Design and Mechanical Properties of ZTA–Niobium Composites with Reduced Graphene Oxide

Abstract

Niobium–graphene oxide–zirconia-toughened alumina (ZTA) composites were produced by wet mixing and spark plasma sintering. The microstructure and mechanical properties of this novel composite have been studied. The results show that niobium particles are homogeneously dispersed in the ZTA matrix. Raman spectroscopy confirmed the thermal reduction in graphene oxide during sintering. The presence of ductile metal and graphene flakes leads to an increase in the crack resistance value of the ZTA matrix. The developed composites demonstrate a fracture toughness of 16 MPa∙m^1/2, which is three times higher than ZTA ceramic composites. The high toughness values found in this new composite are a consequence of the strong interaction between the simultaneous action of several toughening mechanisms, specifically involving crack trapping, crack blunting, crack renucleation, and the bridging mechanisms of the metallic and graphene particles. Moreover, this increase has also occurred due to the enhancement of the transformability of zirconia in ceramic–metal composites.

1. Introduction

Aluminum oxide (Al₂O₃) has been the focus of close attention from developers of new materials for several decades [1,2,3,4,5,6,7,8,9]. The possibility of synthesizing alumina ceramics from readily available and cheap raw materials contributes to its widespread use; for example, as high-speed cutting tools, implants and prostheses, electrical and thermal insulation, wear-resistant parts, coatings, etc. Its exceptionally high-temperature resistance and chemical inertness are very attractive properties that open up the possibility of using oxide materials at high temperatures or aggressive environments, as well as both simultaneously. High hardness, chemical inertia, high compressive strength, and high thermal insulation properties make aluminum oxide a potential candidate for use as a structural material in high-temperature conditions; however, the main obstacle to its use in this capacity is the inherent fragility of ceramics due to low fracture toughness and low bending strength [10,11]. Zirconium oxide (ZrO₂) has a record of high bending strength and fracture toughness among oxide materials due to its ability to convert the tetragonal phase of ZrO₂ into a monoclinic one; in other words, the t → m transformation. Since the monoclinic phase occupies a larger volume than the tetragonal one, compressive stress arises in the matrix when this transition is activated (for example, due to external stress). Compressive stresses act as locking forces, closing the surfaces of cracks, thereby reducing the intensity of their propagation. With the discovery of this mechanism, a vast number of studies on ceramic materials of the ZrO₂-Al₂O₃ (ZTA) system and methods for their production have emerged. The degree of the t → m transition affects the amount of material capable of transformation. If the zirconium oxide is completely stabilized, the ability to transition is lost, and this mechanism is not activated, despite the presence of ZrO₂ in the composite. Partially stabilized zirconium dioxide is used to produce ceramic materials with increased bending strength and crack resistance [12,13,14]. Such materials are used in industry to create high-strength, wear-resistant ceramics that work at moderate temperatures. In addition, Pecharroman et al. reported that it is very important to control the level of ZrO₂ in ZTA composites [15]. In order to avoid spontaneous transformation of ZrO₂, this limit must be below the percolation threshold, which has been found to be 16 vol.% or 22 wt.%. However, rapid crack propagation under severe loading conditions, even if it is below the strength of the material, reduces the reliability of ceramics. Therefore, much attention is dedicated to studying the mechanisms of crack propagation and the toughening of ceramics with such components that activate toughening mechanisms and increase the fracture toughness of the material. Technological techniques for increasing the strength of monolithic oxide ceramics are based on two main approaches: reducing the grain size, which helps to reduce creep by providing an initial defect size comparable to the grain size, and the introduction of reinforcing additives that provide a moderate (without catastrophic failure) reaction to existing or emerging defects. The concept of reinforcing brittle ceramics with ductile metals has been employed to enhance the mechanical properties of ceramics. Rodriguez-Suarez et al. discussed ceramic (alumina, zirconia, spinel)–metal (nickel, niobium, tungsten) micro/nanocomposites mechanical and tribological properties, such as hardness, wear resistance, and toughness [16]. The presented results demonstrate that the studied composites and ductile phase toughening concepts can be used to enhance the toughness and crack growth resistance in custom-made ceramic-based materials using smart microstructural control from the nano to the micro-scale. Mattern et al. studied interpenetrating alumina–aluminum composites obtained by the infiltration of molten aluminum into porous ceramic preforms by a direct squeeze casting process [17]. The structure of micropores was studied based on the content and size of spherical pore-forming agents. The microporosity of the pore walls was affected by changes in the sintering temperature. Yeomans presented a review of the alumina matrix with discrete ductile (nickel, iron, molybdenum, copper, and silver) particle composites with [18]. It has been reported that the reinforcing particles provide an increase in toughness due to plastic deformation and bridging an advancing crack. Diaz et al. fabricated zirconia/zirconium composites using a wet processing route and spark plasma sintering [19]. The main toughening mechanisms operating in the studied composites are the blunting of cracks and their branching. There was no clear evidence of the formation of a bridging mechanism of metal particles due to their strong adhesion to the ceramic matrix. Thus, the toughness value of composites was slightly higher compared to that of monolithic ceramics. The works of [20,21,22] show the possibility of using mechanical alloying of a powder mixture containing 10 vol.%, 30 vol.%, and 50 vol.% of niobium (Nb) and Al₂O₃ to obtain dense sintered samples. The strength and fracture toughness of the composites increase as the niobium content increases due to the strong interfacial bonding of the composite and extensive plastic bridging of the niobium ligaments. In addition, various intermetallic phases (AlNb₂, Al₃Nb) and niobium oxide (NbO) were detected in the sintered samples.

Several investigations of zirconia–niobium composites concerning mechanical properties [23], the interface between Nb and Al₂O₃ [24], and low-temperature degradation [25] were performed. The mechanical [26,27,28,29] properties of aluminum oxide composites reinforced with niobium particles were also studied. In addition, the damage tolerance [30] and mechanical properties [31] of alumina–zirconia–niobium micro–nano-hybrid composites were studied. The synergistic effect resulting from the interaction of ZrO₂ transformation toughening mechanisms and the bridging of niobium inclusions was found to increase the toughness, flaw tolerance, and crack growth resistance of these composites. It should be noted that all of the above studies were conducted at room temperature.

Weidner et al. studied microstructure and high-temperature mechanical properties under compressive loading of Nb-Al₂O₃ refractory composites with both 11 vol. % and 21 vol. % of niobium [32]. It has been demonstrated that the compressive strength of these composites decreases significantly with increasing temperature. On the other hand, even serious damage due to cracks running along the ceramic–metal interface from coarse-grained Al₂O₃ particles did not lead to complete destruction of the material during compression tests at temperatures from 1300 to 1500 °C.

The particles used can be represented as whiskers, fibers, flakes, etc. The bridging ligaments exert closure stresses that reduce the stress intensity at the crack tip and offer resistance to further crack propagation. As plotted in Figure 1, the bridging effect occurs when the metal particles either dissipate the elastic energy upon breaking (1, 9), debond from the ceramic matrix and pull out (2, 8), or shed load from the crack tip while remaining intact (3, 7). Additionally, the crack deflection effect occurs when a crack tip meets with a particle that forces the crack to change the direction of its growth (4). On the other hand, ductile inclusions provide blunting of crack tips (5) or stress-releasing plastic deformation in the process zones of growing cracks (6). The magnitude of the crack closure is based on the stress–strain relation of the ligaments, which is influenced by the ligament yield stress, diameter, orientation, volume fraction, and the nature of the ceramic–metal interface.

Metal phases have a plasticity that ceramics do not possess; therefore, in metal–ceramic composite materials, the energy of plastic deformation can be activated when a crack propagating in the ceramic matrix reaches a metal particle. Part of the crack propagation energy is absorbed by plastic deformation, and the remaining energy is no longer sufficient for further crack propagation. In addition, deformed particles can connect cracks, thereby overlapping stresses on each other and reducing the intensity of crack propagation. Moreover, alumina, zirconia, and niobium have found applications in clinical practice, mainly in dental and joint replacement applications. Bartolomé et al. investigated the biological tolerance of zirconia/Nb composites under both in vitro and in vivo conditions [33]. In vitro studies demonstrated that composites showed no deleterious effect on cell proliferation, extracellular matrix production, or cell morphology. In vivo studies concluded that composites are biocompatible and osteoconductive. The dense composite laminates made from Al₂O₃ and Nb have been studied in vitro for potential use as an alternative femoral head material by Rahman et al. [34]. This composite showed low wear on the dense surface of Al₂O₃ with the ductility of the metal femoral head, which can reduce the risk of catastrophic brittle fracture of femoral heads made of Al₂O₃ in vivo.

However, when choosing materials for medical use, it is necessary to take into account not only their mechanical properties and biocompatibility but also tribological properties. Portu et al. investigated the sliding wear behavior of Al₂O₃-based composites containing 15 vol.% and 25 vol.% of niobium against tungsten carbide (WC) disk [35]. The friction coefficient decreased with the increase in load due to the presence of transient phases triggered by the reaction of Nb with W, C. The wear behavior of these composites was found to be very similar to that observed for pure alumina. Gutiérrez-González et al. evaluated the friction and wear behavior of ultra-high molecular weight polyethylene (UHMWPE) against 3Y-TZP/niobium composite after accelerated aging [36]. The results show that the surface roughness of the metal–ceramic composite remains unchanged during the aging test; therefore, there is a lower coefficient of friction and wear rate for composite/UHMWPE pairs than for 3Y-TZP/UHMWPE pairs.

The friction pair in endoprosthetics should be an analog of the joint and allow it to perform the necessary functional roles, replacing the natural joint. The choice of the materials for the head and the cup took into consideration not only properties such as mechanical resistance and biocompatibility but also friction and wear. Graphene has a friction-reducing potential for solid lubricating and can decrease the friction force acting on the contact surfaces at the micro- and nano-scale [37,38]. Recent studies report on the enhancement of wear performance of ceramic-based nanocomposites in case graphene is added [39,40]. Therefore, graphene oxide (GO) was added to the ceramic matrix to improve the wear behavior of ZTA. Niobium was chosen as the second phase of the metal not only because of its plasticity and mechanical properties but also because of its biocompatibility and corrosion resistance [41,42]. The aim of the current study is to investigate the mechanical performance of Al₂O₃-ZrO₂ (ZTA) composites with reduced graphene oxide (rGO) reinforced with 20 vol.% of niobium.

2. Materials and Methods

2.1. Graphene Oxide Preparation

A modified Hummers method was used to obtain graphene oxide (GO) from graphite powder. This process has been described in more detail in previous papers [43,44]. In this work, commercial graphite powder (Plasmotherm, Moscow, Russia) with a median particle size d₅₀ = 3 mm was used.

2.2. Powder Processing and Sintering

Commercial t-ZrO₂ (3Y-TZP, 3 mol% Y₂O₃; TZ-3YS-E, Tosoh Corp., Tokyo, Japan), alumina (α-Al₂O₃, A16SG, Alcoa, New York, NY, USA), and niobium (Nb, Goodfellow, Huntingdon, UK) powders with particle size d₅₀ = 0.26 µm, d₅₀ = 0.30 µm and d₅₀ = 74 µm, respectively, were used in this work (Figure 2).

Niobium powder was milled in an HD-01 attritor (Union Process, Akron, OH, USA) with zirconia balls in a Teflon container for 4 h, using isopropyl alcohol as the liquid medium (Figure 3).

Due to this process, niobium particles with a lamellar–flaky shape, a high aspect ratio, and a d₅₀ = 44 µm were obtained. To fabricate the ZTA/20Nb ceramic composites, 56.48 wt.% Al₂O₃ + 9.20 wt.% ZrO₂ + 34.10 wt.% Nb of powders was placed in a plastic container with Al₂O₃ balls (diameter 3 mm), distilled water, and Dolapix CE 64 as a dispersant. The obtained mixture was wet-mixed in a multi-directional mixer for 24 h at 150 rpm and subsequently dried by FreeZone2.5 freeze-drying system (LabConco, Kansas, MO, USA, Figure 4).

Distilled water with a pH value of 10 was added to the produced mixture, which was then dispersed under mixing for 30 min. To obtain compositions with the required content of graphene oxide, it was added drop by drop to the suspensions and stirred for 1 h. To obtain the powder mixtures for further consolidation by SPS, resulting suspensions were dried at 110 °C in a Lab spray dryer (B-290, Buchi, Flawil, Switzerland, Figure 5).

Thereby, ZTA/Nb powders with GO content of 0.5 vol.% (0.22 wt.%) were prepared. Powder mixtures were sintered in an H-HP D 25 SD Spark Plasma Sintering machine (FCT Systeme GmbH, Rauenstein, Germany) in a vacuum at 1500 °C, applying a heating rate and pressure of 100 °C/min and 80 MPa, respectively, (Figure 6).

The isothermal hold at the final temperature was 3 min. The produced samples had a diameter of 20 mm and a height of 3 mm (Figure 7). During sintering, the in situ thermal reduction in the GO takes place. In this work, the carbon material obtained after consolidation will be referred to as reduced graphene oxide (rGO).

2.3. Microstructural Characterization

A field emission scanning electron microscope, LYRA3 (Tescan, Brno, Czech Republic), equipped with an X-Act energy dispersive spectroscopy detector (Oxford Instruments, Abingdon, UK) and a focused ion beam scanning electron microscope (FIB-SEM) AURIGA 60 CrossBeam Workstation (Carl Zeiss Microscopy, Jena, Germany), was used to study the microstructure (Figure 8). Before the studies, the samples were polished using diamond polishing slurries with grit sizes ranging from 9 μm to 1 micrometer. After polishing, the samples were washed in an ultrasound bath in ethanol for 15 min and dried using compressed air.

2.4. X-Ray Diffraction (XRD) and Raman Characterization

Phase identification of sintered samples and raw powders was carried out by X-ray diffraction in an Empyrean diffractometer (PANalytical, Almelo, The Netherlands) with radiation source Cu–Kα (λ = 1.5405981 Å) operated at an intensity of 30 mA and 40 kV [18]. The analysis was carried out with a scanning speed of 0.06/min and a step size of 0.05. The analysis of m- and t-ZrO₂ phases on the polished and fractured surfaces of the samples was carried out using the XRD technique. The amounts of m-ZrO₂ and its volume fraction were estimated according to the Garvie and Nicholson [45] as well as the Toraya et al. [46] methods, respectively. Raman analysis was achieved in a Raman analyzer DXRTM2, (Thermo Fisher Scientific, Waltham, MA, USA) using a 532 nm laser with a power of 2.0 mW for the control of the graphene-based mixtures and sintered composites. A 50× optical microscope objective was applied to focus the laser beam on the studied area into a 2 μm spot with an accumulation time of about 10 s [47].

2.5. Density and Mechanical Properties Characterization

The density of the sintered composites was measured using the Archimedes’ method in distilled water. The theoretical density was calculated using a ZrO₂ density of 6.05 g/cm³, Al₂O₃ density of 3.98 g/cm³, Nb density of 8.58 g/cm³, and GO density of 2.2 g/cm³.

Their Vickers hardness (Hv) was measured from 10 imprints (indenter Qness, Salzburg, Austria) per sample under load and loading time 294 N and 10 s, respectively, (Figure 9). To estimate the average of hardness values, the following equation was used:

$${\mathrm{H}}_{\mathrm{V}}={\displaystyle \frac{0.1891\mathrm{P}}{{\mathrm{d}}^2}}$$

where P was the set load (N); and d is the average length of two diagonals (mm).

The values of Vickers indentation fracture toughness (KC) were estimated as proposed by Miranzo and Moya [48]. The flexural strength (σf) was evaluated through a biaxial bending test (ISO 6872). Each sample was placed onto a device with three 3 mm balls made of hardened steel and disposed on a holder (10 mm in diameter) at 120° to each other. The load was applied with an AutoGraph AG-X (Shimadzu Corp, Kyoto, Japan) universal testing machine by means of a plain head with a diameter of 1.6 mm at a speed of 1 mm·min⁻¹ until failure (Figure 10). The specimen thickness was measured at the breakage point. The strength value was calculated by averaging the data after testing 12 sintered disks. The equation for strength calculations was described in a previous work [49].

3. Results and Discussion

Representative X-ray diffraction (XRD) patterns corresponding to raw and sintered materials are shown in Figure 11.

According to the information provided by the International Center of Diffraction Data (ICDD), the position and intensity of the peaks found in the studied materials corresponded with the characteristic peaks of zirconia (monoclinic, 24-1165, and tetragonal, 83-0113), alumina (81-1667), and niobium (35-0789), respectively. As expected, the found phases fitted with the crystallographic forms of zirconia, alumina, and niobium, indicating the non-existence of further compounds as a result of external contamination or reactions during the wet processing (Figure 11e). The diffractograms of Nb powder before and after milling (Figure 11a,b) show that pronounced texturing and broadening (2θ~56°) occur as a result of deformation and an increase in the lattice strain. In addition, XRD analysis of the samples showed a complete conversion of the monoclinic ZrO₂ to tetragonal ZrO₂ after SPS (Figure 11f).

Raman spectroscopy is a powerful tool to analyze the presence of graphitic structure in the composites. The Raman spectra of studied specimens are presented in Figure 12.

Raman spectrum of the graphite powder (Figure 12a) demonstrates the characteristic peaks located at ~1584 cm⁻¹ (G) and it is more intense than the 2D (2720 cm⁻¹) and D (1350 cm⁻¹) bands. In the GO spectrum (Figure 12b), the G band is expanded and shifted to 1594 cm⁻¹. The D band (1363 cm⁻¹) appears due to the intense oxidation of graphite, and the 2D band, centered at about 2700 cm⁻¹, also appears. Figure 12c demonstrates that the wide G peak and slight second-order area are features of sp¹, sp², and sp³ hybridized C-C bonds in graphene [50], while the D band (~1350 cm⁻¹) shows structural flaws such as lattice deformation in ZTA-20Nb-GO powder [51]. The Raman spectrum for the sintered composites (Figure 12d) indicates that GO was in situ reduced (rGO) during SPS. This is confirmed by the decrease in the intensity ratio between D- and G-bands (I_D/I_G). It was found that this intensity in raw powder reached 1.01, while in sintered composites, I_D/I_G diminished to 0.36, which proves the lower defectivity of the rGOs in the thermally treated samples. In addition, a well-resolved two-dimensional symmetric peak at ~2700 cm⁻¹ appears. An increase in the I_2D/I_G ratio to 0.67 compared to the raw powder mixtures (~0.13) was found in sintered composites. It also confirms the restoration of the graphene structure after sintering. The results confirmed the thermal reduction (including the reduction in large sp² regions) of the graphene oxide during sintering at 1500 °C.

Representative electron microscope images of the microstructure of sintered samples are presented in Figure 13. Darker and lighter phases are ceramics and Nb metal, respectively. The metallic particles are uniformly dispersed in the matrix and no porosity is observed. Analysis of the fracture surfaces shows that the failure mechanism was complete or partial plastic deformation and debonding of niobium particles (Figure 13B).

The evolution of crack propagation in the composite’s microstructure under the indentation load has been studied using the FIB technique (Figure 13C). These observations confirm the presence of plastic deformation and crack formation in niobium particles. SEM observations on crack paths (Figure 13D) suggest the occurrence of both crack-tip blunting and renucleation and crack bridging by niobium particles, as operating toughening mechanisms in these composites. The crack can be arrested at the Nb phase particle such that it must renucleate on the other side. Also, metallic inclusions may act as bridges and deform plastically during the period of bending. Consequently, the stress reduction at the crack tip becomes the primary source of the increased fracture toughness. The composite mechanical properties are determined by a complex balance among interactions at various levels and are largely dependent on the metal–ceramic interface adhesion. If the interfacial surface is very strong, the high degree of stress to which it is subjected will lead to the brittle fracture of the reinforcing particles; consequently, there will be no significant increase in fracture toughness. On the contrary, if the interface is very weak, the metallic particles will simply break out of the matrix, and there will also be no increase in fracture toughness. Therefore, the ceramic–metal interface must support cleavage and therefore act as a crack entrapment if the metal particle is large or has low plasticity. Various modes of fracture of niobium particles of different shapes are observed in the sintered composite (Figure 14).

On the one hand, these are metal particles that have been destroyed as a result of plastic deformation, complete or partial due to stretching (Figure 14B). In these cases, the fracture of lamellar metal particles manifests itself in the form of ductile dimpled fracture (holes and mounds on the fractured surface). On the other hand, there are niobium particles that have separated from the matrix, as well as those that have been fractured predominantly by cleavage (Figure 14A). Thus, it can be concluded that an increase in fracture toughness will occur with moderate separation of the interface surfaces, which allows metal particles to deform plastically; metal particles in the matrix exhibit elastic behavior until a crack appears. When this happens, the surface of the particle undergoes plastic deformation, but the particle itself remains intact, connected to the matrix material, even after the particle is surrounded by a crack. Subsequently, as a result of plastic deformation, the metal stretches, which eventually leads to deformation. The intense plastic deformation of the reinforcing particles before destruction eventually leads to a significant increase in fracture toughness.

The relative density of the composites was then calculated as the ratio of bulk density to theoretical density. Density values showed that all compositions can achieve almost theoretical compaction (99.8%ρ_th). The mechanical properties of the new ceramic–metal composites are presented in

Table 1. Comparison of mechanical properties of ZTA, ZTA-rGO, ZTA-Nb, and ZTA-Nb-rGO composites.

Specimen	Vickers Hardness HV (GPa)	Fracture Toughness Kc (MPa∙m1/2)	Flexural Strength σf (MPa)
ZTA [47]	16.8 ± 0.2	5.2 ± 0.3	847 ± 30
ZTA-rGO [47]	17.6 ± 0.3	6.8 ± 0.3	876 ± 43
ZTA-Nb	8.6 ± 0.4	15.3 ± 0.4	801 ± 27
ZTA-Nb-rGO	8.5 ± 0.8	16 ± 0.2	839 ± 24

and compared with ZTA and ZTA-rGO materials obtained from previous studies.

It has been published in other sources [47,52,53,54,55] that small amounts (0.1–0.5 vol.%) of graphene can improve the mechanical properties of the composite matrix. Maximum flexural strength and Vickers hardness were obtained for the zirconia–alumina–rGO composites. No significant improvement in the Vickers hardness of ZTA-Nb-rGO composites was observed. The hardness values of composites with and without rGO were almost equal. Since graphene is a soft phase, it does not contribute to an increase in hardness. XRD analyses of the polished surfaces of all studied materials revealed the predominance of t-ZrO₂ (

Table 2. Volume fractions of monoclinic and tetragonal zirconia on polished and fractured surfaces of the specimens and transformability of tetragonal zirconia.

Specimen	Volume Fractions of m- and t-ZrO2 (vol.%)
	Polished		Fractured		Transformability of t-ZrO2 Vtrans (%)
	m	t	m	t
ZTA	99	1	96	4	3
ZTA-rGO	99	1	96	4	3
ZTA-Nb	95	5	80	20	15
ZTA-Nb-rGO	95	5	80	20	15

). Additionally, the results proved that the fraction of transformed zirconia considerably increases during the fracture process (Table 2). As expected, the lowest fracture toughness values were demonstrated by samples without a reinforcing metal phase. Although, despite the same values (3%) of transformability, the ZTA-rGO composite has a value of fracture toughness 30% higher than the ZTA composite [47].

This can be explained by the fact that the presence of rGO in sintered ceramic composites creates opportunities for deflection of cracks and bridging, which, in turn, increases its fracture toughness. Composites with Nb particles exhibit a completely different behavior. The transformability of zirconia in these composites is higher than in the composites without niobium and is equal to 15%. There is always an oxide (Nb₂O₅) layer on the surface of the particles of the initial niobium powder, which, in turn, increases the ability to transform zirconium oxide. We believe that after SPS the entire oxide layer is dissolved in a solid solution. The higher transformability is probably due to the alloying effect on the tetragonality [56] of tetragonal zirconia. The EDS analysis was performed to determine the content of a solid solution of Nb in zirconium dioxide. The proportion of niobia was below 2 mol%, without any additional phases. Consequently, the presence of niobium oxide increases the transformability of zirconia and serves as an additional source of increased fracture toughness. The present study has demonstrated that the high toughness values found in ZTA-Nb and ZTA-Nb-rGO composites are based on synergies between crack bridging and plastic deformation of the metallic particles and rGO, together with crack deflection and the phase transformation of zirconia as well.

4. Conclusions

For the first time, dense ZTA-Nb and ZTA-Nb-GO composites were manufactured through a wet processing route and spark plasma sintering. The microstructure analysis showed a uniform distribution of niobium metal particles in a ceramic matrix. The proposed sintering temperature (1500 °C) leads to a thermal reduction in graphene oxide, which was confirmed by Raman spectroscopy. Simultaneous addition of 20 vol.% Nb and 0.5 vol.% GO in the ceramic matrix increases the fracture toughness to 16 MPa·m^1/2. However, it should be noted that the increase in fracture toughness values for composites with and without graphene is insignificant and amounts to only about 4%. The major contributions to toughening are the resulting crack bridging and plastic deformation of the metallic particles, together with crack deflection and interfacial debonding. In addition, the phase transformation of zirconia also leads to an increase in the fracture toughness of ceramic–metal composites. Graphene is a soft phase and does not lead to an increase in hardness. Therefore, the hardness values of the studied composites were almost identical. Based on the results of the study of the mechanical properties, we can conclude that the developed composites can be promising candidates for designing components for dental or load-bearing applications. To evaluate the possibility of using these composites for biomedical purposes, studying the mechanical properties alone is not enough. Future research directions will include the study of various methods for modifying the surface of bioinert ceramics in order to improve the process of osseointegration of the ZTA-Nb-GO composites with surrounding bone tissue. Surface modification will include various methods; for example, changing the surface relief and porosity in the intended area of contact with bone tissue or spraying with bioactive components. Also, conducting in vivo studies to assess long-term biocompatibility and osseointegration while simultaneously exposing composites to dynamic loads for more accurate prediction of clinical parameters.