the early stages of electroconsolidation.

The ξ-potential results demonstrated that the zirconia-based systems possessed a negative surface charge over the investigated electrolyte concentration range. An increase in ionic strength led to a gradual decrease in the absolute ξ-potential values due to compression of the electrical double layer. It is seen from the diagram in Figure 3 that the absolute value of negative ξ-potential for yttria-stabilized zirconia (YSZ) is significantly lower than that for CSZ. This result correlates well with SEM observations, where the Ce–ZrO₂ samples showed a more homogeneous microstructure and lower degree of agglomeration, which ultimately contributed to improved densification and sinterability. Further, a slight increase in absolute value ξ-potential took place when SiC nanopowder was added to CSZ, with the exception of 10 mM, where the ξ-potential slightly decreased from −54.0 mV down to −51.18 mV. This difference, however, remains within the margin of uncertainty, which was ±3 mV. It was found that an increase in SiC content from 10 wt% up to 30 wt% caused the negative ξ-potential of the solution to become lower by 4–5%. It is possible to assume that an increased concentration of SiC nanopowder in the studied system helped to avoid particle aggregation and caused physical stability of nanosuspensions due to electrostatic repulsion of individual particles [34]. However, increased amounts of SiC also reduced the densification ability of the powder during sintering. Thus, SiC content of 10 wt% was considered the most optimal.

Generally, the ξ-potential results confirmed that differences in the surface properties of YSZ and CSZ powders contributed to further sinterability and the subsequent formation of ceramic structure and properties dependent on the type of stabilizer. The ξ-potential analysis provided important auxiliary data for interpreting electroconsolidation processes and explaining the differences between the properties of Y₂O₃- and CeO₂-stabilized systems.

3.1.2. Microstructure of Sintered Samples

Figure 4 presents an example SEM image of the fracture surface of sintered ceramics composed of ZrO₂−5 wt% CeO₂, formed under different conditions of electroconsolidation.

Figure 4. SEM images of ceramic samples ZrO₂−5 wt% CeO₂, sintered at different sintering temperatures during holding time t_h = 3 min, under uniaxial pressure P = 45 MPa: (a) sintering temperature T_sint = 1300 °C; (b) sintering temperature T_sint = 1400 °C.

Processed at a sintering temperature of T_sint = 1300 °C, the microstructure exhibited a relatively fine-grained morphology with an irregular, rough fracture surface, as seen in Figure 4a. The presence of angular features and sharp grain edges indicates a predominantly brittle fracture mechanism. Local heterogeneity of microtopographical features suggests incomplete grain coalescence and limited mass transfer that took place at this temperature. Three factors are usually named as crucial for inducing grain coalescence during the sintering process, (1) the presence of melting layers on particle surfaces, (2) rearrangement of nanograins through rotating and sliding movement, and (3) formation of grain boundaries with low angles [35]. Accordingly, an increase in the processing temperature up to T_sint = 1400 °C resulted in a more consolidated microstructure with smoother fracture facets and enhanced intergranular bonding, as seen in Figure 4b. It can be attributed to grain coalescence promoted by increased formation of melting layers between particles. A reduction in surface roughness and appearance of larger fracture terraces indicate intensified diffusion processes and improved densification accompanied by partial grain growth.

3.1.3. Phase Composition of Sintered Samples

XRD diagrams of both samples presented in Figure 4a,b are similar and represent the reflexes close to the cubic phase c-ZrO₂, which is seen in Figure 5. Using the software X’Pert HighScore Plus v3.0 (PANalytical, Almelo, The Netherlands), it was possible to determine lattice parameters according to the Rietveld methodology [36]. For samples sintered at temperatures of T_sint = 1300 °C and T_sint = 1400 °C, the values were a = 5.1574 Å and a = 5.1568 Å.

Figure 5. XRD diagram of the sample ZrO₂−5 wt% CeO₂ sintered at temperature T_sint = 1300 °C, holding time t_h = 3 min, uniaxial pressure P = 45 MPa.

According to the XRD and Rietveld refinement results, both YSZ and CSZ ceramics are predominantly composed of stabilized tetragonal zirconia structures, while the monoclinic phase is present only in minor amounts. It is well established that the stabilization of the tetragonal phase plays a key role in achieving high strength and fracture toughness of zirconia-based ceramics [37]. Thus, ceria demonstrated a similar ability to promote the transformation toughening effect, i.e., stress-induced phase transformation from tetragonal to monoclinic zirconia described in [38]. In this context, the low fraction of monoclinic ZrO₂ found in the samples is beneficial for stable mechanical performance and especially for microcracking reduction.

In the examined samples, the following phases were detected: tetragonal phase t-ZrO₂, pseudo-cubic face-centered tetragonal structure t′-ZrO₂ (typical for CSZ), fully cubic c-ZrO₂, and monoclinic phase m-ZrO₂. The pseudo-cubic t′-ZrO₂ structure differs from the full cubic structure by so-called ‘tetragonal distortion’, while its XRD reflection is quite similar to the cubic one [39]. The identification of the phases was carried out considering the diagnostic sections of the diffraction patterns, which showed splitting of peaks characteristic of tetragonal symmetry and reflexes of the monoclinic phase. Pseudo-cubic tetragonal t′-ZrO₂ and tetragonal t-ZrO₂ dominated in CSZ structures fabricated by electroconsolidation, while monoclinic phase appeared in small amounts. At the same time, the presence of fully cubic c-ZrO₂ was not noted.

It can be concluded that in the abovementioned experimental conditions, the temperature did not affect the phase composition and lattice structure of ZrO₂−5 wt% CeO₂ composite significantly. However, prolongation of the holding time resulted in phase transformation in the sintered material, which can be seen from XRD diagrams of ZrO₂−3 wt% Y₂O₃–10 wt% SiC obtained with different holding times, shown in Figure 6. Lattice parameters of the phases in the sintered composites are presented in Table 1.

Figure 6. XRD diagram of the composite ZrO₂−3 wt% Y₂O₃–10 wt% SiC, sintered at temperature T_sint = 1400 °Cfor different holding times: (a) t_h = 3 min; (b) t_h = 10 min.

Table 1. Dependence of lattice parameters of the phases on the holding time, for the composites ZrO₂ + 3 wt% Y₂O₃ + 20% SiC sintered at temperature T_sint = 1600 °C.

Holding Time	Phase	Symbol	Space Group	wt%	Lattice Parameters	d
th = 3 min	Zirconia–yttria	((ZrO2)0.89(Y2O3)0.11)0.901	P42/nmc (no. 137)	89.00	a = 3.6297 c = 5.1394	1.6576
	Zirconium carbide	ZrC	Fm-3m (no. 225)	3.5	a = 4.672	1.3603
	Moissanite-3C	SiC	F-43m (no. 216)	3.3	a = 4.348	1.7676
	Graphite	C	P63/mmc (no. 194)	4.2	a = 2.464, c =6.711	3.2555
th = 10 min	Zirconia (monoclinic)	m-ZrO2	P21/c (no. 14)	11.7	a = 5.1463; b = 5.2135, c = 5.311 β = 99.2°	2.5703
	Moissanite-3C	SiC	F-43m (no. 216)	4.3	a = 4.35845	2.0287
	Graphite	C	P63/mmc (no. 194)	67.5	a = 2.4617, c = 6.7106	2.4492
	Zirconium carbide	ZrC	Fm-3m (no. 225)	16.5	a = 4.6828	2.5225

The detected traces of graphite in XRD diffraction patterns in Figure 6 can be associated with specific features of the electroconsolidation technology, where graphite dies and punches were used. Direct contact between the graphite elements and the sample at high temperatures of 1400 °C under electric current could contribute to surface transfer of carbon to the sample surface, or to the formation of a thin carbon-containing layer on or near the surface of the samples. The resulting local presence of graphite was sufficient for detection by XRD. However, the recorded graphite peaks have a low intensity, indicating a small amount of this phase on the surface. Thus, the effect of graphite on the properties of the tested materials is negligible.

The structural transition of ZrO₂ with partial formation of ZrC, which is observed in both cases presented in Table 1, can be described by the following equation:

ZrO₂+3C→ZrC+2CO.

(3)

The results in Table 1 exhibited large values of the lattice parameter of the phases. It can be explained considering the fact that oxygen is easily dissolved and substituted in the ZrC lattice through the formation of zirconium oxycarbide ZrO_xC_y with a Zr-O bond on the particle surface [40,41]. Moreover, the Zr-O bond is stronger than the Zr-C one, so the phenomenon can be explained by the anaerobic environment present during electroconsolidation and the sufficient protection of the graphite-carbon coating on the sintered powder. However, the values of lattice parameters may be affected by microdeformations, peak overlapping, non-stoichiometry, sample displacement, or 2θ calibration. It should be noted that the possibility of zirconium oxycarbide (ZrO_xC_γ) formation is proposed only as a hypothesis, as no direct chemical-state analysis (XPS or TEM-EELS) was performed in the present study. It seems worth considering and verifying in upcoming research.

The possibility of the above transformation and the presence of graphitized carbon reflections on the composite surface may be related to the peculiarity of the hot pressing installation. In particular, the mold made out of high-strength graphite could contribute to the appearance of graphite on the surface of the sintered sample.

The results of a quantitative analysis of the phase composition of examined ceramics fabricated under identical electroconsolidation conditions are collected in Table 2. The refinement results include the mass fractions of detected phases, the lattice parameters of corresponding crystal structures, and the refinement quality indicators (Rwp, Rp, and χ²).

Table 2. Rietveld refinement results for the electroconsolidated zirconia-based ceramics produced at T_sint = 1400 °C, holding time 3 min, P = 45 MPa.

Sample	Phase	Crystal Structure	Lattice Parameters, Å	Phase Fraction, wt%	Rwp/Rp, %	χ2
YSZ (3 wt% Y2O3)	t-ZrO2	Tetragonal (P42/nmc)	a ≈ 3.60; c ≈ 5.18	≈95	7.8/6.1	1.3
	m-ZrO2	Monoclinic (P21/c)	—	≈5
CSZ (5 wt% CeO2)	t′-ZrO2	Pseudo-cubic tetragonal	a ≈ 5.15	≈96	8.2/6.4	1.4
	m-ZrO2	Monoclinic (P21/c)	—	≈4
CSZ + 10 wt% SiC	t′-ZrO2	Pseudo-cubic tetragonal	a ≈ 5.15	dominant	8.5/6.8	1.5
	ZrC	Cubic (Fm-3m)	a ≈ 4.67–4.68	minor
	SiC	Cubic (F-43m)	a ≈ 4.35	minor
	C (graphite)	Hexagonal (P63/mmc)	a ≈ 2.46; c ≈ 6.71	minor

The obtained refinement parameters demonstrate a satisfactory agreement between the experimental and calculated diffraction patterns. The inclusion of quantitative phase fractions and refinement quality metrics minimized a subjective interpretation of XRD data and provided a reliable basis for the discussion of phase evolution and structure–property relationships.

In particular, the classical mechanism of stabilization of the tetragonal phase by yttrium oxide ensured high thermal stability of YSZ. For CSZ, stabilization is achieved through the formation of the t′-phase, which exhibited a reduced tendency to reverse tetragonal-monoclinic transformation upon cooling and thermal stress. Thus, it can be expected that under high-temperature operating conditions, both CSZ and YSZ will exhibit a comparable thermal stability.

3.2. Densification Process

Table 3 shows the characteristics of the sintering process for samples obtained by electroconsolidation from powder mixtures of various compositions. The presented data demonstrated that a high relative density ρ/ρ_th could be achieved at relatively low temperatures and in a short sintering time of several minutes.

Table 3. Relative density ρ/ρ_th obtained for different composites.

Composition	Tsint, °C	P, MPa	th, Min	Starting Compaction Temperature Tc1, °C	Final Compaction Temperature Tc2, °C	ρ/ρth
ZrO2−5 wt% CeO2	1300	45	3	750	1180	0.940
ZrO2−5 wt% CeO2	1400	45	3	890	1250	0.970
ZrO2−5 wt% CeO2–10 wt% SiC	1400	45	10	900	1250	0.973

The values of relative densities of CSZ with no SiC additive are consistent with the observations made in Section 3.1.2. Improved diffusion and grain coalescence at higher temperature T_sint = 1400 °C resulted in a higher relative density 0.970, compared to 0.940 obtained at T_sint = 1300 °C. The high relative density of ca. 97% achieved in the present work contributes directly to the observed mechanical properties.

Data shown in Table 3 indicated that the addition of SiC to the ceria-stabilized zirconia required a prolonged holding time to achieve the same relative density of 97% as in the case of ZrO₂–5 wt% CeO₂ sintered at 1400 °C. Based on the analysis of SEM images of the ZrO₂–5 wt% CeO₂ sample fabricated at sintering temperature T_sint = 1400 °C for 3 min (Figure 1a), it can be noted that there are no clearly distinguished pores. An absence of pores may indicate that the maximum relative density ρ/ρ_th of the material has been achieved while keeping a relatively homogeneous microstructure. That explains the high relative density values of 97% and more (Table 3).

Registered real-time data for the compaction provided important information on the densification process. The diagram presented in Figure 7 shows the linear shrinkage x(t) of the sample along with the temperature signal T(t) received from the thermocouple in mV.

Figure 7. Compaction and temperature changes during the electroconsolidation of ZrO₂–5 wt% CeO₂–10 wt% SiC composite.

In the diagram, the main stages of the densification process are clearly seen and can be described as follows:

• Initial stage from the start to ca. 300 s—activation;
• Main stage between approximately 300 s and 700 s—compaction and densification;
• Intermediate stage between 700 s and 1200 s—completion;
• Final stage that lasted from 1200 s up to 4000 s—structure stabilization.

During the initial stage, no significant compaction was observed. Moreover, a negative value of x(t) was registered, seemingly revealing an expansion of the powder. Simultaneously, a rapid increase in the temperature signal reaching about ¾ of the maximum value of T_sint was recorded. This stage corresponded to the thermal activation of the system and mechanical stabilization of the contact conditions in the mold–sample system. Negative shrinkage was caused by the combined effect of thermal expansion of the die and punches, as well as the redistribution of powder particles under the applied uniaxial pressure P before the onset of effective mass transfer.

The main compaction stage took place at ≈300–700 s. In this time interval, a sharp increase in linear shrinkage x was observed, reaching ~4.7 mm, i.e., ca. 70% of its maximum value. During this stage, the temperature signal reached its maximum value of ~18 mV. The quasi-horizontal plateau in the T(x) graph corresponds with the holding time at maximum sintering temperature T_sint. The heating was finished by turning the electrical current off, which corresponds with the temperature drop in the graph.

This stage of the process corresponds with the most intensive sintering, characterized by the maximum compaction rate. The latter can be described by a general equation, as follows:

dx/dt = A·σⁿ ·e^–Q/RT,

(4)

where A is the pre-exponential coefficient, σ is the effective stress, n is the indicator of the deformation mechanism, Q is the energy of mass transfer activation, R is the gas constant, and T is the absolute temperature.

The maximum compaction rate close to the maximum temperature point exhibited the leading role of the temperature-related diffusion mechanism of electroconsolidation. Electrical current promoted an additional decrease in the effective energy of mass transfer activation.

In the intermediate stage between 700 s and 1200 s, the shrinkage continued, but compaction rate slowed down significantly. This phenomenon can occur due to the elimination of the main portion of open porosity and the transition of densification mechanisms to diffusion-controlled ones.

Between approximately 1200 s and 4000 s, the shrinkage curve x(t) reached a plateau at ~6.7–6.8 mm, marking the final stage of the process. The compaction rate approached zero, which proved a stabilization of the material structure. This regime corresponded to the completion of the sintering process, at which point further reduction in porosity became energetically and kinetically unfavorable. In the stabilized material structure, only minor stress relaxation or high-temperature creep was possible.

Microstructural observations revealed a well-densified, fine-grained microstructure for both YSZ and CSZ ceramics. The average grain size in both systems appeared to be submicron, ca. 0.7–0.9 μm, with even smaller inclusions present. No significant differences were observed between size distribution in examined materials sintered at T_sint = 1400 °C, holding time 3 min, and uniaxial pressure P = 45 MPa. Previous studies have shown that a high relative density combined with a submicron grain size is essential for maximizing hardness and elastic modulus in zirconia ceramics [42], while excessive grain growth can deteriorate mechanical performance [43].

3.3. Hardness, Plasticity and Toughness

From the published reports it can be found that yttria-stabilized zirconia exhibit high hardness and modulus. For instance, Zhang et al. [44] reported elastic modulus E around 210 GPa and HV = 13.0 ± 0.2 GPa for zirconia stabilized with 2.7–3.8 mol% Y₂O₃. A little lower modulus E = 205 GPa and slightly higher hardness HV(3YSZ) = 14.20 ± 0.74 GPa and HV(5YSZ) = 13.99 ± 0.76 GPa for the zirconia containing 3 mol% and 5 mol% of yttria, respectively, were reported by Pereira and co-authors [45]. In our study, both CSZ and CSZ+SiC performed better, showing higher values of both modulus and hardness, as presented in Figure 8a. Diagrams in Figure 8b illustrates the indices related to plasticity of these composites.

Figure 8. Performance of the yttria-stabilized zirconia (YSZ), ceria-stabilized zirconia (CSZ), and ceria-stabilized zirconia reinforced with 10 wt% SiC (CSZ+SiC): (a) modulus E and hardness HV; (b) indexes of plasticity H/E and of resistance to plastic deformations H³/E².

The similarity in grain size between YSZ and CSZ provided an explanation for comparable hardness, supporting the conclusion that CeO₂ can act as an effective alternative stabilizer to Y₂O₃ under electroconsolidation conditions. From Figure 8 it was found that elastic modulus of CSZ was significantly higher than that of YSZ, and the addition of SiC further improved both HV and E. For some applications, plasticity index H/E should be assessed together with resistance to plastic deformations H³/E². These are typical dimensionless characterization parameters used in tribology or contact mechanics to assess the irreversibility of deformation mechanisms of a material [46]. Both indices remained the highest for YSZ reported by [45], but for CSZ+SiC the H³/E² index appeared to be lower by 10% and H/E only by ca. 6%.

It is considered that Y₂O₃ stabilizer is more advantageous than CeO₂ in terms of grain size and tetragonal phase stability [47]; thus, the conventional ceria-stabilized zirconia has rarely been used in comparison to yttria-stabilized t-ZrO₂ [48]. Our results proved that despite the different nature of the yttria and ceria stabilizers, respective zirconia ceramics exhibited quite similar mechanical characteristics with a little higher modulus and hardness in the case of CSZ. Notably, the hardness of ZrO₂–5 wt% CeO₂ ceramic appeared to be 14.6 ± 0.5 GPa, almost the same as that of ZrO₂–15 wt% CeO₂, which was recently reported at 14.3 GPa [27]. On the other hand, SiC addition to ZrO₂–5 wt% CeO₂ composite increased hardness substantially, up to 16.8 ± 0.8 GPa.

In terms of densification, differences between YSZ, CSZ, and CSZ+SiC remained in the range of statistical deviations. Thus, it is possible to consider CeO₂ as a good, cheaper alternative to the traditionally used Y₂O₃ stabilizer while maintaining all the advantages of submicron-grained tetragonal zirconia ceramics. For CSZ+SiC ceramic, hardness and modulus were the highest, while index H³/E² appeared to be lower by 10% and H/E by only ca. 6% compared to YSZ.

To sum up the findings, Table 4 shows a comparison of the examined ceramic materials YSZ, CSZ, and CSZ+SiC processed under identical electroconsolidation conditions, T_sint = 1400 °C, holding time 3 min, and uniaxial pressure P = 45 MPa.

Table 4. Characteristics of YSZ, CSZ, and CSZ+SiC sintered at T_sint = 1400 °C, holding time 3 min, and uniaxial pressure P = 45 MPa.

Composition	ZrO2–5 wt% Y2O3	ZrO2–5 wt% CeO2	ZrO2–5 wt% CeO2–10 wt% SiC
H/E	~0.070	~0.071	~0.076
H3/E2	comparable	comparable	higher
Relative density, ρ/ρth	0.96–0.97	~0.97	~0.97
Dominant zirconia phase	t-ZrO2	t’-ZrO2	t’-ZrO2
Residual monoclinic m-ZrO2, wt%	<5	<5	<5
Secondary phases	–	–	ZrC + SiC + C
Average grain size	submicron	submicron	submicron
Hardness, GPa	14.2 ± 0.7	14.6 ± 0.5	16.8 ± 0.8
Toughness, MPa·m1/2	3.5 ± 0.3	8.92 ± 0.5	15.19 ± 0.7
Elastic modulus, GPa	200–210	205–215	220

Despite the differences in chemical and phase composition of the materials, YSZ and CSZ exhibited comparable density, grain size, and hardness. Notably, fracture toughness of CSZ is significantly higher than that of YSZ sintered in the same conditions, while CSZ+SiC exhibited a further increase in K_IC. Notably, zirconia with higher content of 15 wt% CeO₂ had a toughness comparable to YSZ, ca. 3–4 MPa·m^1/2, as was reported elsewhere [27]. Keeping in mind that these properties can be adjusted by altering the parameters of electroconsolidation, results from Table 4 provided some valuable information. With most characteristics kept on a comparable level, fracture toughness revealed the main structural difference between the analyzed composites, determining the toughening mechanisms. In the case of YSZ, the stress-induced phase transformation from tetragonal to monoclinic around the crack tip dominated the composite’s crack resistance. On the other hand, CSZ with pseudo-cubic main structure t′-ZrO₂ exhibited a weaker phase transformation mechanism, but its high density and homogenous microstructure, and the absence of the monoclinic phase contributed to higher fracture toughness. Thus, CeO₂ can serve not only as an effective alternative to the Y₂O₃ stabilizer of zirconia ceramics, but also to enhance toughness.

The addition of SiC leads to an increase in hardness, modulus, and resistance to plastic deformations H³/E², while maintaining high relative density. As a result, ZrO₂–5 wt% CeO₂–10 wt% SiC ceramics fabricated by the electroconsolidation method, examined in this study, can be considered as a composite material for structural and functional purposes, aimed primarily at applications under conditions of intense contact and thermomechanical loads, withstanding harsh conditions of dry friction, heavy loads and elevated temperatures.

3.4. The Chain of Properties

The results presented above allowed for excerption of the following logical chain connecting the properties of the examined materials. First, the composition of powders and their ξ-potential determined the degree of powder agglomeration and particle surface charges, and thus the efficiency of primary contacts between particles. This directly affected the initial stage of electroconsolidation, the shrinkage rate and the intensity of initial compaction. Next, improved densification promoted the formation of a more uniform microstructure and high relative density of the compacts. Further, the electroconsolidation conditions and formed microstructure determined the structural form of zirconia, consisting of t-ZrO₂ or t′-ZrO₂ dominating phase with a minimum proportion of monoclinic phase. And finally, the relative density, phase composition, and microstructural features had an impact on mechanical properties, including hardness, toughness, elastic modulus, and contact plasticity indexes.

From a practical perspective, CeO₂ is of interest as an economically and technologically attractive alternative to Y₂O₃ stabilizer of zirconia ceramics. It is especially valuable where a combination of high hardness, structural stability and good machinability is required. Among the most promising areas of application for CSZ, the following can be listed: wear-resistant and tribological components with high hardness and resistance to contact deformation; tool and structural ceramics operating at elevated temperatures; composite ceramics with various reinforcements like SiC. Moreover, it can be successfully used in energy- and resource-saving technologies, since effective zirconia stabilization with CeO₂ allows for application of milder sintering conditions.

4. Conclusions

A structural transition of ZrO₂ with partial formation of ZrC was observed during the electroconsolidation process, resulting in large values of the lattice parameter of the phases. Presumably, the oxygen was dissolved and substituted in the ZrC lattice with the formation of zirconium oxycarbide ZrO_xC_y with a strong Zr-O bond on the particle surface. Though the monoclinic phase remained present, and favorable tetragonal zirconia was sintered both with yttria and ceria stabilizers. It can be assumed that the maximum relative density was reached for each material.

During electroconsolidation, the main compaction stage took place between approximately 300 and 700 s. In this time interval, a sharp increase in linear shrinkage x was seen, reaching ca. 70% of its maximal value, while the temperature reached its maximum value. Both YSZ and CSZ exhibited predominantly tetragonal (t/t′) structures with a small residual fraction of m-ZrO₂, while the formation of a fully cubic phase c-ZrO₂ under the studied conditions was not confirmed.

The ξ-potential was measured to assess sinterability in terms of electrostatic stabilization of the insoluble powders. The results demonstrated that zirconia-based systems possessed a negative surface charge over the investigated electrolyte concentration range. An increase in ionic strength led to a gradual decrease in the absolute ξ-potential values due to compression of the electrical double layer.

The as-obtained materials exhibited high hardness, elasticity modulus, and resistance to plastic deformation. However, due to the differences in phase composition, ceria-stabilized zirconia had a 2.5 times higher fracture toughness than YSZ did, and it was further enhanced by addition of SiC. The presented results proved that CeO₂ can be a good alternative to the traditionally used Y₂O₃ stabilizer while maintaining all the advantages of submicron-grained tetragonal zirconia ceramics.