Gd3+ and Bi3+ co-substituted cubic zirconia; (Zr1−x−yGdxBiyO2−δ): a novel high κ relaxor dielectric and superior oxide-ion conductor

Solid oxide fuel cells (SOFCs) offer several advantages over lower temperature polymeric membrane fuels cells (PMFCs) due to their multiple fuel flexibility and requirement of low purity hydrogen. In order to decrease the operating temperature of SOFCs and to overcome the high operating cost and materials degradation challenges, the Cubic phase of ZrO2 was stabilized with simultaneous substitution of Bi and Gd and the effect of co-doping on the oxide-ion conductivity of Zr1−x−yBixGdyO2−δ was studied to develop a superior electrolyte separator for SOFCs. Up to 30% Gd and 20% Bi were simultaneously substituted in the cubic ZrO2 lattice (Zr1−x−yGdxBiyO2−δ, x + y ≤ 0.4, x ≤ 0.3 and y ≤ 0.2) by employing a solution combustion method followed by multiple calcinations at 900 °C. Phase purity and composition of the material is confirmed by powder XRD and EDX measurements. The formation of an oxygen vacant Gd/Bi co-doped cubic zirconia lattice was also confirmed by Raman spectroscopy study. With the incorporation of Bi3+ and Gd3+ ions, the cubic Zr1−x−yBixGdyO2−δ phase showed relaxor type high κ dielectric behaviour (ε′ = 9725 at 600 °C at applied frequency 20 kHz for Zr0.6Bi0.2Gd0.2O1.8) with Tm approaching 600 °C. The high polarizability of the Bi3+ ion coupled with synergistic interaction of Bi and Gd in the host ZrO2 lattice seems to create the more labile oxide ion vacancies that enable superior oxide-ion transport resulting in high oxide ion conductivity (σo > 10−2 S cm−1, T > 500 °C for Zr0.6Bi0.2Gd0.2O1.8) at relatively lower temperatures.


Introduction
Since aer the discovery of superior oxide-ion conductivity in Mg doped perovskite structure Na 0.5 Bi 0.5 TiO 3 , 1 a well known piezoelectric material that possesses high leakage conductivity that makes the material unsuitable for piezo-and ferroelectric applications, newer interest is open to developing superior oxide-ion conducting materials through controlling the nature of dielectricity of the materials. The fast oxygen ion diffusion of Na 0.5 -Bi 0.5 TiO 3 (NBT) is attributed to the high polarizability of Bi 3+ and is mediated by oxygen vacancies 1 that can be introduced either by changing the NBT compositions through Bi deciency or by Mg doping. 1,2 Dielectric leakage or relaxor-like characteristics of ferroelectrics or high k dielectric materials reveal as a strong temperature and frequency dependence in the maximum of both real and imaginary parts of the dielectric permittivity. However, relaxors not only show particular and intriguing behaviours in the dielectric response, but also show promising activity in fast-ion conduction to be applied as oxide-ion conductors for application in solid oxide fuel cells (SOFCs), oxygen separation membranes, oxygen sensors and oxygen pumps. [3][4][5][6][7][8][9] Intermediate temperature solid oxide fuel cells (IT-SOFCs) have gained recent attention due to their potential long-term durability, shorter start-up times and economic competitiveness for a wide range of applications, such as small-scale portable devices, automotive auxiliary power units and large distributed power generation systems. [10][11][12][13][14][15] However, signicant increases in power losses factors especially ohmic and activation losses, due to relatively high temperature (T > 800 C) operations reduce the cell performance. [16][17][18][19] The ohmic and activation losses are primarily related to oxide-ion transport through the electrolyte and the sluggish reaction kinetics on the electrode surfaces. These losses can be reduced by using the electrolyte materials with high ionic conductivity at low temperatures, reducing electrolyte thickness, increasing reactant concentration, and a number of potential reaction sites, and decreasing the activation barrier. In the past few decades, signicant research has been done in the development of perovskite and uorite based oxide ionic conductors, e.g., LaGaO 3 based (Sr and Mg doped) perovskites, [20][21][22][23] rare earth doped ceria based materials, 6,9,[24][25][26]  Ceria-based materials, especially rare earth-doped ceria (GDC and SDC), have been considered strong candidates for IT-SOFCs electrolytes due to their high ionic conductivity in intermediate temperature range. But their performance suffers/ degrades due to electronic conduction resulting through partial reduction of Ce 4+ into Ce 3+ at low oxygen partial pressures. 33,34 This chemical instability of ceria restricts the application of the electrolyte resulting the issue of the stability of the cell. Further, high cost of gallium and formation of inactive secondary phases during the preparation of LaGaO 3 -based electrolyte is a serious concern that hamper the applicability of the material as an oxide-ion electrolyte in SOFCs. That is why more attention is given on the fabrication of thin electrolytes supported SOFCs relying on yttria-stabilized zirconia (YSZ) that has been widely used as an electrolyte material at high temperatures rather than the ceria-based electrolyte. [35][36][37][38][39][40] In ZrO 2 -based materials, combination of high dielectric permittivity and thermal stability with low leakage current due to a reasonably high barrier height that limits electron tunnelling, counts it to further research as oxide ion conductor for SOFCs application. [36][37][38] Also, Bi-based based oxide ion conductors demonstrate the remarkable ionic conductivity due to highconcentration of intrinsic oxygen vacancies and high polarizability of Bi 3+ with 6s 2 lone pair electrons. 39 A zirconia-based electrolyte (YSZ) is considered the most effective candidate as a solid electrolyte for electrochemical cells working either in open-circuit mode (oxygen sensor) or in a power application (oxygen pump and solid-oxide fuel cell) due to its robustness. [35][36][37][38][39][40] However; it seems to be an arduous task to achieve IT-SOFCs at a commercial scale using YSZ based electrolytes due to its relatively low oxide-ion conductivity at intermediate temperatures.
Recent studies demonstrate that the relaxor nature of high k dielectricity and higher polarizability of Bi 3+ ion seems to play a directive role in providing superior oxide-ion transport throughout the lattice at temperature close to dielectric relaxation temperatures. 1,4,47 ZrO 6 octahedra was a feudal point in developing superior high k dielectric/ferroelectric materials especially in PZT based perovskite structures. Relaxor type of high dielectric materials were also reported for doped cubic zirconia phases. [41][42][43][44] We have envisaged that high polarizability of Bi 3+ ion couple with high k dielectric relaxation (high dielectric leakage) can generate superior oxide-ion conduction near T m (the temperature of the maximum dielectric permittivity). To realize the concept, we attempted the suitable dopping of Bi 3+ and Gd 3+ ions into ZrO 2 lattice to stabilize cubic phase of zirconia and found that the synergistic interaction by introducing a secondary substituent (Gd 3+ ions) enhances the oxide-ion vacancy transport within the percolation limit of ion transport inside the host structure at lower temperatures. In cubic ZrO 2 , the theoretical ratio of the ionic radius of the cation to anion (O 2À ) for fully packed FCC lattice is 0.73 at room temperature, but the ratio is 0.59 for tetragonal phase of ZrO 2 stabilized at room temperature. 45 Hence, doping of other elements with larger ionic radius than Zr at Zr site is an efficient way to stabilize the hightemperature cubic phase at room temperature by the formation of solid solutions. Our study show that the co-doping of Bi 3+ and Gd 3+ ions (ionic radii in 8 coordinations, Bi 3+ ¼ 1.17 A and Gd 3+ ¼ 1.053 A) 46 in ZrO 2 lattice resulted the formation of Zr-Bi-Gd-O solid solution in cubic uorite structure at room temperature and also it resulted the superior oxide-ion transport (oxide ion conductivity $ 10 À2 S cm À1 above 500 C) at lower temperatures. Material also showed relaxor type dielectric nature of solid solution coupled with synergistic interaction of Gd and Bi in solid solution Zr 1ÀxÀy Bi x Gd y O 2Àd . Here, we present the synthesis, characterization, permittivity and oxide-ion conductivity studies of Bi 3+ and Gd 3+ substituted cubic zirconia in this manuscript.

Material's synthesis and characterization
Zr 1ÀxÀy Bi x Gd y O 2Àd samples were synthesized by employing solution combustion method by dissolving stoichiometric amount of ZrO(NO 3 )$xH 2 O, Bi 2 O 3 and Gd 2 O 3 in 100 ml of 40% nitric acid solution with continuous stirring at 90 C for 4-5 hours. Further for auto-combustion, glycine was used as the fuel and was added in a molar ratio of 1.5 : 1 to total moles of metal ions present in the solution. The temperature of the hot plate-magnetic stirrer was increased to 250 C for combustion to start. Reaction ends up with vigorous combustion aer the evaporation of water at gelation point due to constant heating. The material le behind aer combustion was collected, and multiple calcinations were carried out at 900 C for 12 hours to get single-phase materials. For conductivity measurement, the powder was made into pallets of 10 mm diameter and $0.2 cm thickness by pressing it to $8 ton weight on a hydraulic press. These pallets were red at 1000 C for 10 hours for densication. Density of the pellet was measured by using Archimedes method and it was found to be $97% of the apparent density obtained from geometrical analysis.
The phase formation study was carried out through Rigaku Miniex desktop X-ray diffractometer (XRD) with Cu Ka radiation (l ¼ 1.54 A) in the range 2q $ 10-90 with a step size of 0.02 . The structures were rened by the Rietveld renement method using the FULLPROF suite soware package and cubic uorite ZrO 2 (space group: Fm 3m) as model structure. The microstructures of the sintered samples were investigated by using scanning electron microscopy (EVOscanning electron microscope MA15/18). The average grain size was calculated using the linear intercept method. The composition of the compounds was examined by energy dispersive X-ray (EDX) spectroscopy with a probe attached to the SEM instrument. Raman spectroscopy of powdered sample was carried out by using STR-300 micro-Raman spectrometer with a laser excitation wavelength of 532 nm and step size of 1.9 cm À1 .
Pt paste was used as a blocking electrode for conductivity measurements. For this purpose, the sintered pellets were coated with platinum paste and cured at 800 C for 30 minutes. The conductivity measurements were performed using Autolab potentiostat as a function of frequency from 1 MHz to 1 Hz at different temperatures varying from 100 C to 650 C. All measurements were taken during the cooling cycle from 650 C to 100 C.

Result and discussion
As Gd 3+ ions can stabilized the ZrO 2 in cubic uorite structure, 37 role or promoting effect of Bi 3+ ion were utilized to develop superior oxide-ion conductors. Several compositions of Zr 1ÀxÀy Gd x Bi y O 2Àd in cubic uorite structure were synthesized and few important data were presented in the manuscript. We have found that at max, total 40% ions can be substituted at Zr site to make single phase cubic uorite material using Gd 3+ and Bi 3+ as simultaneous substituent. Thus, up to 20% of Bi and 30% Gd was co-substituted in ZrO 2 lattice (Zr 1ÀxÀy Gd x Bi y O 2Àd , x + y # 0.4, x # 0.3 and y # 0.2) in different combinations and several solid solutions were synthesized in cubic uorite structure. The synthesized Bi 3+ and Gd 3+ substituted ZrO 2 powder was in off-white in colour. The crystal structure and phase purity of the material was analyzed by powder XRD study. Powder XRD pattern of Zr 0. 6  Rietveld renement of powder XRD pattern is given in Table 1. Due to substitution of larger Bi 3+ and Gd 3+ cations on Zr site in ZrO 2 lattice, there was increase in the lattice parameter of the materials with increase in concentration of dopants.
SEM micrographs of Zr 0.6 Bi 0.2 Gd 0.2 O 1.8 (powder, top view and cross section of the pellet used for conductivity measurements) are shown in Fig. 3(a-c). The SEM study shows that the powders are made of with interconnected grains in size of 4-10 mm. Fig. 3(b and c) show the top and cross section images of the pellet. Crystal growth during sintering resulted microstructure is having nearly no or very low porosity and the grains are     Fig. 3(d)) conrms that the elements Zr : Gd : Bi were present in the ratios of 0.589 : 0.195 : 216 that is very much close to the elemental ratios used for the synthesis.
Thus XRD study and SEM study coupled with EDX study conrms that Bi 3+ and Gd 3+ ions are substituted at Zr 4+ sites in stabilized cubic ZrO 2 lattice. Considering the Schottky defect formation due to substitution of Bi 3+ and Gd 3+ at Zr 4+ sites that will create oxygen vacancy generation in the lattice and the oxygen defect formation equation using Kröger-Vink notation can be represented as: Further the lone pair of Bi 3+ ions can stabilize the oxide-ion vacant uorite structure as represented in Fig. 4. The lone pair of Bi 3+ ion are known to implant higher vacancy mobility as it was witnesses in the case of Bi based oxide-ion conductors. 1,4,[27][28][29][30][31]47    The Raman spectrum for cubic ZrO 2 is characterized by a narrow band at 145 cm À1 and broad bands centered around 250, 305, 440, and $601 cm À1 . The stabilized ZrO 2 sample in this study clearly showed the broad peak between 500 to 650 cm À1 that is related to the disordered oxygen sub-lattice along with mass-related disorder indicating of a large disorder in the cationic cage upon Gd and Bi ion substitution in cubic ZrO 2 lattice whereas monoclinic ZrO 2 exhibits several well dened sharp bands because of the symmetry reduction. 48 Since the cations are much heavier than the oxygen atoms, they are the major contributors to the vibrations associated with the acoustic branches indicating a periodic arrangement of the vacancies in stabilized Gd/Bi co-doped cubic-zirconia the lattice. 49-51 A careful examination of the Raman spectra also shows weak bands around 620, 660, and 815 cm À1 that could be associated with the rearrangement of the anionic sub-lattice, i.e. oxygen ions and vacancies containing Bi cage in stabilized Bi/ Gd codoped cubic zirconia. 52 Further observed bands around 535 nm and 790 nm can be assigned to Raman vibrations of Gd containing sub lattice of stabilized Bi/Gd codoped cubic zirconia. 53 Thus the Raman spectroscopy study clearly reveal the formation of oxygen vacant Gd/Bi co-doped cubic zirconia lattice.
To see the effect of Gd and Bi substitution in the cubic uorite ZrO 2 structure, impedance spectroscopy was utilized to study the oxide-ion vacancy conduction process and dielectric constant of the materials at various temperatures in different environments. Fig. 6(a)  increasing temperature for all the compositions. The best electrical conductivity of this series was observed for the composition Zr 0.6 Bi 0.2 Gd 0.2 O 1.8 . At 550 C, the measured conductivity was $10 À2 S cm À1 for Zr 0.6 Bi 0.2 Gd 0.2 O 1.8 , which is better than that of Zr 0.92 Y 0.08 O 2 (YSZ) at 670 C at and of La 0.8 -Sr 0.2 Ga 0.83 Mg 0.17 O 3 (LSGM) at 600 C. Here, a careful study was made on development to superior ZrO 2 based electrolyte. As Gd can stabilized the ZrO 2 in cubic uorite structure, role or promoting effect of Bi 3+ ion were utilized to develop superior oxide-ion conductors. Systematically, we carried out ionic conductivity study of Gd stabilized ZrO 2 for 15%, 20%, 25%, 30% Gd 3+ ions doping in ZrO 2 and ionic conductivity data of these materials are provided in Table 2. It was found that 20 and 25% Gd stabilized cubic ZrO 2 showed almost similar conductivities. Further to improve the conductivity of Gd stabilized cubic ZrO 2 , Gd 3+ and Bi 3+ ion co-substituted cubic zirconia was synthesized. We have found that at max, total 40% ions can be substituted at Zr site to make single phase cubic uorite material. In rst attempt additional 15% Bi 3+ co-doping was attempted along with Gd 3+ ions. However, in case of 30% Gd stabilized zirconia, only 10% additional Bi 3+ can be doped in single phase. Among them, we Zr 0.65 Gd 0.20 Bi 0.15 O 2Àd showed highest conductivity. Further we extended Bi 3+ ion substitution in Gd stabilized cubic zirconia and found that highest conductivity can be achieved with Zr 0.6 Gd 0.2 Bi 0.2 O 2Àd sample. Thus this study can conrm that maximum 40% substitution in ZrO 2 lattice can be achieved using Gd 3+ and Bi 3+ ions together, and the highest conductivity was achieved for cubic uorite Zr 0.6 Gd 0.2 Bi 0.2 O 2Àd sample. The data for oxide-ion conductivity of different samples of Zr 1ÀxÀy Bi x Gd y O 2À(x+y)/2 at different temperature along with the data of other competitive oxide-ion electrolyte in the same temperature range is given in Table 2. As shown in Table 2, highest conductivity (1.1 Â 10 À2 S cm À1 at 550 C) was observed for Zr 0.6 Gd 0.2 Bi 0.2 O 2Àd . As evident from the study, the conductivity of the materials was increased with increasing Gd content in the cubic uorite ZrO 2 lattice. At initiation of co-doping, the promoting effect of Bi is clearly visible on co-doping of Bi along with Gd in ZrO 2 lattice and we found that the maximum conductivity was observed for Zr 0.6 -Gd 0.2 Bi 0.2 O 2Àd sample as maximum co-doping or simultaneous substitution of Gd and Bi in cubic uorite ZrO 2 lattice is limited to 40%.
Impedance study of Zr 0.6 Bi 0.2 Gd 0.2 O 1.8 was also carried out at different temperatures in dry hydrogen (UHP H 2 ) and dry nitrogen (UHP N 2 ) environment ( Fig. 6(b)) also to see the effect of absorbed moisture, and impurities present in the air on the surface or at oxide-ion vacancy sites of the sample and also the stability of the material in reducing environment in presence of hydrogen. Below 500 C, the total conductivity of Zr 0.6 Bi 0.2 -Gd 0.2 O 1.8 was found little lower in hydrogen and nitrogen atmosphere compared to air. The cubic uorite phase of ZrO 2 , YSZ is predominantly a total oxide-ion conductor. Below 500 C, Gd and Bi doped ZrO 2 sample in air atmosphere may have little bit associate protonic conduction contribution due to presence of existing moist into the air. As moisture present in air can result absorption of moisture on the surface of the sample at low temperature contributing to additional conductivity at those temperatures. As we have not found any increase in total conductivity in hydrogen atmosphere even at higher temperatures, this suggests the stability of Bi 3+ ion in ZrO 2 lattice that does not allow the reduction Bi 3+ ions in hydrogen media. The Cole-Cole plot at 500 C in air atmosphere for Zr 0.6 Bi 0.2 Gd 0.2 O 1.8 is shown for understanding the polarization and oxide-ion transport nature of the sample (Fig. 6(c)). The linear tail present in the plot clearly suggests ionic conduction pathways. Thus the total conductivity in Gd and Bi co-doped ZrO 2 sample is ionic in nature and facilitates oxide-ion conductivity due to oxygen vacancy migration. Further, we have also characterized the Zr 0.6 Bi 0.2 Gd 0.2 O 1.8 sample heated in 10% hydrogen balanced in nitrogen atmosphere at 800 C for 6 h by powder XRD study and we have not found any diffraction peaks for Bi metal in the XRD of the sample as all the peaks were identied to cubic phase of zirconia only (Fig. 6(e)). Further no colour changes were observed for the sample heated in H 2 atmosphere at 800 C for 6 h. These study clearly suggest the stability of the material in reducing media and also suggest that the total conductivity of our samples are predominantly an oxide-ion conduction as Gd 3+ and Bi 3+ doped sample contain oxide-ion vacancy. A comparison of oxide-ion conductivity of Zr 0.6 Bi 0.2 Gd 0.2 O 1.8 in the same temperature range of 300-650 C with other established oxideion conductors having crystal structures of uorite or perovskite also presented in (Fig. 6(d)). Oxide-ion conductivity of Zr 0.6 Bi 0.2 Gd 0.2 O 1.8 (ZBGO) is very much comparable to Sr and Ga doped LaGaO 3 ; LSGM and KTa 0.4 Ti 0.3 Ge 0.3 O 2.7 (KTTGO). The activation energy for oxide-ion conductivity was found as low as 0.42 eV. In the case of all samples, a sudden increase in oxideion conductivity was found at or around 450 C.
Further to understand the sudden increase in conductivity if it is associated with any phase transformation, thermogravimetric and differential scanning calorimetry (TGA-DSC) analysis at a constant heating rate of 10 C per minute in the temperature range of 30-900 C in N 2 atmosphere. Fig. 7 shows the TGA plot for Zr 0.6 Bi 0.2 Gd 0.2 O 1.8 sample preheated at 120 C. The lack of physically adsorbed water on the sample was demonstrated as marginal weight loss was observed up to 150 C followed by very little weight loss ($0.5%) up to 900 C. The TGA analysis conrms the relatively low hygroscopicity or dry nature of the material. The DSC curve shown in Fig. 7 does not show any signicant feature for any associated phase change that may arrive from oxide-ion vacancy or structure reorientation. Thus the TGA/DSC studies conrm the structure stability of the material in the temperature range of 30-900 C. In addition, an FT-IR study was also performed to monitor the presence of hydroxide ions or water absorption at the oxygen vacancy position or at the surface of the Bi and Gd cosubstituted ZrO 2 samples. Fig. 8 displays the FT-IR spectra of Zr 0.6 Bi 0.2 Gd 0.2 O 1.8 sample preheated at 120 C for about 1 hour. Absence of peaks between 3300 and 4000 cm À1 clearly suggests the absence or insignicant presence of hydroxide ions or physio-adsorbed water on the surface of the material. This conrms that the conductivity observed for cubic uorite Zr 1ÀxÀy Bi x Gd y O 2À(x+y)/2 samples are only due migration of oxideion vacancies in the lattice.
The sudden change or activation of oxygen vacancy migration above 450 C conrms the oxide-ion transport within the percolation limit of a conductive phase transition coupled with thermal activation. Further to understand the effect of dielectric polarizability on oxide-ion conductivity, the dielectric constant in the frequency range of 20 kHz to 100 kHz at different temperatures is plotted in Fig. 9(a). The dielectric studies show a relaxor type behavior coupled with diffusive phase transition $9725 at 600 C at applied frequency 20 kHz, much higher than those for the pure ZrO 2 . With increasing applied frequency, the T m had varied and observed to decrease. This gradual decrease with increasing applied frequency conrms the relaxor behaviour of this high k dielectric material. Also, these compositions show a rather high dielectric loss (tan d > 100) above $400 C that increases exponentially with temperature above 600 C ( Fig. 9(b)) suggesting high leakage current at elevated temperatures. The dielectric relaxation of the dipole moment can lead to the material's superior oxide-ion conductivity at the temperature close to T m . The relaxation of net dipole moment generated over oxygen vacant octahedra can play a vital role in reorientation of the polyhedra at elevated temperature to provide the short transport pathways for the oxide-ion vacancy migration. Thus this giant loss (high leakage) seems to be associated with conduction or migration of oxide vacancy. It is well known that higher concentration of oxide-ion vacancies at lattice sites and their high mobility are two key factors for achieving high ionic conductivity in typical oxide ion conductors. For cubic zirconium oxides, oxide ionic conduction is primarily associated with the conducting passageway through a cubic block and the movement of oxygen vacancies dominates. Similar to PZT-based ferroelectric ceramics, the vibration of a smaller-sized Ti 4+ and Zr 4+ cation from its mean position in octahedral coordination was shown to have the high dielectric constant and relaxor type behaviour and associated high oxideion conductivity in KTa 0.4 Ti 0.3 Ge 0.3 O 2.7 . 5 Here, Bi 3+ and Gd 3+ ions were doped into ZrO 2 lattice to stabilize the materials in cubic phase and the lone pair of Bi 3+ ions can also play a vital role in enhancing the polarizability of the solid solution. Further, the synergistic interaction by introducing a secondary substituent (Gd 3+ ions) seem to play important role to enhance the oxide-ion vacancy transport within the percolation limit ion transport inside the host structure at lower temperatures. High polarizability of Bi 3+ with 6s 2 lone pair electrons has been viewed as a key factor for high ionic conductivity in Bi-based oxide conductors, e.g., d-Bi 2 O 3 , 30 Bi 4 Ti 3 O 12 , 5 g-Bi 4 V 2 O 11 . 47 Similarly, in the present material, the 6s 2 lone pairs of Bi 3+ ions can be oppositely pointed toward a vacant central of the Zr plane of ZrO 8 polyhedra in the parent-phase as shown in Fig. 4. Also, this structure can have a relaxed unit cell with longer Zr-O bonds, where the oxygen vacancy can jump by thermal activation to the energetically equivalent neighbouring oxygen sites of the lattice. However, more advanced structural and phase transition studies such as neutron powder diffraction (NPD) or EXAFS studies at various temperatures are important for the making mechanistic propositions about the associated phase transition responsible for sudden increase in the conductivity of the materials at elevated temperatures. However, the direct correlation of dielectric relaxation of dipole moments to superior oxide-ion transport was also observed previously for Na 0.5 Bi 0.5 -TiO 3 , 1-3,27,28 KTa 1ÀxÀy Ti x Ge y O 3Àd , 5 20% Sm doped CeO 2 (Ce 0.8 -Sm 0.2 O 2Àd ) 6 and La 2 Mo 2 O 9 . 7 However further studies are necessary to validate the relationship of dielectric relaxation and associated phase transitions that provide shorter conduction pathways for material's superior oxide-ion conductivity.