Phase Composition and Stability, Sintering and Thermal Conductivity of Gd₂O₃ and Yb₂O₃ Co-Doped YSZ

Abstract

Y₂O₃-stabilized ZrO₂ (YSZ) has been the material of choice for thermal barrier coatings (TBCs) in the past decades, yet its phase decomposition limits its application above 1200 °C. In this study, Gd₂O₃ and Yb₂O₃ co-doped YSZ powders were produced, in which some amounts of monoclinic (m) phase were introduced into the cubic (c) phase matrix. XRD results showed that the fabricated powders obtained by a solid phase synthesis were composed of m and c phases, and hat the m phase content decreased in a sequence of 4Gd-2Yb-4Y, 2Gd-2Yb-6Y and 2Gd-4Yb-4Y powders. This indicated that Yb³⁺ is an excellent stabilizer in the ZrO₂-based lattice, which could largely suppress the formation of the m phase. The m phase content in the powders was almost kept unchanged with heating at 1300 °C, which could provide a toughening effect to the ceramic. All the powders exhibited no obvious sintering at 1300 °C for 150 h. As compared to YSZ, the three fabricated ceramics had lower thermal conductivities, and they increased in a sequence of 4Gd-2Yb-4Y, 2Gd-4Yb-4Y and 2Gd-2Yb-6Y.

1. Introduction

ZrO₂ stabilized by oxides has been extensively studied, and one of the representative stabilizers is Y₂O₃ [1,2,3]. Due to its excellent high-temperature mechanical and thermal–physical properties, Y₂O₃ partially stabilized ZrO₂ (YSZ) has been practically applied in electronics [4,5], atomic energy, the aviation industry, and especially to the thermal barrier coatings (TBC) which offer thermal insulation to the hot components of aero-engines and thus effectively improve their combustion efficiency [6,7,8]. YSZ TBCs have been successfully used for several decades [9,10], but they have become less and less reliable with the development of aero-engines [11,12,13,14]. One crucial aspect in this is issue of the phase decomposition and phase transformation [1,15,16]. As is known to all, YSZ TBCs are usually prepared either by electron beam physical vapor deposition (EB-PVD) [17,18] or by air plasma spraying (APS) [19,20,21]. Additionally, they exhibit a non-transformable metastable tetragonal (t’) phase, which is desirable for thermal insulation yet. At temperatures above 1200 °C [22], the t’-YSZ decomposes into cubic (c)-YSZ and t YSZ, the latter transforming into a monoclinic (m) phase accompanied by some volume expansion [23,24], which is fast and uncontrollable [25]. When the phase-transformed part reaches a critical value, the resulting volume expansions lead to crack initiation, propagation, and even coating spallation.

Many efforts have been made to enhance the high-temperature phase instability of the t’-YSZ, such as selecting alternate stabilizers to Y₂O₃ for ZrO₂-based systems. Cairney et al. reported that the phase stability of ZrO₂, stabilized by Yb₂O₃, is better than that of YSZ at 1300 °C [26]. Er₂O₃, Sm₂O₃, Nd₂O₃, Sc₂O₃ and Gd₂O₃ have also been studied as stabilizers for ZrO₂, and Er₂O₃-ZrO₂ has the best phase stability [27,28,29,30,31,32]. It is worthy noted that the Gd₂O₃ dopant improves the sintering resistance of both bulk Gd₂O₃-ZrO₂ (4 mol% stabilizer) and YSZ TBCs, although it decreases the t’ phase stability, as reported by Rahaman [27]. Er₂O₃, Sm₂O₃, and Nd₂O₃ have also been studied as stabilizers for ZrO₂, and the t’ phase stability of Er₂O₃-ZrO₂ is the best [28]. Additionally, some reports have revealed that co-doping two or more rare earth oxides (RE₂O₃) into ZrO₂ has an effective function in improving the t’ phase stability. Guo et al. found better t’ phase stability in Sc₂O₃-Gd₂O₃-ZrO₂ ceramics with increasing Sc₂O₃ content [33]. Moreover, Guo et al. has analyzed the t’ phase stability of RE₂O₃ (RE = La, Nd, Gd, Yb)-Yb₂O₃-YSZ at 1400 °C [34].

Although the modified t’ phase stability has improved at high temperatures, it still inevitably undergoes phase decomposition with long-term annealing. Despite its excellent thermal conductivity and high-temperature phase stability, C-ZrO₂ is unfortunately not tough enough for use in TBCs because of its low toughness [15]. The volume expansion, resulting from the t to m phases transformation, can absorb or consume the energy at the crack tip and generate compressive stresses in the region of the main crack action, something which improves the strength and fracture toughness of the coating [20]. Therefore, it is reasonable to believe that setting an appropriate amount of m phase in the c-ZrO₂ and maintaining the m phase content during service can effectively improve the performance of ZrO₂ ceramics.

In this study, we prepared YSZ co-doped with Gd₂O₃ and Yb₂O₃ via a solid-state reaction, which consisted of c and m phases. An investigation was conducted into the effect of co-doped RE₂O₃ on the monoclinic. The phase stability of the fabricated powders which were heat-treated at 1300 °C was evaluated. In addition, the sintering behavior and thermal conductivity of the co-doped ZrO₂ ceramics were investigated and compared with those of YSZ. Finally, the mechanisms of the reduced thermal conductivity, resulting from the rare earth co-doping, were discussed.

2. Experimental Procedure

The co-doped ZrO₂ were prepared by solid-state reaction at 1400 °C for 12 h, using Gd₂O₃, Yb₂O₃, Y₂O₃ (99.99% purity) and ZrO₂ (99.9% purity) as raw materials. The total stabilizer content of each doped zirconia is 10 mol%, and the specific components are expressed in mol% as follows: 2Gd₂O₃-2Yb₂O₃-6Y₂O₃-90ZrO₂ (2Gd-2Yb-6Y), 2Gd₂O₃-Yb₂O₃-Y₂O₃-90ZrO₂ (2Gd-4Yb-4Y), and 4Gd₂O₃-2Yb₂O₃-4Y₂O₃-90ZrO₂ (4Gd-2Yb-4Y). The raw powers were mixed in a planetary ball mill with absolute ethanol for more than 12 h to obtain a homogeneous mixture. The homogeneously mixed solution was dried, placed in a zirconia crucible at 1400 °C and sintered for 12 h to produce RE₂O₃-stabilized ZrO₂ ceramics. Meanwhile, some mixed powder was cold-pressed at 250 MPa and then sintered at 1400 °C for 10 h to obtain the bulk material for thermal conductivity measurements. A scanning electron microscope (SEM; JEOL 7900F) was used to characterize the morphologies of ceramic powders after calcination and annealing at 1300 °C for 50h, 100, and 150 h. Phase stability was characterized by X-ray diffraction (X′Pert³ MRD, Malvern Panalytical, Cu Kα radiation). The scanning frequency was 0.05°/s. The molar content of the m phase was calculated using Miller’s formula [35], given in Equations (1) and (2):

$$\frac{M_m}{M_{c,t}}=0.82\frac{I_m\left(\overline{1}11\right){+I}_m\left(111\right)}{I_{c,t}\left(111\right)}$$

(1)

$$M_m{+M}_{c,t}=1$$

(2)

where M_m and M_c,t are the mole fractions of the m and c/t phases, respectively, and I is the integrated intensity corresponding to the peaks concerned.

The thermal diffusivity (α) of RE₂O₃ co-doped ZrO₂ was measured between 20 °C and 1200 °C using a laser flash meter (LFA 427; Netzsch, Selb, Germany), with an interval temperature of 200 °C. Before measuring the thermal diffusivity, the front and back surface of the samples were coated with a graphite film to allow thermal absorption of the laser pulse. Each sample was measured three times at selected temperatures. The specific heat capacity (C_p) is calculated from the heat capacity values of the individual constituent oxide based on the Neumann-Kopp rule [36]. The density (ρ) was measured by the Archimedes’ method. The thermal conductivity (λ) was calculated by Equation (3):

$${\lambda =\rho ·\alpha ·C}_p$$

(3)

Since the sintered sample cannot be completely dense, the influence ($\varphi$) on the thermal conductivity of the material can be fixed using the following equation, in which λ₀ expresses the full density of the sample (λ₀) [37]:

$$\frac{\lambda }{{\lambda }_0}=1-\frac{4}{3}\varphi$$

(4)

3. Results and Discussion

Figure 1 shows the XRD patterns of the as-fabricated 2Gd-2Yb-6Y, 4Gd-2Yb-4Y and 2Gd-4Yb-4Y powders. For all samples, the primary phase is the c phase. In the range between 2θ ≈ 28.3° and 31.5° of all the patterns the diffraction peaks were found, which are characteristic peaks of the m phase. Hence, it could be identified that the fabricated powders are composed of c and m phases. The m phase contents in the three types of co-doped ZrO₂ were calculated and illustrated in Figure 2. It could be found that the m phase content in the samples increases in a sequence of 2Gd-4Yb-4Y, 2Gd-2Yb-6Y and 4Gd-2Yb-4Y powders.

The different m contents in the three samples could be attributed to the difference in the stabilization ability of Yb, Y and Gd. When rare earth cations are used to stabilize ZrO₂, a smaller size misfit between RE³⁺ and Zr⁴⁺ ions in the ZrO₂ lattice is desirable. The cation radii of Yb³⁺, Y³⁺, Gd³⁺ and Zr⁴⁺ are 0.985 Å, 1.019 Å, 1.053 Å and 0.84 Å in the 8-coordination, respectively. Thus, doping Yb³⁺ into the ZrO₂ lattice is considered to have the most ability to obtain a stable c phase [38]. Based on this consideration, the variation tendency of the m phase content in the samples could be well explained. Comparing samples 2Gd-4Yb-4Y and 2Gd-2Yb-6Y, they have the same range of Gd, but the first sample is rich in Yb, while the second is rich in Y. Since Yb³⁺ is smaller than Y³⁺, it is reasonable that the 2Gd-4Yb-4Y sample has a higher content of c phase content than the 2Gd-2Yb-6Y sample. Thus, the content of m phase in the 2Gd-4Yb-4Y sample is lower than that in the 2Gd-2Yb-6Y sample. The Yb contents of the 2Gd-2Yb-6Y and 4Gd-2Yb-4Y samples are identical, but the former contains lower Y³⁺ levels. Hence, the 2Gd-2Yb-6Y sample has a higher ability to produce the c phase, and its m phase content is lower. Therefore, the m phase content follows the sequence m_2Gd-4Yb-4Y ˂ m_2Gd-2Yb-6Y ˂ m_4Gd-2Yb-4Y, further evidence that Yb³⁺ is an excellent stabilizer in the ZrO₂ lattice that can largely suppress the formation of the m phase.

The lattice parameters of the c phase in the powders 2Gd-2Yb-6Y, 4Gd-2Yb-4Y and 2Gd-4Yb-4Y were calculated, and the results are listed in

Table 1. Lattice parameters in different ceramic samples.

Gd2O3 mol%	Yb2O3 mol%	Y2O3 mol%	a = b = c (Å)
		10	5.138
2	2	6	5.133
4	2	4	5.151
2	4	4	5.131

. It was found that the lattice parameters of the samples increase in a sequence of 2Gd-4Yb-4Y, 2Gd-2Yb-6Y and 4Gd-2Yb-4Y. This could also be explained based upon the difference in the RE³⁺ radius. By comparing 2Gd-2Yb-6Y and 4Gd-2Yb-4Y, their Yb contents are identical, but the latter contains more Gd³⁺. Thus, 4Gd-2Yb-4Y has a more significant lattice parameter than 2Gd-2Yb-6Y. For the case of 2Gd-2Yb-6Y and 4Gd-2Yb-4Y, the Yb contents are the same, but 2Gd-2Yb-6Y contains more Y and thus shows a smaller lattice parameter.

To investigate the changes in the phase structure of the co-doped zirconia powder at 1300 °C, heat treatment tests were performed at different times. Figure 3 shows the XRD patterns of 2Gd-2Yb-6Y, 4Gd-2Yb-4Y and 2Gd-4Yb-4Y powders at different heat treatment times. The c and m phases are still present at each time point, and the changes in m phase content in the samples are shown in Figure 2. It can be seen that the m phase contents in all samples remain almost unchanged with the heating time. 2Gd-2Yb-6Y, 4Gd-2Yb-4Y and 2Gd-4Yb-4Y are composed of c and m phases. It is known that the c phase does not undergo any transformation during heat treatments. Therefore, it is reasonable to keep the content of m phase constant after heat treatment. However, the m phase undergoes transformation during heat treatment. It transforms into a t phase, and the t→m phase transition occurs during the subsequent cooling process. This process belongs to the martensitic transformation, a phenomenon which can provide a toughening effect to the ceramic. As a result, 2Gd-2Yb-6Y, 4Gd-2Yb-4Y and 2Gd-4Yb-4Y are expected to have higher toughness than those with a pure c phase. Even though, it should be noted that the proportion of m phase in the sample should not be too high as this would be detrimental to the ceramics for TBC applications [39,40].

Figure 4 shows the microstructure of the powders after 150 h of heat treatment. All samples showed good sintering resistance, with most powders maintaining an independent grain shape. In the 2Gd-2Yb-6Y powder, many separated grains were connected by a neck structure, the neck between the grains grew significantly, the deformation of the grains was more significant, and a clear trend of grain boundary movement was found. Conversely, in the 4Gd-2Yb-4Y and 2Gd-4Yb-4Y powders, most of the grains still maintained an independent polygonal structure with fewer and smaller necks, indicating a lower degree of sintering in these two ceramic samples. The microscopic mechanism of the initial sintering of ceramic powders is the lattice diffusion from the inner grain boundary to the neck. The low sintering rate of polyoxides-doped ZrO₂ is reasonable, Gd³⁺ has a large ionic radius, and its distribution requires higher energy. In contrast, Yb³⁺ has a smaller ionic radius, but its larger atomic mass also limits the grain boundary diffusion, a factor which improves the sintering resistance of ceramics. This is consistent with the cation distribution in Figure 5 over 150 h. Therefore, the degree of sintering is also lower. Since the ceramic samples in this study are a mixture of c and m phases, which increases the energy required for grain boundary movement, and thus doping with Gd₂O₃, Yb₂O₃, and, Y₂O₃ helps to improve the resistance of the ZrO₂-based ceramics.

Figure 6 shows the thermal conductivities of 2Gd-2Yb-6Y, 4Gd-2Yb-4Y, 2Gd-4Yb-4Y and YSZ ceramics. The error bars are neglected because they are smaller than the symbols. Firstly, the thermal conductivity of ceramics decreases with increasing temperature in the room temperature range up to 800–1000 °C and increases when the temperature exceeds 1000 °C. Secondly, it is found that all the co-doped ceramics have lower thermal conductivities than YSZ. At 1000 °C, the thermal conductivities of 2Gd-2Yb-6Y, 2Gd-4Yb-4Y, and 4Gd-2Yb-4Y are 1.889 W/mK, 1.663 W/mK and 1.420 W/mK, respectively.

It is well known that thermal conductivity is highly dependent on the non-harmonic effects caused by phonon scattering and on the vacancies caused by defects [41,42]. In this study, co-doping with RE₂O₃ causes more defects and elastic strain fields, reducing phonons average free range and leading to the lower thermal conductivity of ZrO₂ co-doped with Gd₂O₃, Yb₂O₃, and Y₂O₃ than YSZ.

For the ZrO₂ doped with RE₂O₃, the phonon scattering rates can be expressed with the following Equation (5) [43] due to the different atomic masses of doped RE³⁺ and Zr⁴⁺:

$$\frac{1}{{\tau }_{\mathsf{\Delta }M}\left(\omega \right)}=\frac{{ca}^3{\omega }^4}{4\pi v_s{}^3}{\left(\frac{\mathsf{\Delta }M}{M}\right)}^2$$

(5)

where τ is the relaxation time, α³ is the atomic volume, ν_s is the transverse wave velocity, ω is the phonon frequency, c is the ratio of the number of defects per unit volume focus to the number of dot positions, M is the Zr⁴⁺ mass of cation, and ∆M is the mass difference between the doped RE³⁺ and Zr⁴⁺. The scattering rate of the phonons is proportional to the square of ∆M. Thus, the larger the mass difference between the doped cations and Zr⁴⁺ is, the stronger the phonon scattering will be, resulting in lower thermal conductivity. Therefore, ceramics prepared in this study have lower thermal conductivities.

The doping of RE₂O₃ to YSZ produces an elastic field due to the difference in ionic radii between RE³⁺ and Zr⁴⁺, dislodging surrounding atoms and altering the phonon frequency.

Due to the difference in ionic radius between RE³⁺ and Zr⁴⁺, doping YSZ with RE₂O₃ produces an elastic field, which dislodges atoms and affects phonon frequencies, enhancing phonon scattering [43]:

$$\frac{1}{{\tau }_{\mathsf{\Delta }R}\left(\omega \right)}=\frac{{2ca}^3{\omega }^4}{{\pi v}^3}{J}^2{\gamma }^2{\left(\frac{\mathsf{\Delta }R}{R}\right)}^2$$

(6)

where J is a constant, γ is the Gruneisen’s constant, R is the Zr⁴⁺ cation radius, and ∆R is the radius difference between the doped RE³⁺ and Zr⁴⁺. Thus, the larger the radius difference between the doped cations and Zr⁴⁺, the stronger the phonon scattering, resulting in lower thermal conductivity. Hence, the fabricated ceramics in this study are shown to have lower thermal conductivities compared to YSZ.

As shown in Figure 6, the thermal conductivities of the co-doped samples increase in an order of 4Gd-2Yb-4Y, 2Gd-4Yb-4Y and 2Gd-2Yb-6Y, which could be explained by the basis of the above discussion on the effects of ∆M and ∆R on the phonon scattering. In comparison, the Y2O3 content of 4Gd-2Yb-4Y and 2Gd-4Yb-4Y ceramics is the same. Thus, only the impact of Gd and Yb on the phonon scattering needs to be considered. Compared to Yb³⁺, Gd³⁺ is larger and lighter. Thus ∆R (Gd³⁺-Zr⁴⁺) is larger than ∆R (Yb³⁺-Zr⁴⁺), but ∆M (Gd³⁺-Zr⁴⁺) is smaller than ∆M (Yb³⁺-Zr⁴⁺). Since 4Gd-2Yb-4Y has lower thermal conductivities than 2Gd-4Yb-4Y, it could be implied that the ∆R plays a more dominant role in affecting phonon scattering than the ∆M for those two ceramics. For the 2Gd-4Yb-4Y and 2Gd-2Yb-6Y ceramics, ∆M (Yb³⁺-Zr⁴⁺) is larger than ∆M (Y³⁺-Zr⁴⁺), but ∆R (Yb³⁺-Zr⁴⁺) is smaller than ∆R (Y³⁺-Zr⁴⁺). 2Gd-4Yb-4Y has lower thermal conductivities than 2Gd-2Yb-6Y, which implies that in this case, the ∆M plays a more dominant role in affecting the phonon scattering than the ∆R.

4. Conclusions

10 mol% for Gd₂O₃, Yb₂O₃, and Y₂O₃ co-doped ZrO₂ was synthesized by the solid-state reaction method:

• Analysis of the phase content analysis of the synthesized powders revealed that the content of m phase gradually decreased for the powders 4Gd-2Yb-4Y, 2Gd-2Yb-6Y, and 2Gd-4Yb-4Y. This suggests that Yb₂O₃ effectively suppresses the formation of the m phase and is an excellent stabilizer for ZrO₂. A subsequent heat treatment test at 1300 °C, conducted for 150 h, showed that m phase remained almost unchanged, indicating that it can significantly improve the toughness of a given material.
• No sintering was observed in any of the powders, suggesting that the addition of Gd₂O₃ and Yb₂O₃ improves the sintering resistance of YSZ materials.
• Co-doping Gd₂O₃ and Yb₂O₃ can effectively reduce the thermal conductivity of YSZ. The co-doping increases the discrepancy between the radius and atomic mass of RE³⁺ and Zr⁴⁺ ions, which increases the phonon scattering and reduces the thermal conductivity of the material, resulting in increased phonon scattering. The experimental results indicate the Gd₂O₃ is more effective at decreasing thermal conductivity than co-doped ZrO₂.

As a result of these comprehensive properties, 4Gd-2Yb-4Y ceramics may be considered a promising candidate for TBC manufacturing. Additionally, more work will focus on the mechanical properties of 4Gd-2Yb-4Y materials and the performance of coatings with this composition.