The Effect of Sintering Temperature on the Phase Composition, Microstructure, and Mechanical Properties of Yttria-Stabilized Zirconia

It is known that the yttria-stabilized zirconia (YSZ) material has superior thermal, mechanical, and electrical properties. This material is used for manufacturing products and components of air heaters, hydrogen reformers, cracking furnaces, fired heaters, etc. This work is aimed at searching for the optimal sintering mode of YSZ ceramics that provides a high crack growth resistance. Beam specimens of ZrO2 ceramics doped with 6, 7, and 8 mol% Y2O3 (hereinafter: 6YSZ, 7YSZ, and 8YSZ) were prepared using a conventional sintering technique. Four sintering temperatures (1450 °C, 1500 °C, 1550 °C, and 1600 °C) were used for the 6YSZ series and two sintering temperatures (1550 °C and 1600 °C) were used for the 7YSZ and 8YSZ series. The series of sintered specimens were ground and polished to reach a good surface quality. Several mechanical tests of the materials were performed, namely, the microhardness test, fracture toughness test by the indentation method, and single-edge notch beam (SENB) test under three-point bending. Based on XRD analysis, the phase balance (percentages of tetragonal, cubic, and monoclinic ZrO2 phases) of each composition was substantiated. The morphology of the fracture surfaces of specimens after both the fracture toughness tests was studied in relation to the mechanical behavior of the specimens and the microstructure of corresponding materials. SEM-EDX analysis was used for microstructural characterization. It was found that both the yttria percentage and sintering temperature affect the mechanical behavior of the ceramics. Optimal chemical composition and sintering temperature were determined for the studied series of ceramics. The maximum transformation toughening effect was revealed for ZrO2-6 mol% Y2O3 ceramics during indentation. However, in the case of a SENB test, the maximum transformation toughening effect in the crack tip vicinity was found in ZrO2-7 mol% Y2O3 ceramics. The conditions for obtaining YSZ ceramics with high fracture toughness are discussed.


Introduction
Modern techniques of designing novel structural materials intended for applications in various high-temperature structural components are currently being developed. These techniques include fine-grained microstructure formation due to the optimization of material manufacturing modes allowing materials to reach excellent high-temperature strength and crack growth resistance as well as thermal stability [1][2][3][4][5][6][7][8][9]. calcination for 3 h at 1000 • C and subsequently underwent a heat treatment at 1300 • C for 50 h. By using XRD and Raman analyses, the authors confirmed the formation of the non-transformable (t ) ZrO 2 phase as well as the stability of this phase after heat treatment. Based on the properties of nano-CYSZ, they suggested the material as promising for advanced TBCs in aero-engine and power generation applications.
In the work [19], the effects of powder particle morphology and size on microstructure and phase composition of Y 2 O 3 -ZrO 2 polycrystals were studied. The authors used dilatometry measurements to investigate the powder's compact behavior during sintering. The SEM and EBSD studies allowed for identifying symmetry between the observed grains. Hardness, fracture toughness, and mechanical strength measurements were also performed. Two populations of grains essentially differing in their sizes were found. Surprisingly, the EDS line scan of the bigger grains displayed substantially higher yttrium content than in the much smaller grains surrounding them. The X-ray diffraction of the material revealed the presence of 46.6% t-ZrO 2 phase, 15.6% c-ZrO 2 phase, and 37.8% m-ZrO 2 phase. Using EBSD analysis, the authors tried to attribute corresponding symmetry to the grains observed in the specimen microstructure. They stated that the preferential matter transport from nanometric Y 2 O 3 -ZrO 2 particles towards sub-micrometer particles led to the transformation of the latter to form a higher-symmetry part of the system. They also suggested two mechanisms related to this phenomenon. One resulted from the yttrium concentration gradient, but it would lead to the chemical homogenization of the system. The second mechanism was related to the high curvature of the contact points between small and larger grains. This leads to the matter diffusion of smaller grains toward larger ones. These grains, coming initially from sub-micrometric monoclinic particles, become sufficiently rich in yttrium to develop t-ZrO 2 phase and c-ZrO 2 phase symmetry. Simultaneously, nanometric particles initially rich in yttrium transfer to the part of the microstructure featuring monoclinic symmetry. The described phenomenon does not allow for the chemical homogenization of the system.
The authors [20] studied the microstructure evolution of two ZrO 2 -SiO 2 nanocrystalline glass-ceramics (NCGCs) in relation to thermal treatment modes. NCGCs were composed of monoclinic (m) and tetragonal (t) ZrO 2 nanocrystallites and an amorphous SiO 2 compound. During thermal treatment, both m-ZrO 2 and t-ZrO 2 nanocrystallites were metastable. The metastability of m-ZrO 2 and t-ZrO 2 nanocrystallites was explained using a size-driven phase transformation approach. It was shown that the percentage of m-ZrO 2 in the undoped ZrO 2 -SiO 2 NCGC increased due to thermal treatment at 850 • C for 5 h and decreased due to thermal treatment at a temperature above 950 • C. A rapid phase transformation of t-ZrO 2 nanocrystallites was reached due to thermal treatment at 1250 • C for 5 h followed by cooling, with the formation of 88.6 vol% m-ZrO 2 . In contrast, the yttria additive was the reason for improved t-ZrO 2 phase stability to a temperature of 1250 • C, since the percentage of m-ZrO 2 in the yttria-doped ZrO 2 -SiO 2 NCGC continuously decreased with an increase in a temperature of thermal treatment up to 1250 • C. The sizes of both m-ZrO 2 and t-ZrO 2 nanocrystallites increased with an increase in the temperature of thermal treatment for both the NCGCs.
In the work [21], a comparative study on densification and microstructural evolution of 8 mol% YSZ sintered ceramics reinforced with CeO 2 particles (10, 12, and 14 wt.% CeO 2 ) has been performed. The specimens were fabricated via both microwave and conventional sintering methods. In both cases, the sintering temperature was 1400 • C, and the holding time was 20 min and 5 h for the microwave and conventional methods, respectively. The materials were characterized in terms of densification, microstructure, and mechanical behavior. For both methods, the sintered densities of 8YSZ specimens increased with the addition of CeO 2 amount. In these conditions, no destabilization of the 8YSZ cubic crystal structure was found. It was revealed that the grain size of the 8YSZ specimens decreased with the addition of CeO 2 . Respectively, Vickers hardness of the ceramics increased with increasing CeO 2 amount. All the mentioned effects were found to be more pronounced in microwave sintered specimens compared to those obtained by the conventional method.
The authors of the works [22,23] studied the phase changes in plasma sprayed YSZ coatings during annealing. It was found using neutron scattering and XRD studies that the t-ZrO 2 phase decomposed into the m-ZrO 2 phase and c-ZrO 2 phase while the yttria amount in the t-ZrO 2 phase decreased. Using XRD analysis, the authors of the works [24,25] revealed that in the plasma-sprayed and EB-PVD coatings under study, the transformation of the t-ZrO 2 phase into a yttria-depleted t-ZrO 2 phase and a c-ZrO 2 phase or t-ZrO 2 phase with high yttria content occurred in a temperature range of 1300-1400 • C. Transmission electron microscopy (TEM) along with XRD analysis were used in the work [26] for studying the phase transformations in EB-PVD coatings heat-treated in a temperature range of 1100-1500 • C. It was found that the t-ZrO 2 phase decomposed into an yttria-depleted t-ZrO 2 phase and both the t-ZrO 2 phase and c-ZrO 2 phases with high Y 2 O 3 content. It was also revealed that the domain boundaries having a cubic-like structure contain a quite large number of yttrium ions [27,28]. Such knowledge allows for increasing the energy efficiency of high-temperature fuel cells by optimizing the operating modes and microstructure of ceramic electrodes [29][30][31][32]. This urgent task is caused by the deployment of renewables to meet global climate objectives [33,34].
It is critically important to manufacture materials resistive to the aggressive operating environment [35][36][37][38] to avoid microstructure degradation [39][40][41][42]. Therefore, along with strength and tribology tests of materials as the most popular methods for diagnosing their load-bearing capacity, the indentation test, known as the simplest mechanical method [39,43], is widely used. This method is more microstructurally sensitive as compared to the above-mentioned ones and allows for estimating the crack growth resistance of materials [40,44]. Fracture toughness tests employing various specimen shapes and loading schemes are also quite microstructurally sensitive [40,44,45]. Therefore, the application of microhardness and crack growth resistance test methods for diagnostics of the microstructure stability of YSZ ceramics is promising in terms of searching for the optimal sintering and treatment modes.
This work is aimed at evaluating the effect of sintering temperature on the phase composition, the size and morphology of the crystallites, and the mechanical properties of YSZ ceramics stabilized by the various amount of yttria.

Experimental Procedures
In this work, yttria-stabilized zirconia (YSZ) ceramics sintered from commercial starting powders have been studied. The powders were produced at the Vol'nogorskii Mining and Smelting Plant, Vol'nogorsk, Ukraine. Initial particle sizes of the starting powders were as follows: 100-150 nm (ZrO 2 powder) and 10-30 nm (Y 2 O 3 powder). A series of beam specimens of YSZ ceramics stabilized with 6, 7, and 8 mol% Y 2 O 3 (hereinafter: 6YSZ, 7YSZ, and 8YSZ) approximately 4.2 × 4.2 × 50 mm in size were sintered in a furnace for 2 h. Argon was used as an inert sintering atmosphere. Eight variants of material were obtained using four sintering temperatures (1450 • C, 1500 • C, 1550 • C, and 1600 • C) for the 6YSZ series and two sintering temperatures (1550 • C and 1600 • C) for the 7YSZ and 8YSZ series (Table 1). For marking each variant, corresponding chemical composition and sintering temperature were indicated, e.g., 7YSZ-1550. After sintering, the side surfaces of specimens were polished using a grinding and polishing machine for metallographic preparation to reach the required surface quality and avoid phase transformations.
Microhardness of the material variants was measured using a NOVOTEST TC-MKB1 microhardness tester (Novotest, Novomoskovsk, Ukraine). We used the following set of indentation loads: 0.49 N, 0.98 N, 1.96 N, 2.94 N, 4.91 N, and 9.81 N. At least 10 indentations for each level of the indentation load were made to determine the microhardness of each material variant.
The relevant standards [46,47] regulate the microhardness measurement conditions. Vickers microhardness {in GPa} is calculated by the formula [47]: where P is the indentation load {N} and d is the average length of the diagonals of the indentation imprint {mm}. An optical microscope Neophot-21 (Zeiss, Oberkochen, Germany) was used for estimating the imprint and crack geometry.
Along with microhardness, the fracture toughness of material was estimated by to calculating the critical stress intensity factor (SIF), K Ic . This characteristic made it possible to characterize the propensity of a material to brittle fracture due to the nucleation and propagation of cracks [48][49][50]. There exists a wide range of methods for estimating the fracture toughness of materials under Vickers pyramid indentation [42,51,52]. In these works, the formulas for calculating the K Ic values contain both physical and mechanical parameters, as well as empirical coefficients. Due to the comparison of the K Ic values calculated by these formulas with those obtained by conventional methods of fracture mechanics, we recently concluded [53] that the following formula presented by the authors of the work [51] best fits the characterization of the ZrO 2 -Y 2 O 3 ceramics: where E is Young's modulus {GPa}, H is microhardness {GPa}, P is the indentation load {N}, and c is the radial crack length {m}. Therefore, we used this formula to estimate the fracture toughness of the materials under study.
For comparison, we used in this work a single-edge notch beam (SENB) test [54][55][56] to estimate the fracture toughness of the material. An edge notch less than 0.1 mm in width was machined in a beam specimen. The three-point bending SENB tests were carried out on the MTS Criterion E43.104 test machine (MTS Systems Corporation, Eden Prairie, MN, USA) at 20 • C in air. The distance between the supporting rollers of the loading unit was 14 mm. For calculating the critical SIF K Ic of material, we used corresponding formulas [54][55][56]. The average K Ic value was computed for each set of five specimens of the investigated material variants.
The material microstructure and morphology of the fracture surfaces of tested specimens were investigated using a Carl Zeiss EVO-40XVP scanning electron microscope (SEM) (Zeiss, Oberkochen, Germany). The chemical homogeneity of materials was evaluated with an energy-dispersive X-ray (EDX) microanalysis using an INCA Energy 350 system (Oxford Instruments, Abingdon, UK). A DRON-4.07M diffractometer (Bourevestnik, St. Petersburg, Russia) was used to perform X-ray diffraction (XRD) studies of as-sintered specimens. All procedures including indexing, refinement of the profile and structural parameters, as well as calculations/evaluation of the phase weight fractions were performed using the WinCSD program package (WinCSD, https://www.wincsd.eu/, accessed on 20 March 2022). The ZrO 2 phase marking and reference codes were as follows: t-tetragonal (COD ID 2300612), m-monoclinic (COD ID 1528984), and c-cubic (COD ID 2101234).

The XRD Analysis of YSZ Ceramics
The obtained XRD patterns of the materials under study (Table 1) exhibit, in general, the phase balance for 6YSZ, 7YSZ, and 8YSZ ceramics ( Figure 1). We revealed ambiguous changes in the phase balance with changes in the sintering temperature from 1450 • C to 1600 • C for 6YSZ ceramics. A maximum percentage of t-ZrO 2 (over 56 wt.%) along with a decrease in m-ZrO 2 and c-ZrO 2 weight fractions (to 32 wt.% and 11 wt.%, respectively) was revealed for this material sintered at a temperature of 1550 • C.
ies of as-sintered specimens. All procedures including indexing, refinement of the profile and structural parameters, as well as calculations/evaluation of the phase weight fractions were performed using the WinCSD program package (WinCSD, https://www.wincsd.eu/, accessed on 20 March 2022). The ZrO2 phase marking and reference codes were as follows: t-tetragonal (COD ID 2300612), m-monoclinic (COD ID 1528984), and c-cubic (COD ID 2101234).

The XRD Analysis of YSZ Ceramics
The obtained XRD patterns of the materials under study (Table 1) exhibit, in general, the phase balance for 6YSZ, 7YSZ, and 8YSZ ceramics ( Figure 1). We revealed ambiguous changes in the phase balance with changes in the sintering temperature from 1450 °C to 1600 °C for 6YSZ ceramics. A maximum percentage of t-ZrO2 (over 56 wt.%) along with a decrease in m-ZrO2 and c-ZrO2 weight fractions (to 32 wt.% and 11 wt.%, respectively) was revealed for this material sintered at a temperature of 1550 °C.  Table 1). Phase marking: t-tetragonal, m-monoclinic, c-cubic.
Thus, in 6YSZ ceramics, the sintering temperature of 1550 °C allows for providing a relatively high percentage of the t-ZrO2 phase, while the m-ZrO2 phase weight fraction is decreased compared to variants 6YSZ-1450 and 6YSZ-1500. Similarly, the c-ZrO2 phase weight fraction reaches its minimum for variant 6YSZ-1550. In 7YSZ ceramics at the sintering temperature of 1550 °C, a relatively high percentage of the m-ZrO2 phase (over 55 wt.%) was obtained, while the t-ZrO2 phase weight fraction (about 43 wt.%) obtained was lower than for both variants 6YSZ-1550 and 8YSZ-1550.
Thus, in 6YSZ ceramics, the sintering temperature of 1550 • C allows for providing a relatively high percentage of the t-ZrO 2 phase, while the m-ZrO 2 phase weight fraction is decreased compared to variants 6YSZ-1450 and 6YSZ-1500. Similarly, the c-ZrO 2 phase weight fraction reaches its minimum for variant 6YSZ-1550. In 7YSZ ceramics at the sintering temperature of 1550 • C, a relatively high percentage of the m-ZrO 2 phase (over 55 wt.%) was obtained, while the t-ZrO 2 phase weight fraction (about 43 wt.%) obtained was lower than for both variants 6YSZ-1550 and 8YSZ-1550.
The XRD patterns of the selected material variants ( Figure 2, variants 6YSZ-1550, 6YSZ-1600, 7YSZ-1600, and 8YSZ-1600) show in detail the above-mentioned peculiarities of the phase balance of the studied compositions.
The Therefore, we can describe the general tendencies of changes in 6YSZ, 7YSZ, and 8YSZ ceramics phase compositions as follows: (i) with increasing sintering temperature, the content of the tetragonal phase increases when the percentage of the stabilizing Y2O3 additive is quite low (6YSZ ceramics); (ii) the sintering temperature of 1550 °C is critical in 6YSZ, 7YSZ, and 8YSZ ceramics since the content of the tetragonal phase decreases and content of the monoclinic phase increases with a further increase in sintering temperature; (iii) the amount of cubic phase is quite low, especially in 7YSZ and 8YSZ ceramics, so the cubic phase is the balance; (iv) the maximum m-ZrO2 phase percentage was found in variant 7YSZ-1600 as a result of the decrease in the t-ZrO2 phase weight fraction and increase in the c-ZrO2 phase weight fraction.  Table 1). Phase marking: t-tetragonal, m-monoclinic, c-cubic. Therefore, we can describe the general tendencies of changes in 6YSZ, 7YSZ, and 8YSZ ceramics phase compositions as follows: (i) with increasing sintering temperature, the content of the tetragonal phase increases when the percentage of the stabilizing Y 2 O 3 additive is quite low (6YSZ ceramics); (ii) the sintering temperature of 1550 • C is critical in 6YSZ, 7YSZ, and 8YSZ ceramics since the content of the tetragonal phase decreases and content of the monoclinic phase increases with a further increase in sintering temperature; (iii) the amount of cubic phase is quite low, especially in 7YSZ and 8YSZ ceramics, so the cubic phase is the balance; (iv) the maximum m-ZrO 2 phase percentage was found in variant 7YSZ-1600 as a result of the decrease in the t-ZrO 2 phase weight fraction and increase in the c-ZrO 2 phase weight fraction.
In general, the phase balance in the ceramics under study reflects the competing effect of two factors; namely, the sintering temperature of the ceramics and the content of the stabilizing Y 2 O 3 additive. Therefore, to achieve a relatively high content of metastable tetragonal phase, all the ceramics investigated (6YSZ, 7YSZ, and 8YSZ) should be sintered at a temperature of 1550 • C. Under these conditions, a minimum amount of cubic phase is formed.

Mechanical Properties of YSZ Ceramics and the Relations to Their Microstructure
In the work [53], a dependence of the microhardness of yttria-stabilized zirconia (ZrO 2 -8 mol% Y 2 O 3 ) on the indentation load, known as the indentation size effect [57], was revealed. For this material, the average values of microhardness decreased with increasing indentation load from 0.49 N to 9.81 N. Additionally, a tendency was observed with the yield of microhardness values on the plateau at indentation loads in a range of 4.91 N to 9.81 N. It was concluded that the values of fracture toughness and microhardness obtained in this range of indentation loads are invariant.
In our work, the dependences of microhardness on the indentation loads of 0.49 N, 0.98 N, 1.96 N, 2.94 N, 4.91 N, and 9.81 N for the material variants 1-8 were obtained (Figure 3). It was found that the material variants 7YSZ-1600 and 8YSZ-1600 are characterized by decreasing the average values of microhardness with increasing indentation load. The phase balance of these ceramics is characterized by the maximum percentages of the monoclinic phase. However, for other material variants, we can observe the opposite tendency ( Figure 3).
In the work [53], a dependence of the microhardness of yttria-stabilized zirconia (ZrO2-8 mol% Y2O3) on the indentation load, known as the indentation size effect [57], was revealed. For this material, the average values of microhardness decreased with increasing indentation load from 0.49 N to 9.81 N. Additionally, a tendency was observed with the yield of microhardness values on the plateau at indentation loads in a range of 4.91 N to 9.81 N. It was concluded that the values of fracture toughness and microhardness obtained in this range of indentation loads are invariant.
In our work, the dependences of microhardness on the indentation loads of 0.49 N, 0.98 N, 1.96 N, 2.94 N, 4.91 N, and 9.81 N for the material variants 1-8 were obtained (Figure 3). It was found that the material variants 7YSZ-1600 and 8YSZ-1600 are characterized by decreasing the average values of microhardness with increasing indentation load. The phase balance of these ceramics is characterized by the maximum percentages of the monoclinic phase. However, for other material variants, we can observe the opposite tendency (Figure 3).  Table 1) correspond to the symbol numbers.  Table 1) correspond to the symbol numbers.
In general, the yield of average microhardness values on the plateau at higher indentation loads was found for all the material variants.
The invariant values of the material microhardness obtained under the indentation load of 9.81 N were taken to construct a graph for studying the evolution of changes in microhardness of the studied ceramic variants with a change in the sintering temperature from 1450 • C to 1600 • C (Figure 4). It was revealed that, generally, the increase of sintering temperature from 1450 • C to 1500 • C leads to the improvement of mechanical properties of 6YSZ ceramics. In particular, an increase in microhardness (by 5-6%, Figure 4) was observed for this material and the same tendency in fracture toughness (by 3-4%, Figure 5a) was found. The levels of these characteristics remain unchanged while increasing the sintering temperature up to 1550 • C.
The increase in sintering temperature up to 1600 • C leads to intensive grain growth in YSZ ceramics. This, in turn, leads to the suppression of retention of the metastable tetragonal zirconia [58] in the case when the average grain size of the t-ZrO 2 phase is larger than the admissible one. It seems that the last is about 1 µm for ceramics of this type. According to [58], the microhardness of m-ZrO 2 is lower than t-ZrO 2 . During indentation, the t-m transition occurs with the formation of m-ZrO 2 which causes the lowering of microhardness. dentation loads was found for all the material variants.
The invariant values of the material microhardness obtained under the indentation load of 9.81 N were taken to construct a graph for studying the evolution of changes in microhardness of the studied ceramic variants with a change in the sintering temperature from 1450 °C to 1600 °C (Figure 4). It was revealed that, generally, the increase of sintering temperature from 1450 °C to 1500 °C leads to the improvement of mechanical properties of 6YSZ ceramics. In particular, an increase in microhardness (by 5-6%, Figure 4) was observed for this material and the same tendency in fracture toughness (by 3-4%, Figure 5a) was found. The levels of these characteristics remain unchanged while increasing the sintering temperature up to 1550 °C.  Table 1). The microhardness was measured under the indentation load of 9.81 N. Figure 5. Changes in fracture toughness of 6YSZ, 7YSZ, and 8YSZ ceramics depending on the sintering temperature (see Table 1 Table 1). The microhardness was measured under the indentation load of 9.81 N. dentation loads was found for all the material variants.
The invariant values of the material microhardness obtained under the indentation load of 9.81 N were taken to construct a graph for studying the evolution of changes in microhardness of the studied ceramic variants with a change in the sintering temperature from 1450 °C to 1600 °C (Figure 4). It was revealed that, generally, the increase of sintering temperature from 1450 °C to 1500 °C leads to the improvement of mechanical properties of 6YSZ ceramics. In particular, an increase in microhardness (by 5-6%, Figure 4) was observed for this material and the same tendency in fracture toughness (by 3-4%, Figure 5a) was found. The levels of these characteristics remain unchanged while increasing the sintering temperature up to 1550 °C.  Table 1). The microhardness was measured under the indentation load of 9.81 N. Figure 5. Changes in fracture toughness of 6YSZ, 7YSZ, and 8YSZ ceramics depending on the sintering temperature (see Table 1 Table 1 In contrast to 6YSZ ceramics, a substantial increase in the m-ZrO 2 weight fraction for 7YSZ and 8YSZ ceramics was revealed while increasing the sintering temperature from 1550 • C to 1600 • C (Figure 1). This, in turn, leads to a decrease in the t-ZrO 2 or c-ZrO 2 weight fractions in these ceramics.
A common result of the two above-mentioned processes, namely, suppression of the t-m transformation of ZrO 2 with the sintering temperature increase and the stress-induced formation of m-ZrO 2 , is displayed in both Figures 4 and 5a. For the case of the sintering temperature of 1550 • C, the differences between the average values of both microhardness ( Figure 4) and fracture toughness (by the Vickers indentation method, Figure 5a) for 7YSZ and 8YSZ ceramics are observed. The latter has an advantage over the previous one both in terms of microhardness and fracture toughness. Probably, a high percentage of m-ZrO 2 in 7YSZ ceramics (Figure 1b) causes the lowering of microhardness, whereas a slight effect of the t-m transformation occurs at a comparatively low percentage of t-ZrO 2 . In contrast, the microhardness of 8YSZ ceramics is higher because of a lower percentage of m-ZrO 2 (Figure 1c). On the other hand, since the t-ZrO 2 percentage is higher in 8YSZ ceramics (Figure 1c), the t-m transformation has a significant effect on microhardness by lowering it and, oppositely, causes an increase in fracture toughness of the material (Figures 4 and 5a, respectively).
On the contrary, another pattern of calculated critical SIF K Ic values using the SENB method was obtained (Figure 5b). Significantly higher fracture toughness of 7YSZ ceramics as compared to 6YSZ and 8YSZ (by 15% and 46%, respectively) was found. This ambiguous behavior of the material is evidence showing that the sintering temperature of about 1550 • C is critical in the microstructure formation process. Such an ambiguity during the estimation of fracture toughness of material by two different methods showed that the t-m transition that occurred in the crack tip vicinity of a notched beam specimen is more pronounced than in the case of the Vickers pyramid indentation. In contrast, no appropriate conditions were available for enhancing the fracture toughness of 8YSZ ceramics, and the reason for that was a lower percentage of m-ZrO 2 and, especially, c-ZrO 2 .
For material sintered at 1600 • C, the t-m transformation dynamics are less pronounced. In this case, mainly t-ZrO 2 by its strength, without the contribution from the t-m transformation, provides the achieved level of crack growth resistance (Figure 5a) which is reflected in the close values of fracture toughness of 6YSZ, 7YSZ, and 8YSZ ceramics at such same close values of microhardness for these material variants.
Thus, for 6YSZ, 7YSZ, and 8YSZ ceramics, phase balances were defined at which the maximum fracture toughness is reached. The maximum K Ic values using both the SENB and Vickers indentation methods for 6YSZ ceramics were obtained at the maximum content of the tetragonal phase (variant 6YSZ-1550), whereas for ceramics with a higher content of the stabilizing Y 2 O 3 additive (7YSZ and 8YSZ ceramics), the maximum K Ic levels were reached at the maximum content of the monoclinic phase (variants 7YSZ-1600 and 8YSZ-1600, respectively).
Low-magnification SEM images of the microstructure of the material variants were used for performing local and general EDX analyses ( Figure 6). The obtained EDX data sets were ordered in such a fashion ( Table 2) that allows for analyzing the availability of the main chemical elements (oxygen, yttrium, and zirconium) in local areas (dark-gray area, spectrum 1; light-gray area, spectrum 2) as well as in general (spectrum 3).
Only variant 7YSZ-1600 exhibits a microstructure which comprises the pure monoclinic phase (see Table 2, spectrum 1 for oxygen, yttrium, and zirconium). A minimum of yttrium percentage is evidence of this assumption. Additionally, we can clearly observe a tendency to increase yttrium content for variants 7YSZ-1550 and 8YSZ-1550, whereas its percentage decreases for variants 6YSZ-1450, 6YSZ-1500, 6YSZ-1550, 6YSZ-1600, and 8YSZ-1600. According to the literature data [19], two mechanisms related to the phenomenon of the yttrium concentration gradient were suggested. One is the concentration gradient-driven mechanism leading to chemical homogenization. The second mechanism relates to the contact point's geometry which leads to the yttrium diffusion from smaller grains toward larger ones. Therefore, large grains become sufficiently rich in yttrium to form t-ZrO 2 and c-ZrO 2 phases while small grains initially rich in yttrium become depleted. Such a phenomenon causes the chemical inhomogeneity of ceramics.
Based on the above assumptions, we were able to identify the m-ZrO 2 phase areas and analyze their morphology at higher magnification (Figure 7). Only variant 7YSZ-1600 exhibits a microstructure which comprises the pure monoclinic phase (see Table 2, spectrum 1 for oxygen, yttrium, and zirconium). A minimum of yttrium percentage is evidence of this assumption. Additionally, we can clearly observe a tendency to increase yttrium content for variants 7YSZ-1550 and 8YSZ-1550, whereas its percentage decreases for variants 6YSZ-1450, 6YSZ-1500, 6YSZ-1550, 6YSZ-1600, and 8YSZ-1600. According to the literature data [19], two mechanisms related to the phenomenon of the yttrium concentration gradient were suggested. One is the concentration gradient-driven mechanism leading to chemical homogenization. The second mechanism relates to the contact point's geometry which leads to the yttrium diffusion from smaller grains toward larger ones. Therefore, large grains become sufficiently rich in yttrium to form t-ZrO2 and c-ZrO2 phases while small grains initially rich in yttrium become depleted. Such a phenomenon causes the chemical inhomogeneity of ceramics.  (Table 2). Table 2. The percentage (wt%) of the main chemical elements (oxygen, yttrium, and zirconium) present in local areas (spectrums 1 and 2) and in general (spectrum 3) according to the local and general EDX analyses of the investigated material variants (see Figure 6). Based on the above assumptions, we were able to identify the m-ZrO2 phase areas and analyze their morphology at higher magnification ( Figure 7).  For variants 6YSZ-1450 and 6YSZ-1500 (Figure 7a,b) there are no signs of any substructurization of the m-ZrO 2 phase agglomerates, whereas for variant 6YSZ-1550 some signs of separate sub-area formation can be observed (Figure 7c), and for variant 6YSZ-1600, clear rectangle-shaped sub-areas can be seen (Figure 7d). Such observations are consistent with peculiarities of fracture surface morphology of the corresponding specimens examined after fracture toughness tests (Figure 8a-d). In particular, the chaotic micro-areas showing fracture mainly along the boundaries of the m-ZrO 2 phase agglomerates are observed for variants 6YSZ-1450 and 6YSZ-1500 (Figure 8a,b). In contrast, a relief fracture surface can be observed for variant 6YSZ-1550 (Figure 8c) which is evidence of the formation of the microstructure of fully recrystallized m-ZrO 2 phase grains about 1 µm in size. In this case, both the transgranular fracture of the larger grains and intergranular fracture along distinct boundaries of smaller grains were noted. Additionally, nanoparticles of the t-ZrO 2 phase about 20-100 nm in size can be clearly seen on the boundaries. Similar to this, a relief fracture surface showing the t-ZrO 2 phase grains of the increased size (about 150-400 nm) can be observed for variant 6YSZ-1600 (Figure 8d). Such relief fracture patterns for variants 6YSZ-1550 and 6YSZ-1600 are consistent with corresponding average values of both microhardness ( Figure 4) and fracture toughness ( Figure 5). intergranular fracture along distinct boundaries of smaller grains were noted. Additionally, nanoparticles of the t-ZrO2 phase about 20-100 nm in size can be clearly seen on the boundaries. Similar to this, a relief fracture surface showing the t-ZrO2 phase grains of the increased size (about 150-400 nm) can be observed for variant 6YSZ-1600 (Figure 8d). Such relief fracture patterns for variants 6YSZ-1550 and 6YSZ-1600 are consistent with corresponding average values of both microhardness ( Figure 4) and fracture toughness ( Figure 5). For variant 7YSZ-1550 (Figure 7e), separate round-shaped sub-areas about 300-400 nm in size, slightly different in color but without visible boundaries, were detected. It may be suggested that the mentioned sub-areas are yttrium-enriched since maximal yt- For variant 7YSZ-1550 (Figure 7e), separate round-shaped sub-areas about 300-400 nm in size, slightly different in color but without visible boundaries, were detected. It may be suggested that the mentioned sub-areas are yttrium-enriched since maximal yttrium percentage was revealed by EDX analysis in the m-ZrO 2 phase agglomerates ( Table 2). In contrast, in variant 7YSZ-1600, no signs of any sub-area formation of the m-ZrO 2 phase agglomerates can be seen (Figure 7f). The formation of such fully recrystallized grains of the pure m-ZrO 2 phase due to minimal yttrium percentage was evidenced by EDX analysis (see Table 2). The morphology of microstructural components was found to be in full accordance with fracture surface patterns of the corresponding specimens (Figure 8e,f). In particular, intergranular fractures along boundaries of the t-ZrO 2 phase nanoparticles about 200-300 nm in size as well as cleavage facets of the m-ZrO 2 phase particles can be clearly seen (Figure 8e). For variant 7YSZ-1600, both the transgranular fracture of the larger m-ZrO 2 phase particles and intergranular fractures along distinct boundaries of the t-ZrO 2 phase particles about 1 µm in size were noted (Figure 8f). These fracture mechanisms correspond to the highest level of fracture toughness determined by the SENB method (Figure 5b), whereas corresponding average values of microhardness are lower as compared to 6YSZ variants of ceramics ( Figure 4).
For variants 8YSZ-1550 and 8YSZ-1600, the average value of microhardness is intermediate among the variants of ceramics under study ( Figure 4). Simultaneously, the fracture toughness of both 8YSZ variants determined by the SENB method does not reach high values (Figure 5b). Such behaviors of these material variants can be substantiated taking into account the peculiarities of the morphology of microstructural components (Figure 7g,h) and the corresponding fracture micromechanisms (Figure 8g,h). In particular, from the point of view of yttrium distribution between the t-ZrO 2 and m-ZrO 2 phases (Table 2), a steep decrease in yttrium percentage was observed for the m-ZrO 2 phase agglomerates with a sintering temperature increase. However, this did not allow the (t-m) ZrO 2 phase transformation [58] to occur to a great extent and thus did not provide high values of fracture toughness even for variant 8YSZ-1600. The morphology of the m-ZrO 2 phase agglomerates was characterized by blurred contours and the t-ZrO 2 phase particles were quite distinct for variant 8YSZ-1550 (Figure 7g). In contrast, greatly disintegrated m-ZrO 2 phase agglomerates were seen in the microstructure of variant 8YSZ-1600 (Figure 7h). The corresponding fracture mechanism in a specimen of variant 8YSZ-1550 was intergranular due to relatively large pores (Figure 8g), whereas in a specimen of variant 8YSZ-1600, it comprised both transgranular and intergranular fracture microregions (Figure 8h).
Thus, based on the results of the microhardness test and fracture toughness test using the SENB method, it can be concluded that the best variant of material is that of 7YSZ ceramics sintered at a temperature of 1600 • C (variant 7YSZ-1600).

1.
In this work, the conditions for the formation of tetragonal, monoclinic, and cubic phases of zirconia in 6YSZ, 7YSZ, and 8YSZ ceramics have been substantiated.

2.
The dependences of phase composition and mechanical properties of the studied YSZ ceramics on the sintering temperature have been analyzed. 3.
The fracture toughness of YSZ ceramics is related to the phase transformations occurring in the material in the presence of a stabilizing additive. By comparing the mechanical behaviors of studied material variants, it was found that 7YSZ ceramics sintered at 1600 • C have the highest level of fracture toughness. Institutional Review Board Statement: Not applicable, due to the studies not involving humans or animals.
Informed Consent Statement: Not applicable due to this studies not involving humans.
Data Availability Statement: All the supporting and actual data are presented in the manuscript.