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Vorotnikov V.A., et al. Nanosystems: Phys. Chem. Math., 2024,15 (1), 65-79.
http://nanojournal.ifmo.ru
Original article
DOI 10.17586/2220-8054-2024-15-1-65-79
Equilibrium of intrinsic and impurity point defects in Ca-doped Sm2Zr2O7
Vladimir A. Vorotnikov1'2, Semyon A. Belyakov3, Alexey V. Ivanov1'2, Yulia V. Novikova3, Anna Yu. Stroeva1, Vadim V. Grebenev4, Dmitry N. Khmelenin4, Olga V. Emelyanova4, Maksim S. Plekhanov5, Anton V. Kuzmin1'2
1 Institute of Chemistry and Ecology, Vyatka State University, Kirov, Russia
2 Institute of Solid-State Chemistry and Mechanochemistry SB RAS, Novosibirsk, Russia 3Institute of High-Temperature Electrochemistry UB RAS, Yekaterinburg, Russia
4Shubnikov Institute of Crystallography of Federal Scientific Research Centre "Crystallography and Photonics" of the RAS, Moscow, Russia
5Institute of Crystallography, RWTH Aachen University, Aachen, Germany Corresponding author: Anton V. Kuzmin, [email protected]
PACS 82.45.Gj, 82.33.Pt
Abstract In this work, a doping strategy was used to achieve a good conductivity in samarium zirconate which crystallizes in the pyrochlore. The production of nanopowders made it possible to form high-density ceramics with an optimal microstructure. It is shown that intrinsic and impurity defects coexist in Sm2-xCaxZr2O7-5, impairing ion transport at high doping levels. Despite this, Sm195Ca0.05Zr2O7-,s maintains low activation energy of the parent and has good ionic conductivity (10-3 Scm-1 at 600 °C) which is one of the largest among oxide pyrochlores. It has been shown to have a good chemical stability. The material has a thermal expansion coefficient (TEC) of 12 ppm K-1 which is higher than YSZ and provides better compatibility with electrode materials. The above makes it possible to successfully use it as a highly stable oxygen electrolyte or an intermediate thin layer at the electrolyte-electrode interface in electrochemical devices. Keywords solid oxide electrolyte, pyrochlores, grain boundary conductivity, nanoscale powders, combustion method
Acknowledgements The research was partially supported by the Russian Science Foundation (Grant No. 22-23-01121).
For citation Vorotnikov V.A., Belyakov S.A., Ivanov A.V., Novikova Yu.V., Stroeva A.Yu., Grebenev V.V., Khmelenin D.N., Emelyanova O.V., Plekhanov M.S., Kuzmin A.V. Equilibrium of intrinsic and impurity point defects in Ca-doped Sm2Z^O7. Nanosystems: Phys. Chem. Math., 2024, 15 (1), 65-79.
1. Introduction
In recent years, high thermal, chemical and radiation stability has led to great interest in lanthanide zirconates Ln2Zr2O7 with a pyrochlore structure. Due to their good thermomechanical characteristics, compositions such as Gd2Zr2O7, and (La1-xGdx)2Zr2O7 are recommended as thermal barrier coating materials [1-6], which is one of the promising applications of pyrochlores. Having high chemical stability, pyrochlores are used as electrolytes for oxygen sensors for aggressive environments, such as Li-based melts [7-9]. High radiation resistance allows pyrochlores to be used as materials for the encapsulation of nuclear waste [10,11]. Pyrochlores are also considered as promising oxygen-ion conductors for solid oxide fuel cells (SOFCs), so a large amount of work is aimed at studying the relationship between structure and transport properties [12-14].
The mechanism of disorder in pyrochlores is based on its own oxygen deficiency. Partially disordered pyrochlore phases are better ionic conductors than highly disordered fluorite materials of the same composition due to lower activation energies (EA) for the migration of oxygen ions. For example, EA is significantly lower in the pyrochlore-type defective gadolinium zirconate Gd2Zr2O7 than in its fluorite-type analogue [15]. The pyrochlore crystal structure can be derived from the fluorite structure by doubling the unit cell, removing one out of every eight anions, and placing the cations and anions at four crystallographically nonequivalent positions. Using the setup of Tabira et al. [16], the A cation is in the 8-coordinated 16c positions and the B cation is in the 6-coordinated 16d positions (Wyckoff notation). The anions are distributed between two 4-coordinated positions O(48f) and O(8b). One 4-coordinate position O(8b) is free. A decrease in the lanthanide radius in Ln2Zr2O7 leads to a gradual occupation of this position [17-19].
Previously, using atomistic modeling, it was proposed that the most stable intrinsic defect in pyrochlores is the oxygen Frenkel pair, consisting of a vacant O(48f) site and an interstitial ion located at the O(8a) position, and that the diffusion of
© Vorotnikov V.A., Belyakov S.A., Ivanov A.V., Novikova Yu.V., Stroeva A.Yu., Grebenev V.V., Khmelenin D.N., Emelyanova O.V., Plekhanov M.S., Kuzmin A.V., 2024
oxygen ions occurs due to hopping between O(48f) sites [20-24]. Thus, the oxygen conductivity in disordered pyrochlores depends significantly on the energy of formation of this type of defect. The presence of disorder in the cationic sublattice leads to a decrease in this energy value, since cationic disorder increases the similarity between nonequivalent oxygen sites and thereby contributes to the formation of defects in the oxygen sublattice [21,24-28].
Among the zirconate pyrochlore series, Gd2Zr2O7 has the greatest internal disorder [29], demonstrating high conductivity values of about 0.01 Scm-1 at fairly high temperatures ~ 900 °C [30-34]. However, the most disordered pyrochlores are not necessarily the best conductors of oxygen ions and exhibit systematically higher EA for oxygen ion migration. Due to the increase in EA, the conductivity at lower temperatures drops significantly. Thus, the highest conductivity values are usually achieved in partially disordered pyrochlores. This statement can be traced using the example of isovalent doping, when disorder is regulated by varying the ratio of the ionic radii of A and B cations. Doping Gd in Gd2-xMxZr2O7 with larger La or Nd leads to a decrease in the EA and the peak conductivity value reaches Gd1.7Lao.3Zr2O7 and equimolar GdNdZr2O7, respectively [30,31]. In Gd2Sn2_xZrxO7, the conductivity increases with the concentration of Zr which is slightly larger than Sn although the EA changes slightly [32]. For Gd2 Zr2_xTixO7 pyrochlores, a decrease in the EA is observed with increasing content of smaller Ti as the structure becomes more ordered [33]. A decrease in the average size of cations in the A-position, the replacement of Gd by smaller lanthanides such as Dy, Y or Er, increases the intrinsic structural disorder of Gd2Zr2 O7 and induces the pyrochlore-fluorite phase transition [34]. In this case, the replacement of Gd with larger cations, such as Sm, Nd or La, orders the pyrochlore structure.
Heterovalent acceptor-type doping is a strategy to significantly increase oxygen deficiency, as well as cationic disorder. Those dopants that cause the least deformation of the crystal lattice of the base material have the lowest dissolution energy and give the highest conductivity of materials, therefore the replacement of lanthanides with Ca in Ln2Zr2O7 pyrochlores is usually most advantageous [21,23,35]. Against the background of the fairly well-studied mechanism of acceptor doping in both ordered La2Zr2O7 [23, 36-43] and noticeably disordered Gd2Zr2O7 [7-9,44-46,48], moderately disordered Sm2Zr2O7 has been less studied, although it should combine the advantages of the previous two. Liu et al. [49-51] even made attempts to take advantage of equimolar GdSm1-xMxZr2O7_a (M = Ca, Mg) compositions. Without the dopant, the conductivity reaches 0.001 Scm-1 at 600 °C, but decreases for both dopants due to the increasing Ea. Sm and Ca have very similar ionic radii [52]. Despite this, it was reported that the solubility of Ca in Sm2Zr2O7 is low [46, 53], although the mentioned authors used a solid-phase synthesis method. The conductivity of the materials at 600 ° C is about 410-4 S cm 1 and decreases slightly with increasing concentration of Ca dopant, and EA also increase. Our study reveals a noticeable increase in conductivity at low (2.5 at. %) Ca dopant concentrations in Sm2-xCaxZr2O7_^ without a significant effect on the EA.
The microstructure of the samples also affects the level of conductivity; the electrolyte materials must be of high density and approximately equal grain sizes so that the processes occurring in the bulk and at the grain boundary are more understandable. For example, in [46], Sm2_xCaxZr2O7-a ceramics obtained by mechanical activation have a relative density of 89 - 92.6 % and a strong variation in grain sizes from 100-300 nm to several microns is recorded. In work [47], we synthesized the proton conductor La1-xSrxScO3_a by the following methods: solid state reaction, combustion and codeposition. It has been shown that only the combustion method leads to the production of highly dispersed powders with particle sizes up to 100 nm. As a result, this method proved to be the most effective in obtaining a given microstructure of ceramic samples. Hence, a high density relative to the theoretical one (> 95 %) and a uniform size distribution of ceramic grains have been achieved. The influence of the defect formation model and grain boundary effects on the conductivity of materials is also discussed in detail. This work is devoted to the synthesis of nanostructured samarium zirconate powders, the production of dense ceramics and the study of its physico-chemical properties.
2. Experimental 2.1. Sample preparation
The citrate-nitrate combustion technique was used for the preparation of Sm2_xCaxZr2O7-a (where x = 0; 0.05; 0.075; 0.01; 0.125 and 0.15, then SZ; SCZ5, SCZ7.5, SCZ10, SCZ12.5 SCZ15, total - SCZ) samples. Samarium oxide Sm2O3, calcium carbonate CaCO3 and zirconium oxynitrate ZrO(NO3)2, all high purity reagents, were used as precursors. To achieve high homogenization of oxide-forming cations, a solution mixing stage was carried out. For this purpose, a solution of ZrO(NO3)2 with the precise concentration was prepared. The calculated Sm2O3 weights, taking into account the mass loss coefficient during calcination and the drained CaCO3 in a proportional amount of CaO, were transferred to a nitrate solution when interacting with nitric acid (high purity reagent) and an aliquot of zirconium oxyni-trate solution was added. Further, citric acid (high purity reagent) was added to the solution in a ratio of 3/1 to the mass of the resulting oxide and evaporated at a temperature of 130 °C until the combustion reaction took place. The obtained powders were annealed at a temperature of 800 °C (1 hours) to remove the organic residues and carbon black. After this, additional homogenization of the powdered mixture was carried out in an isopropyl alcohol medium using a PM 100 (Retch GmbH, Haan, Germany) planetary ball mill for 1 hour (350 rpm). The dried powders of all compositions
were subjected to preliminary synthesis at a temperature of 1100 °C for 2 hours. After that, the samples were formed by isostatic pressing at 600 MPa and then sintered in air at 1600 °C for 5 hours in a backfill of the appropriate composition.
2.2. Exposure to various atmospheres
The humidity of gas passing through the samples was set using a thermostated vessel with water of a room temperature. The dry atmosphere was reached by a gas passage through synthetic zeolites.
The samples were calcined before analyzes at 1000 °C for 20 hours when blown with dry air, and also at 600 °C without holding when blown in laboratory air. The stability to reduction was checked by holding samples at 850 °C for 250 hours in a flow of dry hydrogen.
2.3. X-ray diffraction
The phase composition of the samples was characterized by a X-ray diffractometer TDM-20 (Tongda, Hong Kong, China) with Cu Ka (A = 1.5418 A) radiation with step scan 0.0095° (29 = 12 - 135°, exposure 1.3 s). The cubic lattice parameters were refined from XRD patterns by Le Bail method. The initial symmetry group and lattice parameters were taken for Sm2Zr2O7 (ICSD 261415).
2.4. Electron microscopy
The surface morphology and cross section of the ceramics were studied by scanning electron microscopy (SEM). The cross sections were performed by grinding and polishing samples with water-based diamond abrasives. SEM images of the cross sections were obtained on a microscope JSM-6510 LV (JEOL, Tokyo, Japan), and the elemental distribution was studied by energy-dispersive X-ray (EDX) spectroscopy using an Inca X-MAX spectrometer system (Oxford Instruments, Abingdon, UK).
In order to investigate microstructure and elemental content of selected samples at subgrain scale, transmission electron microscopy (TEM) study was carried out using OSIRIS TEM (Thermo Fisher Scientific, Waltham, USA) equipped by high-angle annular dark-field (HAADF) detector (Fischione, Corporate Circle, Export, USA) and energy-dispersive spectrometer Super-X EDXS (Bruker, Billerica, USA). TEM data processing was performed in Digital Micrograph 3.4 (Gatan Inc., Pleasanton, USA) and Esprit (Bruker, Billerica, USA) software. The cross-sectional TEM samples were prepared by means of focused ion beam FIB lift-out technique using Scios Dual Beam (Thermo Fisher Scientific, Waltham, USA).
2.5. Dilatometry
Measurements of the thermal expansion were carried out in a quartz cell using the Tesatronic TT-80 (TESA, Renens, Switzerland) meter with a high-precision measuring probe GT 21HP (TESA, Renens, Switzerland). Before the experiment, the cell was calibrated to deduct the contribution of quartz's own expansion. Thermal expansion measurements were implemented during heating/cooling (speed 120 °C/hour) in the temperature range of 50 - 900 °C in moist (pH2O = 2.8 kPa) and dry (pH2O < 0.1 kPa) air. Thermal expansion coefficients (TEC) were calculated from the linear sections of the AL/L0 - T dependencies. Humidity induced chemical expansion was measured during the transition from dry atmosphere (pH2O < 0.1 kPa) to humidified atmospheres (pH2O = 0.6 and 2.8 kPa) at isothermal exposure (T = 600 °C).
2.6. DC and AC conductivity
The total electrical conductivity of ceramic samples was measured by a four-probe DC method using a RM3545-02 (Hioki, Nagano, Japan) ohmmeter. The prepared samples had the geometry of rectangular bars with dimensions of about 12 x 3 x 3 mm. Pt paste and wires were used as electrodes, which was baked to the faces of the samples at a temperature of 1100 °C during 1 hour. The measurements were carried out in the temperature range of 400 - 900 °C in heating/cooling modes in increments of 20 °C. An isothermal exposure at each temperature until the equilibrium resistance values were reached.
The electrochemical impedance spectroscopy (EIS) was applied in a two-probe mode. The samples were tablets with a diameter of 10 mm and a thickness of 1 mm. Ag paste was applied by painting in the form of a 5 by 5 mm area and baked at 800 ° C for 2 hours. In order to reduce the polarization resistance of the electrodes, they were impregnated with a solution of cerium and terbium nitrates and slowly heated to 800 °C at a rate of 20 °C/hour, according to the method of [54]. The measurements were implemented on a SP-200 (Bio-Logic, Seyssinet-Pariset, France) equipment in the temperature range 300 - 400 °C with the step of 25 °C, and in the frequency range 3 MHz - 0.1 Hz at an AC voltage amplitude of 100 mV. The distribution of relaxation times (DRT) technique was employed for the analysis of the spectra [55]. The Tikhonov regularization method was applied [56]. The accuracy of the selection of the regularization coefficient was based on a comparison of the DRT functions of the experimental spectrum and the simulated DRT spectrum from the equivalent circuit [57].
3. Results and discussion
3.1. Crystal structure and morphology
Figure A1 shows typical XRD patterns for SCZ samples. Fig. 1 shows the values of unit cell parameters for samples of two series, as well as literature data. Samples of the first series were calcined in dry air at 1000 °C for 20 hours. Samples of the second series were calcined in laboratory air at 600 °C without holding. One can see a noticeable scatter of parameters for the two series, as well as for the literature data. Thus, heat treatment conditions affect the value of the unit cell parameter. It is known that the formation of the pyrochlore phase passes through the fluorite phase, therefore the unit cell parameter depends on the annealing temperature [58]. First, the synthesis method influences. In all the cited literature references, except for [59], the materials were obtained by the solid-phase synthesis method. Further, the maximum annealing temperature, cooling rate and atmosphere influence. One can compare the annealing conditions using the example of SZ material: in [60], double calcination was used at 1400 °C for 10 hours; in [61], samples were sintered at 1400 - 1600 °C for 8 hours; in [53], there was annealing at 1700 °C for 10 hours; and in [46], there was annealing at 1600 °C for 4-10 hours. In [59], a solution citrate synthesis method was used with final calcination only at 800 °C, although one should expect a much higher combustion temperature for a mixture of citrates and ethylene glycol.
In this work, we chose the method of slow combustion (about 12 hours at a temperature of 130 °C) with the addition of three times the amount of citric acid in relation to the mass of the resulting oxide. This method allows us to obtain powders having the specific surface area of after annealing at 800 °C amounted to 80.000 - 85.000 cm2/cm3 which defines their nano-dimensional character even at this stage of synthesis. At the same time, the powders have already been heat treated, which inevitably led to particle agglomeration. It should be borne in mind that the granulometric composition of powders without heat treatment does not make sense to evaluate due to the presence of additional components, organic residues, and/or the presence of an oxide-forming mixture still in the salt state. The obtained highly dispersed powders led to the production of high-density ceramics (more than 95 % relative to the theoretical one) with a uniform grain size distribution. The final sintering was implemented at 1600 °C for 5 hours. Synthesis and heat treatment conditions primarily affect the concentration of antistructural defects such as Zr in the Sm sublattice and, conversely, Sm in the Zr sublattice [24,26,27]. This can also affect the distribution of the dopant over both sublattices.
Fig. 1. Cubic unit cell parameter of SCZ as a function of dopant concentration. Series 1 means samples pre-calcined at 1000 °C for 20 hours in dry air, Series 2 means samples pre-calcined at 600 °C in laboratory air without holding. Literature results from papers by Shlyakhtina et al. [46], Xia et al. [53], Govindan Kutty et al. [60], Shimamura et al. [61] and Uno et al. [59] are also given.
Figure 2 shows cross sectional images of SCZ ceramic samples. The relative density of SCZ5-SCZ7.5-SCZ10 ceramics varies about 94 %, whereas in samples SCZ12.5 and SCZ15 the density slightly decreases to 91 %. According to the EDX data analysis, all elements in SCZ5-SCZ12.5 are distributed uniformly. Whereas, a clear Ca enrichment is observed in the local areas of SCZ15 sample (Fig. 2c). The presence of such areas indicates the formation of second phase Ca-enriched precipitates. These precipitates have apparently low number density, which is below the detection limit of XRD.
3.2. Thermal expansion
The results of dilatometric measurements of SCZ ceramic samples are presented in Fig. 3. The initial dilatometric dependences are shown in Fig. A2. Average TEC values for Ca-doped samarium zirconates is about 11.510-6 K-1, which is higher than for the conventional oxide-ion electrolyte - YSZ (10.410-6 K-1 [62]). It is known that with a decrease in the ionic radius of the lanthanide in Ln2Zr2O7, the TEC values increase [60,61]. A higher TEC values make it possible to use samarium zirconates as materials of thermal barrier coatings (TBCs) [1-6] which is not available for
Fig. 2. Cross-sectional SEM images of SCZ5 (a), SCZ10 (b) and SCZ15 (c) along with EDX maps of Ca distribution
some other rare earth element zirconates [63]. Also, higher TEC values make the materials more compatible with SOFC electrode materials and similar electrochemical devices [64].
: i • • • • • * •
o • • SCZ dry atm. • SCZ wet atm. O SZ, Z. Qu • YSZ, X. I.iu • ?
0 5 7.5 10 12.5 15
x (Ca), at.% x 2
Fig. 3. Average thermal expansion coefficients (TEC) for SCZ ceramic samples obtained from dilatometry in dry (pH2O < 0.1 kPa) and wet (pH2O = 2.8 kPa) atmospheres. Our experimental data are compared with literature data for Sm2Zr2O7 (SZ) [63] and yttria stabilized zirconia (YSZ) [62].
According to dilatometry data, thermal expansion of SCZ materials does not depend practically on air humidity, and the difference in the TEC in two atmospheres is at the error level. Under stepwise hydration, SCZ5 and SCZ10 samples experience a weak response to a change in humidity, because oxygen vacancies in samarium zirconates are stable to hydration (Fig. A3). Relaxation of chemical expansion occurs abruptly and quickly. It is possible that only partial surface hydration occurs.
3.3. Total conductivity and defects equilibrium
Figure 4 shows the temperature dependencies of the total conductivity for SCZ samples. SCZ5 has a systematically higher conductivity than SZ, while a further increase in Ca concentration generally leads to a decrease in conductivity. Only the conductivity of SCZ10 dominates the others in the temperature range of 800 - 900 0C. The conductivity of all doped SCZ samples is higher than that of undoped SZ at temperatures above 500 0C. Table 1 lists the apparent activation energies (EA). SZ and SCZ5 have comparable EA, while at higher Ca concentrations, the EA increases markedly. In general, the increase in the EA of materials with Ca concentration in Sm2_xCaxZr2O7_<5 is consistent with literature results [46,53]. In our case, the conductivity of SCZ5 at relatively low temperatures of about 600 0Creaches 10-3 Scm-1 which is a record value in the series of oxide pyrochlores, even exceeding the conductivity of the most conducting phases based on doped Gd2Zr2O7 [30-34,49-51].
We did not find a noticeable effect of air humidity on the conductivity of the materials studied, although some researchers have considered these materials as potential proton conductors [46,65]. The air humidity has a slightly visible effect on the conductivity only at temperatures below 700 0C (Fig. 4b). It is difficult to estimate the level of proton conductivity in samarium zirconates according to our data, while we assume that the bulk proton conductivity is very small, and these materials can be successfully considered as oxygen-ion conductors.
Figure 5 shows the conductivity versus pO2 dependence for Sm2_xCaxZr2O7_<5 (x = 0.05, 0.1,0.15). A small effect of pO2 on the conductivity is seen at all temperatures studied. Analogous dependencies can be found in the literature for Gd2_xMxZr2O7_<5 (M = Li, Mg) systems [8,9]. Taken into account the specifics of disorder in the studied pyrochlore-like
Table 1. Apparent activation energies of total conductivity (ETot) for Sm2-xCaxZr2O7-,s (x = 0, 0.05, 0.1, 0.15)
Composition Activation energy of total conductivity (ETot), eV
Dry air, PH2O < 0.1 kPa Wet air, PH2O = 2.8 kPa
Sm2Zr2O7_ä 0.73 —
Sm1.95Ca0.05Zr2O6.975 0.75 0.75
Sm1.9Ca0.1Zr2O5.95 0.94 0.90
Sm1.85Ca0.15Zr2O6.925 1.06 1.04
T,"C
900 800 700 600
1.0 1.2 1.4
1000/T, K'1
-2.0
-2.5
-3.0
-4.0
-4.5
(b)
—- V 900"C
V"' / ■ - Ö 800"C
Q--- ---
. cf / 700"C
X a • ""A 600T
wet atmosphere 8 500°C
0 dry atmosphere
0 5 10 15
x (Ca), at.% x 2
1.2
1.0
<0.9 u
(C)
,y
- dry atmosphere wet atmosphere
0 5 10 15
x (Ca), at.% X 2
Fig. 4. The Arrenius plot of total conductivity for SCZ samples in dry air (a), isothermal dependences of total conductivity on Ca dopant concentration at different temperatures in dry and wet air (b) and apparent activation energies of total conductivity (c). Conductivity values for Sm2Zr2O7 according to Shlyakhtina et al. [46] data are also given.
materials [22,25], several processes affecting the concentration of defects can take place. First, there are oxygen vacancies formed due to doping:
CaO(-SmOi.5) ^ CaSm+^VO'+OO, (1)
where CaSm are Ca atoms in Sm positions, VO' are oxygen vacancies, and OO are oxygen atoms in regular lattice sites. With an increase in temperature, the probability of the formation of anti-Frenkel defects increases according to the equation:
Oo ^ oO+vO', (2)
where OO is the interstitial positions of oxygen. Oxygen vacancies formed by both (1) and (2) reactions can interact with the molecular oxygen:
2o2+vO' ^ OO + 2h. (3)
This creates two electron holes h'. In the case of the dominance of reaction (3), the conductivity should be a linear function of pO2/4. However, at 800 and 700 ° C, the conductivity is linearized only in the plots a = f (pO-1/6) (Fig. 5d), while the a = f (pO2/4) dependence can be found only below 600 °C (Fig. 5e). This indicates the occurrence of several competing processes at temperatures above 700 °C. One of the side processes can be a change in the number of oxygen vacancies, for example, during reaction (2). However, Hagiwara et al. [66] showed by XRD that the occupancy of the O(8a) position in Eu2Zr2O7 at 900 °C is even less than that at room temperature. Shehu [67] also showed by ND method a slight decrease in the occupancy of the O(8a) position at a temperature of 800 °C relative to room temperature for Nd2-xCaKZr2O7-a.
In fact, the O(8a) site in the pyrochlore structure is often called interstitial, although this is not exactly like that. This position is free in ideal pyrochlore, for example in La2Zr2O7 [66]. With a decrease in the ionic radius of the lanthanide, the probability of the formation of a split oxygen vacancy increases, when the oxygen atom is displaced from the O(48f)
Fig. 5. Total conductivities of SCZ as a functions of the pO2 in wet atmosphere for SCZ5 (a), SCZ10 (b), SCZ15 (c). Dependences of total conductivity as a function of pOl/4 for SCZ10 at 700 °C (d) and 600 °C(e).
site towards the O(8a) site, and as a result, the interstitial oxygen atom is between two vacancies in the O(48f) sites [22]. Thus, the oxygen atom never occupies O(8a) site, but only approaches it. If the migration of an "ordinary" oxygen vacancy (at O(48f) or O(8b) sites) is preferable in the crystallographic direction (100), then the split vacancy migrates on average in the direction [111] [22]. In fact, hole migration is also anisotropic and passes along the [110] direction through the Zr and O(48f) positions [19]. This migration direction is perpendicular to the migration direction of ordinary oxygen vacancies. We can assume that the dependence of the total conductivity of SCZ on pO2 at temperatures above 600 ° C is a function not only of the concentration of electron holes but also of the concentration ratio between ordinary and split oxygen vacancies as well as the difference in their mobility due to anisotropic migration.
In contrast to the O(48f) sites, oxygen vacancies in the O(8b) site have a high energy of both formation and migration, as indicated by a number of theoretical studies [20-23]. The oxygen atom O(8b) is surrounded exclusively by A cations. It is to be expected that when the Ca dopant occupies a position in the A sublattice it contributes to the formation of an oxygen vacancy in its immediate environment, including in the O(8b) position. It is worth noting that Shehu [67] did not observe an appreciable deviation from 1 in the occupancy of O(8b) positions in Nd2-xCaxZr2O7_,s by ND even at a very high dopant concentration (x = 0.5). The pair distribution function (PDF) analysis showed a decrease in the average coordination number in the Nd sublattice, but its increase for the Zr sublattice [67]. Thus, the coordination environment of all cations changes [28,67].
It can be seen that the Ca dopant concentration has a progressive effect on disordering in SCZ, in accordance with the ideas developed by van Dijk et al. [15,20]. Formally, the total number of oxygen vacancies increases with increasing Ca concentration which increases the pre-exponential factor of conductivity, but the vacancies are localized in unfavorable crystallographic positions which complicates their migration and increases the activation energy of conductivity.
The presence of a weak but noticeable dependence of conductivity on pO2 in reducing atmospheres confirms the evolution of the number of oxygen vacancies in materials when oxygen leaves the interstitial positions of the material lattice, interacting with hydrogen:
oOO+H2 ^ H2O + 2e'. (4)
As a result, electrons e' are generated for charge compensation, which leads to a slight increase in conductivity. These electrons are probably localized and may be the reason for the partial reduction of Sm3+ to Sm2+. This behavior
is typical for both undoped [68] and Ca-doped samarium zirconate [46]. We additionally tested the stability of SCZ5 sample to reduction through long-term calcination in an H2 atmosphere (Fig. A4). No signs of new phases were found, but the unit cell parameter increases noticeably from 10.583 A after calcination in dry air at 1000 °C for 20 hours to 10.618 A after calcination in dry hydrogen at 850 °C for 250 hours. This may be caused by an increase in the ionic radius of samarium from Sm3+ (r = 1.08 A, CN = 8) to Sm2+ (r = 1.27 A, CN = 8) [52].
The question remains for our case: why does a low concentration of Ca dopant (2.5 %) in Sm2Zr2O7 increase the conductivity of the material and does not influence on the EA. It differs from the results of Xia et al. [53], who received a clear negative reaction? We exclude the influence of sample density, since high-density samples were obtained in both our and Xia et al. [53] works. It is more likely that the difference is due to the specific features of the local structure of pyrochlores where disorder in the anion and cation sublattices is closely related [21,24-28]. The dopant causes a change in the coordination environment in both A and B sublattices [28,67]. Van Dijk et al. [15,20] argue that the oxygen vacancy has the lowest energy when passing between O(48f) positions near the O(8a) position which is essentially identical to the mechanism of migration of a split vacancy. Probably, disruption of such optimal pathways causes a decrease in the mobility of oxygen ions. Unlike the works [46,53], where the SCZ materials were obtained by a solid-phase reaction, in our case, a solution synthesis method with a precursor combustion stage was used. This condition affects the proportion of antistructural defects, as well as the distribution of Ca dopant over the A and B sublattices. It should be expected that the stages of high-temperature combustion and subsequent high-temperature annealing contribute to greater pyrochlore ordering and the formation of a smaller number of antisite pairs, and the solution stage of synthesis promotes the localization of the Ca dopant in the A sublattice. The effect of annealing temperature on the degree of ordering of the pyrochlore structure was shown in the paper [58]. Thus, a small dopant concentration apparently promotes the formation of an increased number of "free" oxygen vacancies and the favorable migration paths are not interrupted. An increased concentration of the dopant causes a violation of the optimal coordination of A and B cations which leads to high energy consumption for the migration of oxygen ions. We believe that an ND study could directly show the occupancy of oxygen positions, unfortunately, ND is difficult for Sm due to its high neutron absorption cross sections. According to ND data for Nd2-xCaxZr2O7_,s [67], the occupancy of the O(8a) position increases noticeably with the dopant concentration and slightly decreases when the material is heated to 800 °C. In addition, Shehu [67] also found a clear increase in the activation energy of conductivity with increasing dopant concentration.
The above allows us to make two general assumptions:
(1) The migration of a split oxygen vacancy, despite the low energy barrier, is disrupted at high defect concentrations. This should also be related to the occupancy of O(8a) positions. A significant change in the coordination environment of cations disrupts the optimal migration routes of oxygen ions.
(2) If the localization of the Ca dopant is preferable in the A position of Sm2Zr2O7, and accordingly, the effect only on coordination in the A sublattice occurs, then this contributes to the formation of energetically favorable oxygen vacancies with low migration activation energy. If part of the dopant ends up in the B position, this will lead to a more significant mismatch of ionic radii between Ca and Zr. The oxygen ion, when migrating, will need to overcome a higher energy barrier in order to change the local coordination of the B cation.
3.4. Grains boundaries state
Since polycrystalline ceramics are under study, one might suspect that the trends in material conductivity discussed above are not related to the equilibrium model of defects in the grain volume, but to the contribution of grain boundary (GB) conductivity. EIS was used to clarify this. Unfortunately, EIS measurements were carried out at lower temperatures compared to the 4-probe method in order to be able to observe the high-frequency response of the impedance spectrum. Fig. 6a shows typical impedance spectra at 350 °C for ceramic samples SCZ5 and SCZ10. A large semicircle emerging from the origin, as well as an additional semicircle, are clearly visible. It should be noted that the analysis may complicate the response from the electrode, so we achieved the lowest resistance of the electrodes by activating them, as was indicated in the experimental section. DRT analysis was used to separate the contributions of relaxation processes on the impedance spectra. Fig. 6b shows the results of DRT analysis at 350 ° C. High-frequency processes with characteristic relaxation times of ~ 10-55 s and capacitances of ~ 10-11 F correspond to the bulk resistance. Low-frequency processes with relaxation times of about 10-4 s and capacitances of ~ 10-9 F can correspond to the GB resistance. The relaxation times of the GB process for SCZ10 are systematically slightly longer than for SCZ5 for all temperatures studied. At the same time, the relaxation times of the bulk process are close for the two samples.
It is worth noting that the resistance of the GB process observed from the impedance spectra is apparent and, before calculating the specific values, it is lower than the bulk one for the both samples. Under the approximation of equality of dielectric constants of bulk and GB [69], the specific GB resistance can be expressed through the ratio of process capacitances:
Rgb = Rgb • n gb . (5) Cbulk
After recalculation, the specific GB conductivity turns out to be noticeably lower than the bulk resistance.
Figure 6c shows the temperature dependences of the bulk and GB conductivity. The dependences of the bulk conductivity are in good agreement with the higher temperature region of conductivity determined by the 4-probe method (Fig. 4a). More importantly, the apparent activation energies of grain volume conductivity are in agreement with higher temperature data which makes our previous discussions about the equilibrium model of defects in the materials under study valid.
Fig. 6. (a) EIS results, (b) DRT spectra and (c) temperature dependencies of bulk and grain boundaries conductivities of SCZ5 and SCZ10 ceramics
To directly confirm or deny formation of single pyrochlore structure at sub-grain scale, we investigate the GB chemistry for SCZ5 and SCZ10 samples. Fig. 7 shows typical bright-field TEM (BF TEM) images and corresponding selected area diffraction (SAED) patterns of SCZ5 and SCZ10 samples. Both samples contained no visible precipitates. In addition, SAED patterns of the SCZ5 and the SCZ10 samples, clearly demonstrate that the local crystal structure of both samples fully matches the cubic phase of Sm2Zr2O7, in the PDF-2 database (ICSD 261415) as follows: the space group Fd3m (227).
Fig. 7. BF TEM images and corresponding SAED patterns of SCZ5 and SCZ10 ceramics cross-sections in the triple junction area
Figure 8 shows the element concentration profiles across the GB for samples SCZ5 and SCZ10 which indicate a uniform distribution of all elements. This contrasts with the results for La2-xCaxZr2O7-a which show clear segregation of the Ca dopant to the GB [43]. Fig. A5 provides the EDX compositional maps of SCZ5 and SCZ10 samples obtained in the triple junction region using the scanning transmission electron microscopy regime (STEM-EDX). Based on the analysis of STEM-EDX compositional maps, both samples under study contain homogeneously distributed Sm, Ca, Zr and O; neither Ca precipitates inside grains, nor Ca segregation at the GB were formed in these samples.
SCZ5 SCZIO
r .,.,.,. 1 " -.-1-.-1-,-1-.-1-.-
U.000 0.004 0.008 0.0L2 0-000 0.004 0.008 0.012 0.016 0.020
Distance, nm Distance, nm
Fig. 8. HAADF images of SCZ ceramics cross-sections of in the region of the triple junction of grains with EDX-profiles across the grain boundaries
Thus, from the totality of the results, it can be expected that the conductivity trends at high temperatures are observed for both bulk and GB. Disorder at GB in pyrochlores is known to occur due to an increased number of antisite pairs, although this effect becomes less significant as cationic disorder in the bulk increases [70]. Apparently, GB, having a similar structure and composition to the bulk, exhibits a similar model of defect equilibrium and only the usual space charge is realized [69].
4. Conclusion
In this work, we obtained high-density (> 95 %) ceramic materials based on Sm2Zr2O7 by combustion method. It was the use of nanoscale pre-ceramic powders that ensured the formation of a given microstructure of the studied samples. We investigated the effect of acceptor doping in Sm2Zr2O7 with a pyrochlore structure on the model of defect formation and conductivity. The low concentration of Ca dopant in the Sm sublattice increases ionic conductivity and almost does not reduce the activation energy (EA), which makes Smi.9sCao.osZr2O7_5 one of the most highly conductive oxide pyrochlores. The coexistence of intrinsic and impurity defects in Sm2_xCaxZr2O7_a gradually impairs ion transport at high doping levels. This leads to high EA due to disruption of the optimal route of the oxygen ions migration which is typical for pyrochlore-type materials. Apparently, the localization of Ca predominantly in the Sm sublattice, which is a consequence of the solution synthesis of materials used, contributes to less distortion of the coordination environment of cations and an increased number of oxygen vacancies, simultaneously. This leads to an increase in conductivity without negatively affecting the EA. Also, the influence of grain boundaries (GB) on the conductivity was tested using impedance spectroscopy. Transmission electron microscopy (TEM) shows the purity of the GB, so it is only the space charge that causes the reduced conductivity of the GB. Sm1.95Ca0.05Zr2O7_5 has been shown to have good chemical stability and a coefficient of thermal expansion (TEC) of 12 ppm K-1 which is higher than YSZ and provides better compatibility with electrode materials. This allows it to be used as a highly stable oxygen electrolyte in electrochemical devices.
Equilibrium of intrinsic and impurity point defects in Ca-doped Sm2Zr2O7 Appendix
i . 1 I . x=0.15 .
1 1 I . x=0.125.
1 . 1 1 > x=0.1 .
1 11. x=0.075 .
». 1 1 . x-0.05 .
1 . A 1> x=0 _ . .
1-I-I-I-r
20 30 40 50 60 70
2Theta, deg.
FIG. A1. XRD patterns for Sm2-xCaxZr2O7-(5 (x = 0; 0.05; 0.075; 0.1; 0.125; 0.15) powders calcined at 1000 °C for 20 hours in dry air
Fig. A2. Temperature dependencies of relative linear expansion and narrowing for Sm2_xCaxZr2O7_,s (x = 0; 0.05; 0.075; 0.1; 0.125; 0.15) dry (a) and wet (b) atmospheres
Fig. A3. Chemical expansion of SCZ5 and SCZ10 samples during the transition from a dry atmosphere (pH2O « 0.1 kPa) to humidified atmospheres (pH2O = 0.6, 2.8 kPa) at a constant temperature T = 600 °C
3
ca
£ "(W
c
o
(b) , 1. 1 1 i . 1
(a) , 1. 1 i . 1
20 30
40 50 60 2Theta, deg.
70
80
Fig. A4. XRD patterns of the SCZ5 sample before (a) and after (b) exposure in hydrogen (250 hours) at 850 °C
Fig. A5. HAADF and SAED cross-sectional images of SCZ5 and SCZ10 samples in the region of the triple junction of grains with STEM-EDX maps of elements distribution
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Submitted 25 November 2023; revised 1 February 2024; accepted 6 February 2024
Information about the authors:
Vladimir A. Vorotnikov - Institute of Chemistry and Ecology, Vyatka State University, Kirov 610000, Russia; Institute of Solid-State Chemistry and Mechanochemistry SB RAS, Novosibirsk, 630090, Russia; ORCID 0000-0002-0247-6198; [email protected]
Semyon A. Belyakov - Institute of High-Temperature Electrochemistry UB RAS, Yekaterinburg 620137, Russia; ORCID 0000-0001-9237-8307; [email protected]
Alexey V Ivanov - Institute of Chemistry and Ecology, Vyatka State University, Kirov 610000, Russia; Institute of SolidState Chemistry and Mechanochemistry SB RAS, Novosibirsk, 630090, Russia; ORCID 0000-0002-7666-831X; [email protected]
Yulia V Novikova - Institute of High-Temperature Electrochemistry UB RAS, Yekaterinburg 620137, Russia; ORCID 0000-0002-0369-3274; [email protected]
Anna Yu. Stroeva - Institute of Chemistry and Ecology, Vyatka State University, Kirov 610000, Russia; ORCID 0000-0002-6772-3321; [email protected]
Vadim V. Grebenev - Shubnikov Institute of Crystallography of Federal Scientific Research Centre "Crystallography and Photonics" of the RAS, Moscow 119333, Russia; ORCID 0000-0002-5484-4642; [email protected]
Dmitry N. Khmelenin - Shubnikov Institute of Crystallography of Federal Scientific Research Centre "Crystallography and Photonics" of the RAS, Moscow 119333, Russia; ORCID 0000-0003-0894-5087; [email protected]
Olga V Emelyanova - Shubnikov Institute of Crystallography of Federal Scientific Research Centre "Crystallography and Photonics" of the RAS, Moscow 119333, Russia; ORCID 0000-0003-3951-8985; [email protected]
Maksim S. Plekhanov - Institute of Crystallography, RWTH Aachen University, Aachen 52066, Germany; ORCID 0000-0002-2701-4619; [email protected]
Anton V. Kuzmin - Institute of Chemistry and Ecology, Vyatka State University, Kirov 610000, Russia; Institute of Solid-State Chemistry and Mechanochemistry SB RAS, Novosibirsk, 630090, Russia; ORCID 0000-0002-0700-662X; [email protected]
Conflict of interest: the authors declare no conflict of interest.
