1 Introduction

Nowadays, zirconia ceramics doped with the stabilizing oxides Y2O3 [1,2,3,4], CeO2 [5,6,7], MgO [8], or CaO [9] are being widely developed. The metastable tetragonal phase of zirconia ceramics is stabilized in such a way at room temperature. It is well known that transformation toughening of zirconia occurs due to the transformation of the tetragonal (t) phase into the monoclinic (m) one [10,11,12,13,14]. This process takes place in the crack tip vicinity under significant stresses, which can effectively improve crack growth resistance of ceramics. Therefore, increased strength and crack growth resistance of the ceramics can be ensured due to the tetragonal phase.

Additionally, other oxides can be added while producing these ceramics to improve their mechanical strength and physical properties [10, 15,16,17,18,19,20,21,22,23,24,25]. Partially stabilized zirconia ceramics having improved mechanical properties can be utilized in various high-temperature structural components [26,27,28,29,30]. These ceramics are also widely used in energy generation, chemical industry, and biomedical applications [31,32,33,34,35,36,37,38,39,40,41,42].

The authors of a number of works [43,44,45] showed that the sintering temperature may affect to a great extent the grain size, porosity, and mechanical behavior of zirconia. It was reported in the work [44] that a single-phase microstructure (m-ZrO2) had been formed already after pressing the zirconia specimens without Y2O3 additive. In contrast, ceramics containing 1.5 mol% Y2O3 exhibited a single-phase microstructure (t-ZrO2) after sintering at 1100 °C. As compared to the two last variants, for zirconia containing 8 mol% Y2O3, a single-phase microstructure (c-ZrO2) was formed during pressing. The maximum pore diameter in all the compacts reached 12.5 nm after pressing. An increase in the pore diameter was found in specimens without Y2O3 additive when the temperature increased to 600 °C. Intense grain growth and lowering of porosity occurred with a further increase in sintering temperature to 1100 °C. In a similar way, specimens containing 1.5 and 8 mol% of Y2O3 additives exhibited significant pore growth with increasing the sintering temperature to 600 °C. Significant densification of the specimens, along with the slight growth of grains, was revealed with increasing temperature up to 1100 °C.

Several authors investigated the effect of conditions of zirconia ceramics preparation on the resulting grain size and mechanical behavior of prepared specimens [46,47,48,49,50]. The authors of the work [47] studied the evolution of grain growth in doped zirconia annealed at 1500 °C for several hours. The maximum grain size for 12Ti–ZrO2 ceramic was found to reach while sintering at 1400 °C with for 70 h. Other authors [48] reported grain size of about 1 μm while sintering at 1500 °C for a few hours. In the work [49], a study on the grain growth of zirconia ceramics was performed. A maximum grain size of 20 μm was found for 8 mol% Y2O3–ZrO2 ceramic after sintering at 1700 °C for 20 h. For Y2O3–ZrO2 ceramic containing 4 mol% Y2O3, a maximum grain size of more than 11 μm was found after sintering at 1800 °C for 100 h. The authors of the work [50] investigated specimens of zirconia ceramics sintered at maximum temperature 1600 °C for 2 h. It was found that the maximum mean grain size is 0.35 μm.

Several scientists applied spark plasma sintering as one of the recently developed techniques for preparing bulk zirconia-based ceramics, namely ceria stabilized zirconia [51,52,53], yttria-stabilized zirconia [54, 55], and other ceramics [56,57,58,59].

In the work [60], the influence of the zirconia percentage on the crack growth resistance and strength of Al2O3–ZrO2 ceramics was studied. The ceramics consisted of the monoclinic and tetragonal ZrO2 phases and the α-Al2O3 phase. The percentage of the t-ZrO2 phase decreased with increasing the total amount of ZrO2. The formation of fine-grained Al2O3–ZrO2 microstructure was reached due to adding 10–20% ZrO2 that allows obtaining ceramics with improved mechanical properties. It was proven in the work [61] that mechanical properties of alumina doped with zirconia and yttria-stabilized zirconia can be effectively optimized due to both the flexural strength and fracture toughness determination.

The authors of the works [62, 63] studied YSZ ceramics stabilized with 3–5 mol% Y2O3. The chemical and phase compositions of the ceramics were found to relate to their mechanical properties [64]. 5YSZ ceramics sintered at 1450 °C for 2 h exhibited the highest fracture toughness. This is consistent with a comparatively high percentage of the tetragonal ZrO2 phase [62]. Effects of the sintering mode and phase composition on mechanical behavior of YSZ ceramics doped with 6–8 mol% Y2O3 were studied in the work [63]. After sintering at 1600 °C for 2 h, 7YSZ ceramics showed the highest fracture toughness, similarly to the results published in [65].

In the work [66], correlations between the yttria percentage, phase balance, grain size distribution, morphology of the fracture surface, and strength for zirconia with addition of 3 to 8 mol% Y2O3 (YSZ ceramics) sintered at 1550 °C for 2 h in argon were analyzed. It was shown that 7YSZ ceramic containing the monoclinic and tetragonal ZrO2 phases exhibited the fine-grained microstructure and, as a result, the highest strength.

This work is aimed at studying the effects of oxide additives and sintering temperature on the microstructure, microhardness, strength, and fracture toughness of fine-grained ZrO2–Y2O3–CoO–CeO2–Fe2O3 ceramics for various applications.

2 Materials and Methods

In this work, zirconia ceramics stabilized with oxides Y2O3, CoO, CeO2, and Fe2O3 have been studied.

Three series of beam specimens of partially stabilized zirconia (PSZ) approximately 2.9 × 2.9 × 45 mm3 in size were sintered in an argon atmosphere for 2 h at temperatures of 1550 °C, 1580 °C, and 1620 °C, respectively. An electric resistance furnace was used for sintering. For marking each series, corresponding sintering temperatures were indicated, i.e., PSZ–1550, PSZ–1580, and PSZ–1620 (Table 1). To avoid phase transformations, the side surfaces of specimens after sintering were processed using a grinding and polishing machine for metallographic preparation. Such processing allowed reaching the required surface quality.

Table 1 Marking of the investigated ceramics, their chemical compositions, and sintering modes
Full size table

To measure microhardness of the studied ceramics, a NOVOTEST TC-MKB1 microhardness tester was used. At least ten indentations under each of the indentation load of 0.25 N, 0.49 N, 0.98 N, 1.96 N, 2.94 N, 4.91 N, and 9.81 N were made to determine the microhardness of the ceramics under study. The test was performed according to the relevant standards [67,68,69].

Various teams of scientists around the world have developed a large number of formulas for calculating the microhardness of brittle materials [70,71,72,73,74,75,76,77,78,79].

Vickers microhardness (in GPa) was calculated using the formula [68]:

$$H = 0.0018544\left( {\frac{P}{d^2 }} \right),$$
(1)

where P is the indentation load (N) and d is the average length of the diagonals of the indentation imprint (mm).

The geometries of imprint and crack were estimated using an optical microscope Neophot-21.

To estimate fracture toughness of the ceramics, the critical stress intensity factor (SIF), KIc, was calculated as an indicator of the propensity of ceramics to brittle fracture [62, 80]. It is known about a number of methods for evaluating the fracture toughness of ceramics and cermets using Vickers indentation [81,82,83]. The developed formulas used for the KIc value calculation include both mechanical and physical characteristics, as well as experimentally determined coefficients. Recently, it was proven in the works [63, 84] that for the characterization of the ZrO2–Y2O3 ceramics, the formula developed in [81] provides the best convergence of results with those obtained by traditional methods of fracture mechanics. This formula used in our work is as follows:

$$K_{{\text{Ic}}} = 0.016\left( \frac{E}{H} \right)^{1/2} \left( {\frac{P}{{c^{3/2} }}} \right),$$
(2)

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).

To study mechanical behavior of the ceramics, the three-point flexure test of ceramic specimens was performed in air at 20 °C with calculating fracture stresses [85, 86].

For comparison with the Vickers indentation method, a single-edge notch beam (SENB) test [87,88,89] was carried out as a conventional method of estimating the fracture toughness of the ceramics. The notch width was 0.1 mm. The SENB specimens were tested under three-point bend at 20 °C in air [90,91,92,93,94,95,96,97,98,99]. Corresponding formulas given elsewhere [87,88,89] were used for calculating the critical SIF KIc of the ceramics.

In both the cases (SENB and Vickers indentation tests), the average KIc value of five specimens for each series was calculated.

The microstructure and fracture surface analyses of material specimens were carried out using a Carl Zeiss EVO-40XVP scanning electron microscope (SEM). For an energy-dispersive X-ray (EDX) microanalysis, an INCA Energy 350 system was used. X-ray diffraction (XRD) analysis of specimens was performed on a DRON-4.07M diffractometer. The WinCSD program package was used to perform the procedures of indexing, refinement, and calculation of the zirconia phase weight percentages.

3 Results and Discussion

The phase compositions of the studied ceramics were determined based on the XRD analysis.

The XRD patterns for the studied ceramics exhibited two-phase structure with clear peaks of the t-ZrO2 and m-ZrO2 phases (Fig. 1). It was found that the peak heights of the t-ZrO2 phase slightly increased for the material sintered at a temperature of 1580 °C, compared to that sintered at 1550 °C, whereas the peak heights of the m-ZrO2 phase slightly decreased. Even more increase of peaks of the t-ZrO2 phase and lowering of peaks of the m-ZrO2 phase were observed for the material sintered at 1620 °C.

Fig. 1
3 X R D profiles plot intensity in an arbitrary unit versus 2 theta. The signal for experimental data follows a slight increasing trend between 20 and 100 angles, with several minimum and maximum peaks. The signal for deviated data is plotted below the experimental line.

XRD patterns of the studied ceramics (Table 1). Notation: t is tetragonal and m is monoclinic ZrO2

Full size image

Such differences in the peak heights for obtained XRD patterns are consistent with the phase balance in the studied ceramics (Fig. 2). For ceramics sintered at 1550 °C, tetragonal phase fraction was about 82.6%, whereas for ceramics sintered at 1580 °C and 1620 °C, the fractions were 84.6% and 85.9%, respectively. For these ceramic series, monoclinic phase fractions were about 17.4%, 15.4%, and 14.1%, respectively. Therefore, correlations between the sintering temperature for the studied materials and percentages of the t-ZrO2 and m-ZrO2 phases were observed.

Fig. 2
A dot plot of phase fraction versus series. The t and m values for P S Z 1550 are marked at 80 and 15. The t and m values for P S Z 1580 are marked at 85 and 14. The t and m values for P S Z 1620 are marked at 13 and 86. Values are estimated.

Phase balance in the studied ceramics (Table 1). Notation: t is tetragonal and m is monoclinic ZrO2

Full size image

There are lots of works on microstructure-related microhardness and strength of zirconia-based ceramics [84, 92, 99,100,101,102,103,104,105,106]. For ZrO2–8 mol% Y2O3 ceramics [63, 84], a gradual decrease in microhardness with an increase in the indentation load from 0.49 to 9.81 N was found. Such material behavior evidences the indentation size effect peculiar to ceramics [107]. However, the microhardness values for ZrO2–8 mol% Y2O3 ceramics were found to yield on the plateau at the loads in a range of 4.91–9.81 N. Such a phenomenon allows concluding that the values of microhardness and fracture toughness in this loading range are invariant.

In our work, the indentation loads of 0.25 N, 0.49 N, 0.98 N, 1.96 N, 2.94 N, 4.91 N, and 9.81 N were set and corresponding microhardness values were determined for the studied materials 1–3 (Fig. 3a).

Fig. 3
2 scatter plots with error bars plot H and K subscript I c versus P. The values are plotted for P S Zs 1550, 1580, and 1620. a. 3 plots are distributed throughout the graph between 6 and 12. b. 3 plots are distributed in an increasing manner between 2 and 8.

Changes in mechanical characteristics of the studied ceramics depending on the indentation load: a Vickers microhardness; b fracture toughness measured by the Vickers method. The symbol marking corresponds to the series marking given in Table 1

Full size image

The obtained microhardness vs load correlations are similar and can be divided into two parts. The first part of each dependence corresponds to low indentation loads (up to 0.98 N), whereas the second one corresponds to the range of 1.96–9.81 N. In this range, the above-mentioned indentation size effect is observed. In contrast, the first part of each dependence exhibits the lowered microhardness values due to the indentation imprints commensurable with the average diameter of pores on the specimen surface. The average microhardness values for all the studied materials were revealed to yield on the plateau under the loads of 4.91–9.81 N (Fig. 3a). For this load range, a trend to increase invariant values of microhardness with an increase in the sintering temperature was found.

It is known that an increase in sintering temperature leads to intensive grain growth in yttria-stabilized zirconia. When the average grain size of the t-ZrO2 phase becomes larger than the admissible one (about 1 μm for ceramics of this type), the retention of the metastable tetragonal zirconia is suppressed [108]. As the microhardness of t-ZrO2 is higher than m-ZrO2 [108], the lowering of microhardness occurs during indentation, due to the t–m transition. In our case, the t-ZrO2 phase is stabilized mainly with CeO2, in which content in the ceramics is 7 mol%. Such a CeO2 amount allows the strong stabilization of the tetragonal crystal structure to be achieved. A slight decrease in the m-ZrO2 phase fraction with an increase in sintering temperature of the ceramics (Fig. 2) is consistent with a slight increase in the microhardness of the ceramics (Fig. 3a).

To construct a graph for studying the changes in fracture toughness of the ceramics with a change in the sintering temperature from 1550 °C through 1580 °C to 1620 °C, the values of the material microhardness obtained under the indentation load in the range of 0.25–9.81 N were taken (Fig. 3b). Surprisingly, a slight difference between obtained levels of fracture toughness for the studied materials was found.

In contrast to such a little difference in values of fracture toughness measured by the Vickers method for the studied materials, distinct changes in both the flexural strength and fracture toughness measured by the SENB method under three-point bending while increasing the sintering temperature were found (Fig. 4).

Fig. 4
2 scatter plots with error bars plot flexural strength versus series. a. The plot for P S Z 1550 is marked at 550, P S Z 1580 is marked at 560, and P S Z 1620 is marked at 600. b. The plot for P S Z 1550 is marked at 5.1, P S Z 1580 is marked at 5.6, and P S Z 1620 is marked at 5.2.

Changes in mechanical characteristics of the studied ceramics (Table 1): a flexural strength; b fracture toughness measured by the SENB method under three-point bending

Full size image

It should be noted that the ceramic sintered at 1550 °C has both the lowest flexural strength and fracture toughness measured by the SENB method (Fig. 4). No significant difference in fracture toughness was found for all the studied ceramics. However, the best result was obtained for the ceramic sintered at 1580 °C (Fig. 4b). It is assumed that in such conditions, the highest bond strength between recrystallized t-ZrO2 grains is achieved. Although no distinct difference was revealed at low magnifications in the microstructure images (Fig. 5a, c, e) and corresponding images of fracture surfaces (Fig. 5b, d, f), the above assumption is confirmed by the images of microstructure (Fig. 6a, c, e) and fracture surfaces (Fig. 6b, d, f) taken at high magnifications. In particular, slightly increased t-ZrO2 grains are observed in the ceramic sintered at 1580 °C compared to those in the ceramic sintered at 1550 °C.

Fig. 5
a, c, and e have the SEM images of 3 sample series of P S Zs 1550, 1580, and 1620 at 25 micrometers. The morphology of the samples reveals agglomeration with uniform grain size. b, d, and f have SEM images of the samples at 25 micrometers. The morphology of the samples has a cracked surface.

SEM a, c, e microstructures (BSD images at low magnifications) and b, d, f fractography (SE images at low magnifications) of specimen series a, b PSZ–1550, c, d PSZ–1580, and e, f PSZ–1620 (Table 1)

Full size image
Fig. 6
a, c, and e have the SEM images of 3 sample series of P S Zs 1550, 1580, and 1620 at 5 micrometers. The morphology of the samples reveals agglomeration with uniform grain size. b, d, and f have SEM images of the samples at 5 micrometers. The morphology of the samples has several cleavage facets.

SEM a, c, e microstructures (BSD images at high magnifications) and b, d, f fractography (SE images at high magnifications) of specimen series a, b PSZ–1550, c, d PSZ–1580, and e, f PSZ–1620 (Table 1)

Full size image

Besides, no signs of transgranular fracture of the m-ZrO2 phase particles were found in the ceramic sintered at 1580 °C (Fig. 6d), whereas in that sintered at 1620 °C a comparatively large number of cleavage facets were revealed (Fig. 6f). This obviously is a reason for the lower fracture toughness of the last material series.

The weaker areas consisting of freely transformed m-ZrO2 grains adjoin the coarse particles of the m-ZrO2 phase in the ceramic sintered at 1620 °C. Such local weakening affects the fracture toughness of the material by lowering it. In contrast, no effect of the local weakening on flexural strength of this ceramic was found. Therefore, coarse particles of the m-ZrO2 phase contribute dominantly to increasing the strength of a bulk material.

4 Conclusion

In this work, a stabilizing effect of small percentages of oxides Y2O3, CoO, CeO2, and Fe2O3 on zirconia ceramics has been studied.

  1. 1.

    It was shown that the metal oxide additives made it possible to reach the stabilization of the tetragonal ZrO2 phase and manufacture ceramics with fine-grained microstructure.

  2. 2.

    The maximum transformation toughening effect in ZrO2–Y2O3–CoO–CeO2–Fe2O3 ceramics was found to correlate with features of fracture surface morphology. Compared to the ceramic sintered at 1550 °C, slightly increased grains of the tetragonal ZrO2 phase were revealed in the ceramic sintered at 1580 °C. The formed microstructure allowed implementing high-energy consuming fracture micromechanism involving tetragonal to monoclinic ZrO2 phase transformation toughening without transgranular cleavage fracture of monoclinic ZrO2 phase particles.