1 Introduction

For engineering applications, ceramics should possess a high fracture toughness to assure safety and durability. Zirconia-based ceramics, stabilized with metal oxides such as yttria, ceria, magnesia, calcia, and others, have obtained increasing attention due to their superior mechanical properties. Additives used to stabilize zirconia not only produce crack-free zirconia, but also the sintered compact can retain the high temperature phase, i.e., the tetragonal phase at room temperature, where the equilibrium phase is monoclinic [1, 2]. Among all zirconia ceramics, yttria tetragonal zirconia polycrystals (YTZPs) are the most attractive, for wide range of energy and biomedical applications, due to their unusual combination of excellent mechanical properties (i.e., high strength and high fracture toughness), as well as low thermal conductivity and biocompatibility [3,4,5,6,7,8,9,10].

The use of a second dopant has been found useful to improve certain properties of zirconia-based ceramics. Doping yttria tetragonal zirconia polycrystals (YTZPs) with CeO2 or addition of alumina and silica were found to improve thermal stability upon aging at relatively low temperatures [11,12,13,14,15]. Moreover, it was found that the ceria addition to zirconia reduces its thermal conductivity, making this ceramic an alternative candidate to be used in thermal barrier coating applications [16]. Gadolinia and ytterbia addition to yttria-doped tetragonal zirconia increased the hardness and the stability of the tetragonal phase, but at the same time decreased the thermal conductivity [17]. Doping yttria tetragonal zirconia with TiO2 promoted grain growth and produced transparent ceramics [18]. It was found also that yttria addition to Mg-PSZ using simple techniques [19] produced a material with long-term thermal stability at relatively high temperatures, a property highly demanded in sensor applications. Addition of dopants like MoO3 to India-stabilized zirconia increased the hardness, ionic conductivity and the thermal expansion [20]. La2O3 addition to ceria-based solid electrolyte enhances its ionic conductivity and thermal stability [21]. The addition of a material having a high elastic modulus to YTZP, formed composites with improved strength and toughness [22], as a result of strong fiber–matrix interface [22, 23].

Tetragonal zirconia produced by doping with ceria (Ce-TZP) is a very attractive ceramic material, which possesses high fracture toughness [12, 24]. Whenever Ce-TPZ was prepared from pure precursors, it had relatively high fracture toughness but low strength. On the other hand, a commercial grade—starting additive powder of CeCl3 containing 7% LaCl3—was found to produce sintered compacts of lower toughness and higher strength [14]. It was found in this latter case that a smaller amount of CeO2 (8 mol %) could stabilize the tetragonal phase, whereas 12 mol % ceria is needed to stabilize the material made from pure precursors. This gave an indication that doping with a trivalent rare earth oxide such as lanthanum oxide could enhance the strength of ceria-doped tetragonal zirconia polycrystals (Ce-TPZ). Further, it has been shown that doping zirconia/6 mol % CeO2 with 1.5 mol % gadolinia could stabilize the tetragonal phase [25], compared to the 12 mol % Ceria which is the minimum amount of ceria required to stabilize the tetragonal phase zirconia ceramic as mentioned above.

In this work, zirconia–ceria ceramic powders with two different compositions were prepared. The effect of gadolinia additions (at percentages ˂ 1 mol %) on the physical and mechanical properties of the sintered compacts made from these ceramic powders were studied, as well as the toughening mechanisms.

2 Experimental work

Co-precipitated powders containing 5 and 10 mol % CeO2 were prepared from 0.2 M mixed salt solutions containing the corresponding concentrations of zirconium oxychloride (Fluka Chemicals, UK, purity 99%) and cerium nitrate (Rare Earth Corporation, purity 99%). The coprecipitation was done at pH = 9, using ammonia solution for control, while stirring. The precipitate gels were dried, calcined at 600 °C, ground, wet ball-milled followed by sieving to produce ceria-doped zirconia powders. Ethanolic solutions of gadolinium nitrate (Rare Earth Corporation, purity 99%) with different concentrations, were added and mixed with each of the above powders to produce zirconia powders containing 0.85 mol % gadolinia/5 mol % ceria and 0.5 mol % gadolinia/10 mol % ceria (designated as 5CeGd and 10CeGd), respectively. The mixed powders were then dried, calcined, ground and finally wet ball-milled once again.

Rectangular samples of 35 mm length, 8 mm width, and 2 mm thick were prepared by pressing at 100 MPa. They had green densities of 43% TD. The samples were then sintered in air at temperatures from 1400–1600 °C for 2 h. The sintered densities were obtained using Archimedes method, taking into account the open porosities if it existed. X-ray diffraction analysis (Shimadzu XRD-3A diffractometer with Cu kα Ni filter) was used for quantitative phase determination [26], where the monoclinic weight and volume fractions Xm and Vm were calculated using the following Eqs. (1 and 2), from the XRD peak intensities I.

$$X_{m} = \frac{{I_{m} (\bar{1}11) + I_{m} (111)}}{{I_{m} (\bar{1}11) + I_{m} (111) + I_{t} (111)}}$$
(1)
$$\mathop \nu \nolimits_{m} = \frac{{PX_{m} }}{{1 + (P - 1)X_{m} }}$$
(2)

Here P is a constant equal to 1.333.

For microstructure examination, the samples were ground and polished using Abramin automatic grinding and polishing machine from Struers Denmark. The grinding and polishing media were diamond disks and diamond paste (29 μm, 7 μm, 2–5 μm) under efficient fluid cooling. The samples were then thermally etched in air at 1350 °C for 20 min. The microstructure examination was done using a scanning electron microscope (SEM: JSM-5400 JEOL). The hardness and fracture toughness were measured using the Vickers Indentation Fracture technique (IF) on the mirror-polished sample surfaces. The polishing has been done with great care in order to minimize surface stress and to reduce tetragonal to monoclinic phase transformation on the samples’ surfaces). Zwick Hardness Tester Germany at 300 N load was used. The Niihara's equation [27] was applied for the fracture toughness calculations,

$$K_{1C} = \, 0.0726 \, P \, C^{ - 3/2}$$

And Vickers hardness

$$Hv \, = 0.47\;{\text{P}}/{\text{a}}^{2}$$

Here P is the indentation load in Newton, a and c are the indentation half diagonal and the half crack length in meter, Hv in GPa and K1C in MPa √m.

A four-point bend test was performed using LLOYD LR10KN universal testing machine with 10/20 mm loading/supporting spans, and 0.22 mm/min cross-head speed for fracture strength determination. The bend strength has been calculated from the following equation:

$$\sigma_{f} = \, 3P_{f} L/4 \, b \, h^{2}$$

Here σf is the bend strength in MPa. L is the loading span, b is the sample width, and h is the sample thickness in meters. The cut surface was obtained by using Accutom-2 automatic cutting machine from Struers Denmark, using a thin diamond circular saw under very efficient cooling to avoid sample heating.

3 Results and discussion

3.1 Densification, phases, and microstructure

The base materials containing 5 mol % CeO2 (5Ce) and 10 mol % CeO2 (10Ce) showed low sintered densities. This was attributed to the small amount of CeO2 dopant, which could not stabilize the tetragonal phase completely. The densities obtained for these samples were corresponding to that of the monoclinic zirconia. The samples were weak and developed cracks, which resulted from the tetragonal to monoclinic phase transformation upon cooling from the sintering temperature. On the other hand, the gadolinia-doped samples, 5CeGd and 10CeGd, showed high densities upon sintering, even at the lowest sintering temperature (1400 °C/2 h), as shown in Table 1. Sintering at higher temperatures, 1500 and 1600 °C, gave almost theoretical densities for both compositions.

Table 1 The densities in % TD of 5CeGd and 10CeGd samples, sintered at different temperatures
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Figure 1 a, b shows the X-ray diffraction patterns for the as-sintered sample surface of 5CeGd and 10CeGd respectively, fired at 1400 °C. The XRD made on the surfaces of the samples sintered at a higher temperature, 1600 °C are shown in Fig. 2 a, b. It can be seen that the tetragonal phase is dominant for the 10CeGd sample at both sintering temperatures (Figs.1b, 2b), while a small amount of the monoclinic phase appeared in the pattern for the 5CeGd sample, as shown in Figs.1a, 2a. It should be noted here that no cubic phase was detected in the XRD patterns, in the 2θ range from 73 to 75. Only the tetragonal phase (004) and (400) lines existed. Even for the sample sintered at 1600 C where there is a possibility of cubic phase formation, no cubic phase was detected but a tetragonal phase prevailed.

Fig. 1
figure 1

XRD patterns made on the as-sintered sample surfaces fired at 1400 °C/2 h: a 5CeGd, b 10CeGd

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Fig. 2
figure 2

XRD patterns made on the as-sintered sample surfaces fired at1600 °C/2 h: a 5CeGd, b 10CeGd

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Urabe et al. [25] proposed a phase diagram for ceria, gadolinia, and zirconia, by detecting the different phases present in the samples fired at 1600 °C for 20 h. Their results showed that the material containing 1 mol % Gd2O3 and 3 mol % CeO2 had 96.4% of a tetragonal phase, while in the present work, the sintered samples containing 0.85 mol % Gd2O3 and 5 mol % CeO2 showed almost completely tetragonal phase upon firing at 1600 °C for 2 h. The slight difference is attributed to the variation in dopant concentration and the preparation method as well as the sintering time. The materials in [25] were prepared by co-precipitation from the homogeneous salt solutions of gadolinium, cerium, and zirconium in one step, while ours were made by adding an ethanolic solution of gadolinium nitrate to the mixed oxides of ceria/zirconia powder—previously prepared by co-precipitation.

Concerning the microstructure, fine-grained ceramics were obtained upon sintering at 1500 oC as shown in Fig. 3 a, b with 2 μm and 1.5 μm average grain sizes for the 5CeGd and 10CeGd materials, respectively. These values are far below the 4 μm obtained by Urabe et al. [25] when sintering at 1600 °C for 20 h. This indicates that the preparation technique used in the present work produced materials with smaller grain size, and consequently greater phase stability [28, 29].

Fig. 3
figure 3

SEM micrographs for the samples sintered at 1500 °C/2 h: a 5CeGd, b 10CeGd

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3.2 Mechanical properties

The four points bend strength, fracture toughness and Vickers hardness results are given in Table 2, for the samples sintered at different temperatures. The values of the bend strength show that the gadolinia addition in small amounts enhanced the strength of the originally weak 5Ce and 10Ce. The bend strength values were found to be higher than those obtained for the 12 mol % ceria-doped zirconia [12, 24]. However, it can be seen that both bend strength and Vickers hardness, decreased with sintering temperatures in agreement with previously published data [30, 31]. This might be due to grain size effect. It should also be noted that the hardness decreased slightly with the Vickers indentation loads.

Table 2 The four-point bend strength (σ), fracture toughness (kIC), and Vickers hardness (Hv) results
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Further, some of the samples prepared by pressureless sintering at 1400 °C were cut to study the effect of cutting on the bend strength. It was found that, when the cut sample surface was subjected to the tensile stress during the four-point bend test, higher strength values were obtained, compared to the case when the ground surfaces were subjected to the tensile stress. Similar results have been obtained for yttria-doped zirconia (TZ3YA) [32].

Figure 4 a, b shows the XRD patterns made on the ground surfaces of 5CeGd and 10CeGd samples sintered at 1400 °C. It can be seen that grinding induced some tetragonal to monoclinic phase transformation. However, the tetragonal peak (111) does not show any anomaly or shoulder corresponding to a cubic phase; in contrast to what has been reported recently for zirconia composites doped with Nd2O3 and Y2O3 [33]. Figure 5 a, b shows the XRD patterns made on the cut surfaces of 5CeGd and 10CeGd sintered at 1400 °C. The presence of the monoclinic phase indicated that cutting induced tetragonal to monoclinic phase transformation, same as in grinding and even greater. In addition, it can be seen that the intensity of the (002) line increases while that of the (200) line decreases. Table 3 summarizes the XRD analysis from which it can be seen that grinding induced phase transformation while cutting induced phase transformation as well as domain reorientation. The domain reorientation (domain switching) is the switching of the intensities of XRD lines (002) and (200). This has been previously observed on the cut surface of some tetragonal zirconia ceramics [34, 35].

Fig. 4
figure 4

XRD patterns made on the ground sample surfaces fired at 1400 °C/2 h: a 5CeGd, b 10CeGd

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Fig. 5
figure 5

XRD patterns made on cut sample surfaces fired at 1400 °C/2 h, a 5CeGd, b 10CeGd

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Table 3 XRD analysis for the samples sintered at 1400 °C/2 h with different surface status
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The above-mentioned XRD results and analysis can explain the higher strength values obtained for the as cut samples compared with those obtained for the as ground samples. The higher strength values might be attributed to the larger portion of constrained expansion volume resulting from the phase transformation on the as cut rather than on the as ground surfaces.

Figure 6 shows the Weibull plot for the bend strengths of the samples 10CeGd sintered at 1400 °C for 2 h, where the side of each sample exposed to the tensile stress was the cut surface. From the figure, a mean bend strength of 795.8 MPa and a high Weibull modulus equal to 34 were obtained.

Fig. 6
figure 6

The Weibull plot (semi-log plot) for the bend strengths of the samples 10CeGd sintered at 1400 °C/2 h

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3.3 Toughening mechanisms

3.3.1 Transformation toughening

Transformation toughening is the increase in resistance to crack propagation, which results from the martensitic tetragonal to monoclinic phase transition. This occurs ahead of the crack tip in the process zone and is characterized by microcrack formation. From Tables 2 and 3, an increase in the bend strength with the formation of monoclinic phase on the as cut sample surfaces can be observed. Also, the fracture toughness (the stress intensity factor KIC) for the sample 5CeGd is greater than that for the sample 10CeGd; where the cut and the ground surfaces show more monoclinic content in the former than in the later. The stresses applied by grinding and cutting trigger the tetragonal to monoclinic phase transformation on both the ground and cut sample surfaces. This phase transformation, as mentioned above, is accompanied by volume expansion in these sample surfaces, putting them under compression which creates a stress field acting in opposition to the stress field that promotes the propagation of the crack. Consequently, the bend stress increases to exceed these compression stresses.

An enhancement in toughness is obtained because the energy associated with crack propagation is dissipated both in the tetragonal/monoclinic transformation and in overcoming the compression stresses due to the volume expansion. The applied stress intensity factor KIC surpasses the real intensity factor at the crack tip KItip, by a KIsh—stress intensity factor—due to crack tip shielding by the transformation in the process zone. This in fact acts to close the crack or to resist its opening and propagation. The KIsh is proportional to the applied KIC, and the amount of phase transformation ahead of the crack tip which corresponds to the amount of transformable tetragonal phase present in the material. The following equation of Chevalier et al. [36], expresses the relations between KIC, KItip and KIsh.

$$K_{Itip} = K_{IC} - K_{Ish}$$

where

$$\begin{gathered} K_{Ish} = C_{sh} K_{IC} \hfill \\ C_{sh} = \frac{{0.214E \, V_{f} \, e^{T} (1 + \nu )}}{{(1 - \nu )\sigma_{m}^{c} }}\frac{\sqrt 3 }{{12\pi }} \hfill \\ \end{gathered}$$

E is the Young's modulus (GPa), Vf the volume fraction of transformable particles, eT (T in K°) is the volume of dilatation associated with the transformation, ν is the Poisson's ratio, and σCm (MPa) is the critical local stress leading to phase transformation. This equation could be made use of qualitatively rather than quantitatively because of the difficulties encountered in determining the parameters.

From the above equation, it can be seen clearly that the transformation toughening, is a function of the volume fraction of the transformable tetragonal phase present in the samples; rather than the whole tetragonal phase present. This is in agreement with previously reported results by Lin et al. [37] which related the increase in fracture toughness to the increase in transformability or operation relevant toughening mechanisms. The increase in bend strength obtained for the samples tested when the cut surface was the one subjected to tension is the result of the high compressive stress induced by cutting. This in fact made the sample to fracture at a higher stress which in turn induced more tetragonal to monoclinic phase transformation. Figure 7 shows 50% monoclinic phase on the fracture surface for the sample 5CeGd sintered at 1400 °C.

Fig. 7
figure 7

The XRD made on the fracture surface of the cut 5CeGd sample sintered at 1400 °C

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3.3.2 Domain switching toughening

In principle, both the domain switching and transformation toughening can occur simultaneously during crack propagation. It can be seen from Fig. 5 a, b for the XRD patterns made on the as cut sample surfaces that: The intensity of the line (002) increased and that of the line (200) decreased; so that the ratio of I(002)/I(200) became greater than one, while it was less than one for the as-sintered and the ground samples (Table 3). The domain switching is ferroelastic and the toughening mechanism is a ferroelastic transformation; which is related to the change from one equilibrium state to another by domain reorientation. Since both are stress assisted, one may ask whether it is the tetragonal to monoclinic transformation or the domain switching. Actually, they can occur simultaneously. The difference between them is that the transformation toughening requires the material to contain a substantial amount of transformable tetragonal phase, while domain switching occurs in transformable phase and non-transformable phase as well. This agrees with the results of Mehta, K. et al. [38] who showed that the domain switching is not related to transformation, reversible or otherwise, but can be explained by ferroelastic domain switching. The XRD made on the fracture surface of the cut sample (Fig. 7) showed a change in the intensity of the (002) line relative to the (200) line indicating that ferroelastic transformation takes place along with transformation toughening.

It has been reported that the intensity ratio I(002)/I(200) changed upon applying electric field for lead titanate zirconate PZT (95/5) [39, 40]. The fracture toughness of PZT increased after poling, which resulted in domain reorientation [41]. The contribution of the ferroelastic domain switching to the stress intensity factor KIC is small. It was approximately estimated to be in the order of 2 MPa√m, or less for tetragonal zirconia ceramics [36]. This might be explained by the small increase of the intensity of the line (002) over the line (200) shown on Fig. 7. In spite of the difficulty of doing XRD on such a very rough and rather irregular surface, the results indicated that a ferroelastic transformation might exist.

4 Conclusions

The effect of gadolinia addition to unstabilized zirconia doped with ceria ceramics has been studied. Doping 5Ce and 10Ce zirconia with small amounts of gadolinia: 0.85 mol % and 0.5 mol %, respectively, could stabilize the tetragonal phase and produce strong and dense ceramics. The method of doping 5Ce and 10Ce previously prepared by co-precipitation with gadolinia powder led to the production of fine-grained ceramics with high stability of the tetragonal phase and consequently high strength. Strengths up to 795.8 MPa could be obtained when the cut surfaces were subjected to the tensile stress during the bending test. High fracture toughness of 17 MPa√m was obtained for the 10CeGd material sintered at 1500 °C for 2 h. The gadolinia-doped materials prepared during this work are tougher than the YTZP ceramics and stronger than the Ce-TZP ceramic. The XRD showed substantial increase in the intensities of the lines (111) and (111) of the monoclinic phase, besides an increase in the intensity of the line (002) relative to the line (200) on the as the cut and fracture surfaces. The toughening mechanisms are mainly transformation toughening and possibly simultaneously ferroelastic domain switching toughening. The domain switching toughening has been discussed in view of Chevalier's model. Extensive work and more detailed studies have to be done on the ferroelastic transformation toughening of the tetragonal zirconia ceramics. This might be an interesting subject for future work.