Proton conductivity of fluorite based rare earth titanates (Ln_xTi_1−x)₄O_8−2x (Ln = Yb, Er, Ho, 0.667 ≤ x ≤ 0.765)†

Nikolay Gorshkov ^a, Egor Baldin ^b, Dmitry Stolbov ^c, Galina Vorobieva ^b, Alexander Shatov ^b and Anna Shlyakhtina *^b
aYuri Gagarin State Technical University of Saratov, Saratov, Russia
^bN.N. Semenov Federal Research Center for Chemical Physics, Russian Academy of Sciences, Moscow, Russia. E-mail: annash@chph.ras.ru; annashl@inbox.ru
^cDepartment of Chemistry, Lomonosov Moscow State University, Moscow, Russia

First published on 6th August 2024

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1. Introduction

The search for solid electrolytes with a high proton conductivity component to make proton-conducting solid-phase fuel cells (PCFCs), which are currently the most promising, is relevant because of the significantly lower operating temperature of PCFCs compared to oxygen-conducting solid-oxide fuel cells (SOFCs). Intermediate-temperature solid oxide devices can be used not only as stationary, but also as mobile energy sources in extreme conditions (lack of electricity, sub-zero temperatures). Despite 40 years of experience in the study of high proton conductivity materials, it has not yet been possible to obtain the values of proton conductivity required for practical use in PCFCs (higher than 10⁻² S cm⁻¹ at 600 °C), moreover, compounds with high proton conductivity have been proven to be unstable in an oxidising or reducing humid atmosphere. Recent years have seen increasing research interest in new, Ba-containing perovskites, aroused by their unusually high proton or oxygen-ion conductivity.^1–7 SOFC devices in which Ba-containing perovskites could be used are intended for operation at high humidity and temperature. Such conditions can lead to rapid degradation of the solid electrolyte material and device performance and, eventually, to failure of the device. The formation of barium hydroxide and carbonates is therefore inevitable during long-term operation under such harsh conditions.⁷ In this context, research into the synthesis of new proton conductors is undoubtedly important. The synthesis of dense ceramics and composites with high oxygen-ion or proton conductivity combined with high electronic conductivity is also of practical interest as membrane materials for separating pure oxygen or pure hydrogen from the products of steam reforming of natural gas and biofuels.

REE (rare-earth elements) titanates with a pyrochlore structure are known to be oxygen ion conductors, as are REE zirconates and hafnates, and have demonstrated high oxygen ion conductivity values. Doping with acceptor impurities significantly increased the conductivity of heavy rare earth titanates.^8–11 In particular, for Yb₂Ti₂O₇ and solid solutions based on it, the oxygen-ion conductivity reached a value of 0.02 S cm⁻¹ at 740 °C.¹¹ As for the fluorite-like “stuffed” pyrochlore Ln₂(Ti_2−xLn_x)O_7−δ (x = 0–0.68), in which some of the titanium positions are replaced by lanthanides, their magnetic properties,^12–14 resistance to radiation effects^15–18 and overall high temperature conductivity^19–21 have been studied. Significant progress has been made in understanding the structure of titanates Ln₂TiO₅, in the context of their investigation as potential materials for nuclear waste disposal.^13,22–24 It has been shown that the cubic and orthorhombic modifications of Ln₂TiO₅ retain its crystallinity under irradiation, in contrast to Ln₂Ti₂O₇ with a pyrochlore structure.²⁵

Polymorphism of Ln₂TiO₅ compounds depends on the ionic radius of the lanthanide.^26,27 It is known that middle rare earth titanates Ln₂TiO₅ (Ln = Ho, Dy, Tb) exhibit the richest polymorphism and can crystallize in three polymorphic modifications: high-temperature cubic (fluorite Fmm), hexagonal (P6₃/mmc) and low-temperature orthorhombic (Pnma). Note that Tb₂TiO₅ and Dy₂TiO₅ are on the stability limit of the cubic phase. For example, Dy₂TiO₅ fluorite was obtained in its pure form by rapid cooling at a rate of 30–50° min⁻¹ from a temperature of 1630 °C.²⁸ It is interesting to note that although Ln₂TiO₅ titanates are fluorites (Fmm) in the long-range order (X-ray diffraction data), electron diffraction studies have shown that they contain nanosized domains of pyrochlore (Fmm) in the short-range order.^14,16,17,29 It is therefore difficult to obtain pure Ln₂TiO₅ (Ln = Ho, Dy, Tb) fluorites due to the rich polymorphism.

They tend to contain an admixture of the hexagonal phase and/or a nanosized phase of pyrochlore.²⁶ Cubic fluorite phases are characteristics of most titanates from the end of the rare earth series Ln₂TiO₅ (Ln = Er–Lu).²⁶ For titanates of the large rare earths Ln₂TiO₅ (Ln = La–Eu), orthorhombic modification occurs at low temperatures, but sometimes in a mixture with the cubic phase.¹⁹ No compounds with an orthorhombic structure have been obtained for Ln₂TiO₅ (Ln = Ho–Lu).²⁸ It has been found that the larger the lanthanide, the more obvious the presence of the pyrochlore phase in the fluorite matrix.²⁸ The number of pyrochlore domains and their size also depend on the conditions of ceramic synthesis, in particular the cooling conditions. Slow cooling leads to the growth of pyrochlore or hexagonal phase domains in the fluorite matrix, which can be identified by XRD.^17,30

On studying the pyrochlore–fluorite disorder in the series of “stuffed” pyrochlore from Dy₂Ti₂O₇ to Dy₂(Dy_0.677Ti_1.33)O_6.677 by ion-irradiation and transmission electron microscopy characterization methods, it was discovered that the initial pyrochlore Dy₂Ti₂O₇ was also partially disordered, i.e. contained fluorite nanodomains.²⁸ Thus, most compounds of the Ln₂TiO₅ series (Ln = Tb–Lu) are typically multiphase and contain a mixture of polymorphic modifications in the form of nano-sized^13–15 or microcrystalline phases.²⁸

Among the rare earth titanates studied, proton conductivity is only observed in acceptor-doped samarium titanates with a pyrochlore structure: Sm₂Ti_1.92Y_0.08O_7−δ and Sm_1.92Ca_0.08Ti₂O_7−δ.³¹ We emphasize that the least explored physical properties of solid solutions with Ln₂O₃ concentrations above 50% in the literature, i.e. the interval for studying different properties is limited to the Ln₂(Ti_2−xLn_x)O_7−δ range (x = 0–0.68). Recently, in an impedance spectroscopy study of the series of Tm₂(Ti_2−xTm_x)O_7−δ (x = 0, 0.1, 0.18, 0.28) and (Tm_xTi_1−x)₄O_8−2x (x = 0.684) prepared by co-precipitation followed by high temperature firing, proton conductivity was discovered for a fluorite-like composition with a high degree of titanium substitution by thulium (Tm_xTi_1−x)₄O_8−2x (x = 0.684) (52% Tm₂O₃).³² The proton transfer coefficients were ∼0.65 in the temperature range of 25–450 °C and decreased to 0.1 as the temperature increased to 700 °C. Thus, like solid solutions based on the pyrochlore Sm₂Ti₂O₇, (Tm_xTi_1−x)₄O_8−2x (x = 0.684) has a significant proton contribution in the range of 300–400 °C. Note that “stuffed” pyrochlores Tm₂(Ti_2−xTm_x)O_7−δ (x = 0, 0.1, 0.18, 0.28) with a low degree of thulium substitution for titanium had no proton contribution. Therefore, there is a need to study the proton conductivity of the whole series of rare earth titanates with a fluorite-like structure. In this work, titanates of medium and heavy lanthanides (Ln_xTi_1−x)₄O_8−2x (Ln = Yb, Er, Ho; x = 0.684) are studied in dry and wet air. In the holmium system, a series of (Ho_xTi_1−x)₄O_8−2x (0.667 ≤ x ≤ 0.765) solid solutions are synthesized and studied to identify the main trend in the change of the proton component with increasing holmium oxide content. All ceramics were obtained by the co-precipitation method, which allows the synthesis temperature of titanates to be significantly reduced compared to the solid phase synthesis method.^32,33 In accordance with literature data, the purest fluorites are expected to be heavy lanthanide titanates (Ln_xTi_1−x)₄O_8−2x (Ln = Er, Yb; x = 0.667).²⁸

2. Experimental

Precursors for the synthesis of (Ln_xTi_1−x)₄O_8−2x (x = 0.684; Ln = Yb, Er, Ho) ceramics and 6 solid solutions (Ho_xTi_1−x)₄O_8−2x (0.667 ≤ x ≤ 0.765) were prepared by coprecipitation. The starting materials used were Ho₂O₃ (99.99%, HoO-1, purity standard TU 48-4-182-74, Russia), Er₂O₃ (TU 48-4-199-72, Russia), Yb₂O₃ (TU 48-4-108-72, Russia) and TiCl₄ (OSCh 12-3, TU-6-09-2118-77, Russia). Ln₂O₃ was dissolved in hydrochloric acid and the titer of the solution was determined gravimetrically. The titanium-containing starting reagent used was TiCl₄ dissolved in concentrated hydrochloric acid. The titer of this solution was also determined gravimetrically. Coprecipitation from the lanthanides and titanium solutions in hydrochloric acid was performed at pH 11 using aqueous ammonia as a precipitant. The resultant precipitates were centrifuged and repeatedly washed with water. The precipitates were dried in a drying oven at 105 °C for 24 h and then decomposed at 650 °C for 2 h, following which the precursors were pressed at 140 MPa and fired at 1600 °C for 4 h. After the firing, the samples were furnace-cooled.

X-ray diffraction spectra of ceramics ground into powder were obtained at room temperature using the Rigaku Smartlab SE X-ray diffractometer (Cu Kα radiation, λ = 1.5418 Å, Bragg-reflection geometry, 40 kV, 50 mA; 2θ range was 10° to 70°, scan step 0.01°, scan rate 5° min⁻¹) in continuous mode. Rietveld refinement was carried out using the SmartLab Studio II software.

The local structure of ceramic materials under investigation was studied by Raman spectroscopy using a SENTERRA Raman microscope spectrometer (Bruker) with excitation at 785 nm.

The microstructure of the ceramic samples was examined using scanning electron microscopy (SEM) on a JEOL JSM-6390LA.

Thermogravimetric studies were carried out on a simultaneous thermal analysis device STA 449C (“NETZSCH”, Germany) at temperatures up to 1000 °C. The sample heating rate was 10 °C min⁻¹. Measurements were carried out under an oxygen atmosphere on a sample with a fluorite structure (Ho_xTi_1−x)₄O_8−2x (x = 0.75): before and after exposure to distilled water at room temperature for one day and two weeks.

Samples for impedance measurements were prepared in the form of discs with a diameter of 7.5 to 8.0 mm and a thickness of 0.7 to 1.5 mm, which were sintered at a temperature of 1600 °C for 4 hours. Pt paste (“ChemPur”, Germany) was used for the contacts. Impedance measurements under excitation of an AC electric field with an amplitude of 100 mV were performed using an impedance meter (Alpha-A Novocontrol, Novocontrol Technologies GmbH & Co. KG, Germany) in the AC frequency range of 10⁻¹–10⁶ Hz. Measurements were made in 50 °C increments over a range of 250–800 °C with isothermal exposure of the sample at each point for 2 hours to achieve thermal equilibrium. Impedance measurements under a dry atmosphere were accompanied by blowing the cell with dried air with a relative humidity of less than 1% at 25 °C, and in a wet atmosphere by blowing with humidified air with a relative humidity of 91% at 25 °C. The parameters of equivalent circuits were constructed using the EIS Spectrum Analyser program (Research Institute of Physical and Chemical Problems, Belarus) using the Powell algorithm.³⁴

3. Results and discussion

3.1. Sample characterization

Fig. 1 shows the diffraction patterns of fluorite-like solutions (Ln_xTi_1−x)₄O_8−2x (Ln = Yb, Er, Ho; x = 0.684;), which initially contain 52 mol% Ln₂O₃ and 48 mol% TiO₂. It is clear that the purest fluorite, as expected from the data in ref. 26 and 27 was formed in the ytterbium system. It can be seen that in the diffraction patterns of erbium and holmium solid solutions of similar composition, diffraction lines of other phases are present in addition to fluorite. Note that in all diffraction patterns of (Ln_xTi_1−x)₄O_8−2x (Ln = Yb, Er, Ho; x = 0.684) ceramics, the main phase is fluorite (F) (Fig. 1). Erbium and holmium titanates also contain a pyrochlore phase (P) with a double fluorite unit cell. The unit cell parameters of the main cubic phase F decrease in the Ho–Er–Yb series (Table 1). The table shows the doubled fluorite cell parameter a = 2a_F for easy comparison with pyrochlore parameters. The hexagonal phase β-Ln₂TiO₅ (P6₃/mmc) is present in small amounts in erbium and holmium titanates. The parameters of the hexagonal phase are: a = 3.6266(9) Å, c = 11.881(6) Å for Er titanate and a = 3.6148(3) Å, c = 11.927(19) Å for Ho titanate. In the (Ho_xTi_1−x)₄O_8−2x (x = 0.684) sample, a second cubic phase of pyrochlore (denoted P2 in Fig. 1) with a smaller unit cell parameter a = 10.0987(1) Å was also detected, which may correspond to stoichiometric pyrochlore Ho₂Ti₂O₇a = 10.1048(2) Å.¹² Similar but much broader lines are present in samples containing Er and Yb. It can be assumed that in these samples a small amount of stoichiometric pyrochlore Ln₂Ti₂O₇ exists in a nanosized/X-ray amorphous state. There are works on thulium and ytterbium systems (Ln₂O₃–TiO₂ (Ln = Tm, Yb)),^35,36 where the authors observed the separation of the pyrochlore phase into two: P1 + P2. In this work, a multiphase system was also obtained, which is consistent with the observation made previously in ref. 28 of the increase in the amount of pyrochlore phase and the growth of its nanodomains in Ln₂TiO₅ fluorites with an increase in the ionic radius of the lanthanide. Thus, although in some cases Er₂TiO₅ ceramics have been found to form pure fluorites by X-ray diffraction,^21,30 they are all metastable phases and contain nanodomains of pyrochlore and/or hexagonal phase, which are often undetectable by this method. Note that a small number of unidentified peaks remain on the diffraction pattern, but their relative intensity and therefore the proportion of this unidentified phase will be extremely small (less than 1%).

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Fig. 2 shows the diffraction patterns of a series of holmium titanates (Ho_xTi_1−x)₄O_8−2x (0.667 ≤ x ≤ 0.765). The (Ho_xTi_1−x)₄O_8−2x (x = 0.667) or Ho₂TiO₅ compound is a mixture of two phases: (1) hexagonal phase β-Ho₂TiO₅ with cell parameters a = 3.6135(4) Å, c = 11.91642 Å, ≈75 wt% and (2) cubic fluorite phase ≈25 wt%. Unfortunately, the presence of the pyrochlore P phase in this sample cannot be accurately assessed because its characteristic peaks (111) and (311) overlap with the hexagonal phase peaks and the next isolated peak (331) is only slightly above noise. In the holmium system, according to preliminary studies,²⁶ it is impossible to obtain pure fluorite Ho₂TiO₅ and it will always contain an admixture of the hexagonal phase and/or pyrochlore. The (Ho_xTi_1−x)₄O_8−2x (x = 0.684) sample was described in the previous paragraph. With increasing holmium content in the titanate, an increase in the unit cell parameter of the main fluorite phase is observed (Table 2). The weak lines of pyrochlore P in the diffraction pattern (Ho_xTi_1−x)₄O_8−2x (x = 0.701) are significantly broadened, indicating that this phase is nanosized. From the width at half maximum of the (311) peak, the size of the crystallites of the pyrochlore phase was estimated using the Scherrer formula and was 46.7 ± 5.5 nm. A further increase in the holmium content of the titanate leads to a decrease in the intensity of the diffraction lines of the nanopyrochlore. In the diffraction pattern of the (Ho_xTi_1−x)₄O_8−2x (x = 0.750) or Ho₃TiO_6.5 sample, the nanopyrochlore lines are completely indistinguishable and it can be argued that it is pure fluorite in the long-range order. The sample (Ho_xTi_1−x)₄O_8−2x (x = 0.765) also contains no traces of nanopyrochlore, but Ho₂O₃ impurities with a bixbyite structure were found in its composition (in Fig. 2 the corresponding peaks are marked with the symbol C). Thus, (Ho_xTi_1−x)₄O_8−2x (x = 0.750) is the only composition that crystallizes to pure fluorite from co-precipitated precursors after annealing at 1600 °C. Rietveld refinement of the XRD pattern of (Ho_xTi_1−x)₄O_8−2x (x = 0.684) and (Ho_xTi_1−x)₄O_8−2x (x = 0.75) (Fig. S1a and S1b†) is presented in the ESI.†

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All prepared holmium titanates (Ho_xTi_1−x)₄O_8−2x (0.667 ≤ x ≤ 0.765) as well as the sample (Yb_xTi_1−x)₄O_8−2x (x = 0.684) were investigated by Raman spectroscopy (Fig. 3). As the holmium content of the titanates increases, the bands become wider and individual peaks are difficult to distinguish. The ordering of the Ho and Ti cations, characteristic of the pyrochlore structure, is lost and the structure becomes disordered fluorite, not only in the long-range order but also in the short-range order. At the same time, the spectra of (Ho_xTi_1−x)₄O_8−2x (x = 0.750) and (Yb_xTi_1−x)₄O_8−2x (x = 0.684), which are the purest fluorites according to the XRD data, practically duplicate each other. An additional narrow band can be seen in the spectrum of (Ho_xTi_1−x)₄O_8−2x (x = 0.765) at 372 cm⁻¹. This band is characteristic of oxides with a bixbyite structure and belongs to Ho₂O₃.³⁷

Unexpectedly, in the Raman spectra of the ceramics (Ho_xTi_1−x)₄O_8−2x with x = (0.677, 0.684, 0.701, 0.718, and 0.734), the pyrochlore bands were found to be weakly expressed, while the origin of most of the strong lines remained unclear.

The main characteristic Raman bands of the pyrochlore single crystal Ho₂Ti₂O₇ are located at 311.09 cm⁻¹ (vibration (F2g) associated with Ho–O stretching) and 522.16 cm⁻¹ (vibration (A1g) associated with Ti–O stretching) according to ref. 38, where Raman spectra of single crystals of the whole series of pyrochlore rare earth titanates RE₂Ti₂O₇ (RE = Sm–Lu, Y) were obtained. In Fig. 3 we observe in holmium ceramics a consistent extinction of bands close to the above-mentioned ∼315 and ∼500 cm⁻¹ with increasing Ho₂O₃ content in samples (Ho_xTi_1−x)₄O_8−2x with x = 0.667, 0.684, 0.701, and 0.718. According to ref. 39, where the Raman spectra of Y₂Ti₂O₇ pyrochlore were studied (the ionic radii of Ho³⁺ and Y³⁺ in the octahedral environment are close: R Ho_CN=8³⁺ = 1.015 Å; R Y_CN=8³⁺ = 1.019 Å), a weaker band at ∼225 cm⁻¹ appeared, also associated with Y–O stretching (F2g). In Fig. 3, for the (Ho_xTi_1−x)₄O_8−2x with x = 0.667, 0.684, 0.701, and 0.718 cm⁻¹, we note the consistent extinction of a similar band at ∼240 cm⁻¹. For ceramics with a higher Ho content, (Ho_xTi_1−x)₄O_8−2x (x = 0.734, 0.75, 0.765), all pyrochlore bands disappear completely.

An increase in the Ho content in (Ho_xTi_1−x)₄O_8−2x with x = (0.667–0.734) solid solutions leads to a slight broadening of the main Raman bands and a slight shift of all spectral lines, but the spectrum changes dramatically for (Ho_xTi_1−x)₄O_8−2x with x = (0.75), which is pure fluorite according to X-ray diffraction data (Fig. 2). It is clearly seen that the splitting into two bands at 716 and 759 cm⁻¹ for the compositions (Ho_xTi_1−x)₄O_8−2x with x = (0.667–0.734) disappears completely during the formation of pure fluorite (Ho_xTi_1−x)₄O_8−2x (x = 0.75). In this case a broadband is formed with a maximum at 743 cm⁻¹. Similar behavior has been observed for fluorite-like compositions in the Y₂O₃–TiO₂ system.⁴⁰ With increasing Y₂O₃ content, the characteristic splitting of the Raman bands of fluorite Y₂TiO₅ into two components at 727 and 760 cm⁻¹ disappears for (Y_xTi_1−x)₄O_8−2x (x = 0.75). One of the reasons for such splitting has been suggested in ref. 40 as residual internal stresses in the unit cells of related phases when they coexist in micro- and nanodomains. In addition, it should be noted that the strong bands at 292, 353, and 397 cm⁻¹ characteristic of the compositions (Ho_xTi_1−x)₄O_8−2x with x = (0.667–0.734) also completely merge into one broadband for pure fluorite (Ho_xTi_1−x)₄O_8−2x (x = 0.75) (Fig. 3). Thus, the typical Raman spectrum of pure fluorites (Ho_xTi_1−x)₄O_8−2x (x = 0.75) and (Yb_xTi_1−x)₄O_8−2x (x = 0.684) consists of 2 broad bands in the region 240–459 cm⁻¹ and 670–800 cm⁻¹ with a pronounced maximum at 743 cm⁻¹. Strong Raman bands at 143, 166, 203, 292, 353, 397, and 743 cm⁻¹ apparently belong to the hexagonal phase (Ho_xTi_1−x)₄O_8−2x (x = 0.667). This is in agreement with the X-ray diffraction data (Fig. 2). Probably, the local hexagonal phase is also preserved in Ho₂O₃-enriched compositions (Ho_xTi_1−x)₄O_8−2x (x = 0.684, 0.701, 0.718, 0.734).

SEM images illustrating the surface morphology of the ceramics discussed are shown in Fig. 4. The surface microstructure of erbium ceramics with the composition (Er_xTi_1−x)₄O_8−2x (x = 0.684) is similar to the microstructure of thulium ceramics with the same composition previously investigated.²⁵ The grains are elongated in one direction and lie close together. Ytterbium ceramics are different from erbium and thulium ceramics, but similar motifs can be seen in them. All holmium ceramics, with the exception of (Ho_xTi_1−x)₄O_8−2x (x = 0.667), which is predominantly hexagonal according to XRD data, are characterised by the formation of large crystallites of the order of 20–40 μm. For two samples of the holmium series (Ho_xTi_1−x)₄O_8−2x (x = 0.701 and 0.75) with high total conductivity (Part 3.3.), extended defects are visible within large crystallites.

On the surface of the (Ho_xTi_1−x)₄O_8−2x (x = 0.765) sample, which contains the highest amount of holmium among those examined, smaller grains of ≈2–5 μm can be observed between and inside large grains. It can be assumed that they belong to the Ho₂O₃ detected in the XRD data. While in all the other samples examined the Ho and Ti cations are evenly distributed, for (Ho_xTi_1−x)₄O_8−2x (x = 0.765) the X-ray mapping shows places where there is no titanium but holmium (Fig. S2†). Notable features of all single-phase materials include low porosity.

3.2. Proton conductivity of ceramics (Ln_xTi_1−x)₄O_8−2x (Ln = Yb, Er, Ho; x = 0.684)

Fig. 5 shows typical impedance spectra for (Ho_xTi_1−x)₄O_8−2x (x = 0.684) at temperatures of 400 °C and 450 °C under dry and wet atmospheres. Hodographs have the general appearance of overlapping large and small semicircles. In general, the impedance of such materials consists of three components: the grain volume impedance, the grain boundary impedance and the electrode contribution. Electrode processes are not the subject of this study and can only be reliably assessed at high temperatures. The equivalent circuit for interpreting the impedance data consists of two serial sections with parallel CPE and resistance: (R_bCPE_b), (R_gbCPE_gb), where R_b is the bulk resistance; R_gb is the grain boundary resistance; CPE_b and CPE_gb are constant phase elements. For the first arc, the capacitance values were in the order of 10⁻¹¹–10⁻¹⁰ F, which can be interpreted as the grain bulk contribution, and the second arc with a capacitance of 10⁻⁹–10⁻⁸ F corresponds to the grain boundary contribution. The grain bulk has a greater contribution to the overall impedance and the effect of grain boundaries is negligible for the overall conductivity. It is interesting to note that under a wet atmosphere the grain bulk/grain boundary ratio increases with increasing x, whereas under a dry atmosphere it decreases for the (Ho_xTi_1−x)₄O_8−2x (x = 0.684–0.75) fluorite series (Fig. S3c†). At the same time, increasing the temperature and humidity of the atmosphere significantly reduces the overall resistance.

The temperature dependences of the total conductivity of the (Ln_xTi_1−x)₄O_8−2x (Ln = Yb, Er, Ho; x = 0.684) series of ceramics are shown in Fig. 6a. Clearly, conductivity increases with lanthanide radius in both dry and wet air. The proton conductivity is higher for the (Ho_xTi_1−x)₄O_8−2x (x = 0.684) solid solution and decreases successively for (Er_xTi_1−x)₄O_8−2x (x = 0.684) and (Yb_xTi_1−x)₄O_8−2x (x = 0.684). Fig. 6a demonstrates that the proton contribution remains up to 700 °C. A similar dependence has been observed for cubic molybdates with a bixbyite structure Ln₆MoO_12−δ (Ln = Er, Tm, Yb).⁴¹ The proton conductivity contribution exists in cubic Ln₆MoO_12−δ (Ln = Er, Tm, Yb) bixbyites up to 450–600 °C and decreases with the decrease of the lanthanide ionic radius.

When Ti⁺⁴ is replaced by Ln⁺³ (Yb, Er, Ho) cations in solid solutions of rare earth titanates with the formula (Ln_xTi_1−x)₄O_8–2x, oxygen vacancies appear:

	(1)

Here, titanate, as a complex oxide, undergoes classic quasi-chemical reactions with atmospheric oxygen (2) and water vapour (3):

	(2)

	(3)

Reactions (2) and (3) are not independent because the concentrations of holes, protons and oxygen vacancies are related by the condition of electrical neutrality, so the dissolution of water vapour, which reduces the concentration of vacancies, also reduces the concentration of holes. The total conductivity of the studied REE titanates is characterized by a mixed type and can be expressed in the general case as:

σ = σion + σh,	(4)

where σ_ion – ionic conductivity, σ_h – hole conductivity.

Under a dry air atmosphere (pO₂ = 0.21 atm), the ionic conductivity σ^dry_ion is determined by the contribution of the oxygen conductivity, then the total conductivity takes the form:

σdry = σ2−O + σh,	(5)

where σ²⁻_O – oxygen ion conductivity.

One can see that the conductivity values under a wet atmosphere are higher than those under a dry atmosphere. The increase in total conductivity under a wet atmosphere is due to the contribution of the proton conductivity to the ionic conductivity (σ^wet_ion = σ²⁻_O + σ_H), where σ_H is the proton conductivity. At lower temperatures, the differences in conductivity between the dry and wet atmospheres become more pronounced. This indicates a hydration process leading to an increase in proton concentration. The analysis of the impedance data allowed us to obtain data σ^dry under a dry atmosphere and σ^wet under a wet atmosphere. Despite the mutual influence of reactions (2) and (3), the equation has been used to estimate proton transfer numbers, assuming that the difference in hole conductivity under dry and moist environments of atmospheric air can be neglected:

	(6)

A full description of the transport properties of the ceramics studied is given in Table 3. The activation energy of proton conduction in the temperature range of 450–200 °C decreases consistently with increasing ionic radius Ln.

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Fig. 6b shows the temperature dependence of the proton transfer coefficient in these ceramics. At 300 °C the proton transfer coefficients are higher for titanates with a smaller ionic radius. However, the difference is insignificant and amounts to 0.1. At 400 °C they are almost the same for all compositions and amount to 0.8. At 600 °C, a drastic decrease in transfer coefficients is observed for all three compositions.

The reason for the increase in proton conductivity of fluorite-like REE titanates with increasing lanthanide ionic radius may be an increase in the basic properties of the REE cations, in the compounds in which proton transfer occurs much more easily. The holmium cation belongs to the middle series of rare earth elements and its compounds, such as Ho₂Ti₂O₇ pyrochlore, are capable of retaining hydroxide water up to 1200 °C.⁴² For some systems, the composite effect is also known, where the conductivity of pure pyrochlore and fluorite is lower than that of the mixed two-phase region P + F or P1 + P2.^33,43,44

3.3. The trend of changes in the potential proton conductivity of ceramics of the (Ho_xTi_1−x)₄O_8−2x (0.667 ≤ x ≤ 0.765) series

The temperature dependence of the total conductivity for (Ho_xTi_1−x)₄O_8−2x (0.667 ≤ x ≤ 0.765) ceramics is plotted in Fig. 7a. It can be seen that the total conductivity in dry and wet air increases as the holmium content increases. The proton contribution for (Ho_xTi_1−x)₄O_8−2x (x = 0.701, 0.75) is retained up to 800 °C. The activation energy decreases for (Ho_xTi_1−x)₄O_8−2x (0.667 ≤ x ≤ 0.765) ceramics with increasing holmium content from 0.667 to 0.701 and then changes very slightly (Table 3). Fig. 7b shows the proton transfer coefficients for these ceramics. It should be noted that the transfer coefficients for the first two compositions (Ho_xTi_1−x)₄O_8−2x (x = 0.667, 0.684) are close and decrease uniformly up to 600 °C. This is probably due to the fact that fluorites, which are the main conducting phase of these materials, are similar in composition due to the formation of the associated hexagonal phase and pyrochlore, as indicated by the similarity of the parameters of the fluorite phase (Table 2). The proton transfer coefficients for all samples show the following general pattern: high values are observed in the range from 250 to 450 °C, then the values decrease with increasing temperature, reaching values below 0.1 at 800 °C. For the ceramic samples (Ho_xTi_1−x)₄O_8−2x (x = 0.667, 0.684, 0.701, 0.718, 0.734, 0.750, 0.765) the proton transfer coefficients in the low temperature range increase with increasing holmium content from 0.75 to 0.90 with the maximum at x = 0.75, while in the high temperature range the nature of the dependence is similar, with the exception of the (Ho_xTi_1−x)₄O_8−2x (x = 0.701) sample, which shows higher values relative to neighbouring holmium concentrations.

The total conductivity for the ceramics (Ho_xTi_1−x)₄O_8−2x (x = 0.667, 0.684, 0.701, 0.718, 0.734, 0.750, 0.765) measured under a dry air atmosphere (Fig. 7c) increases with the holmium content in the series 0.667, 0.684, and 0.701, with the greatest difference at temperatures from 600 °C to 250 °C. A further increase in holmium to x = 0.750 leads to a decrease in conductivity values, while the order of conductivity values at 800 °C for all (Ho_xTi_1−x)₄O_8−2x ceramics (x = 0.667, 0.684, 0.701, 0.718, 0.734, 0.750, 0.765) ∼1 × 10⁻⁴ S cm⁻¹.

For conductivity in humidified air, the nature of the dependence of the values for (Ho_xTi_1−x)₄O_8−2x ceramics (x = 0.667, 0.684, 0.701, 0.718, 0.734, 0.750, 0.765) on the holmium content is not preserved, so that the increase in conductivity for samples x = 0.667, 0.684, 0.701 is replaced by a decrease in values for sample x = 0.718, then an increase in conductivity is observed with increasing holmium content up to x = 0.750. The ceramic sample (Ho_xTi_1−x)₄O_8−2x (x = 0.765) has high conductivity values in the series depending on the holmium content at temperatures from 450 °C to 800 °C, but in the temperature range from 450 °C to 250 °C the conductivity shows a decrease compared to the sample x = 0.750 (Fig. 7d).

The observed dependence of proton transfer coefficients and conductivity under dry and wet atmospheres on the holmium content for the ceramic (Ho_xTi_1−x)₄O_8−2x (x = 0.667, 0.684, 0.701, 0.718, 0.734, 0.750, 0.765) can be explained by changes in the crystal structure of the samples (Fig. 7e). The conductivity in dry air from 600 °C to 250 °C is the maximum for ceramics that contain nano-sized pyrochlore in addition to the main fluorite phase, i.e. (Ho_xTi_1−x)₄O_8−2xx = 0.701. Conductivity in dry air decreases as the amount of nanopyrochlore decreases. In wet air, however, the opposite effect is observed. The conductivity in wet air compared to dry air increases more significantly the higher the parameter of the main fluorite phase and the less nanopyrochlore in the ceramics. The maximum transfer coefficients at 250 °C are obtained for the compound (Ho_xTi_1−x)₄O_8−2xx = 0.750 or Ho₃TiO_6.5, which we consider to be the purest fluorite. Although the parameter of the fluorite phase of the sample with x = 0.765 is higher than that with x = 0.750, its conductivity is lower under both atmospheres, which is explained by the presence of an impurity Ho₂O₃ phase in it.

The time dependencies of the impedance spectra, which allow us to confirm the proton contribution to the conductivity for the best sample (Ho_xTi_1−x)₄O_8−2x (x = 0.75), are shown in Fig. S3a and Fig. S3b† at temperatures of 300 and 450 °C, respectively. The proton conduction will depend on the rate at which water is absorbed in the sample. Considering the rather low conductivity, diffusion coefficients will be very small. Hence, a slow change might indicate proton conduction.⁴⁵

3.4. Thermogravimetric analysis of (Ho_xTi_1−x)₄O_8−2x (x = 0.75) in dry oxygen

The presence of oxygen vacancies is a necessary condition for the transfer of protons in oxides: under a wet atmosphere, dissociative absorption of water vapour from the gaseous atmosphere (eqn (3)) occurs by oxygen vacancies.

At room temperature, proton-conducting complex oxides also hydrate naturally due to the humidity of the surrounding air. It therefore seems reasonable to investigate weight loss in these materials. However, previous studies of rare earth titanates have shown that, in addition to the dehydration process, there is a process of CO₂ release and carbon combustion discovered in the synthesis from coprecipitated precursors during their heat treatment in a wide temperature range of 800–1600 °C.^46–48 X-ray diffraction and Raman spectroscopy methods have shown that Ln₂Ti₂O₇ (Ln = Gd, Tb, Dy) precursors, e.g. obtained by “wet” chemical methods, are either a mixture of amorphous Ln–Ti hydroxycarbonates⁴⁷ or a mixture of rare earth hydroxycarbonates with titanium hydroxide.⁴⁶ At the same time,⁴⁷ a violation of stoichiometry in high temperature ceramics and the presence of up to 3 wt% graphite were noted.

For powdered ceramics with the highest conductivity (Ho_xTi_1−x)₄O_8−2x (x = 0.75), we studied the change in sample mass under an atmosphere of oxygen with two times heating. The ability to hydrate of (1) the original powder and after exposure (2) for 24 hours and (3) for two weeks was investigated.

Fig. 8 shows the behavior of fluorite (Ho_xTi_1−x)₄O_8−2x (x = 0.75) in dry oxygen when heated to 1000 °C in a STA cell. For clarity, the main figure shows the results of a first heating of the original powder, the same hydrated for 24 hours and the same hydrated for two weeks. Note that the weight loss of the original sample was 0.20%, while the sample hydrated for 24 hours showed a slightly greater weight loss of ∼0.24%. After 2 weeks in water at room temperature, the sample showed a slightly lower weight loss. It is known that during hydration at temperatures below 300 °C, hydroxide and hydroxycarbonate impurities may dissolve.⁴⁸ This may be associated with a slightly lower weight loss after 2 weeks of exposure to water. Therefore, keeping the powder for 24 hours is sufficient to ensure maximum hydration, and a soaking time of at least 2 weeks is required to dissolve most of the impurity hydroxides and hydroxycarbonates in water.

Loss of pore water and weakly bound water has been observed in samples up to 530 °C.

However, it should be noted that in the same temperature range CO₂ is released as a product of the decomposition of the remaining REE hydroxycarbonates;⁴⁶ any carbonaceous impurities are burnt out, which is accompanied by an exo-effect at ∼320 °C (Fig. S4†).

The process of decomposition of basic carbonates and hydroxycarbonates of rare earths in air is a multi-stage one and takes place over a wide range of temperatures. For example, the decomposition of lanthanum hydroxycarbonate LaOHCO₃·nH₂O proceeds as follows: first to La₂O₂CO₃ at a temperature of about 500 °C, and then to La₂O₃ in the range 750–830 °C, according to the scheme:⁴⁹

La2(CO3)3·nH2O → La2(CO3)3 + nH2O↑	(7)

La2(CO3)3 ⇌ La2O2CO3 + 2CO2↑	(8)

La2O2CO3 ⇌ La2O3 + CO2↑	(9)

In air, in the presence of CO₂, where ceramic synthesis takes place, this process is reversible. It is therefore not surprising that traces of carbonates can be preserved in ceramics based on lanthanide-containing complex oxides, synthesized by any method using nano-sized precursors, even after annealing at high temperatures (1200–1600 °C).^46,50

At temperatures T > 530 °C, mass loss occurs mainly due to the release of OH⁻ groups and strongly bound protons. In the inset to Fig. 8, the results of the second heating of the samples are shown: (1) initial powder, (2) after 24 h exposure in water, and (3) after two weeks exposure in water. A gradual decrease in mass loss during reheating was observed: 0.22, 0.18, and 0.16 wt% respectively.

Note, however, that the TG curve of the sample under investigation in oxygen does not reach a plateau even when heated to 1000 °C. It is most likely that the final decomposition of carbonaceous impurities occurs at a higher temperature T > 1000 °C, since in ref. 46 it is shown that titanates-pyrochlores of rare earth elements, for example, tend to retain a small amount of carbon-containing compounds even after high temperature annealing in the range of 1200–1600 °C due to relatively slow crystallization.

4. Conclusions

The study of (Ln_xTi_1−x)₄O_8−2x (Ln = Yb, Er, Ho; x = 0.667, 0.684, 0.701) and (Ho_xTi_1−x)₄O_8−2x (x = 0.667, 0.701, 0.718, 0.734, 0.75, 0.765), prepared by the method of co-precipitation followed by high temperature annealing at 1600 °C, has revealed for the first time the proton component of conductivity in fluorite-like REE titanates with proton transfer numbers of up to 0.86 in the range of 200–450 °C. With a further increase in temperature, the proton conductivity drops quite sharply and the transfer coefficients decrease to 0.3 at 700 °C for the best proton conductors in this series, (Ho_xTi_1−x)₄O_8−2x (x = 0.701, 0.75). The increase in proton conductivity in the Yb–Er–Ho series is associated with an increase in the hydrophilic properties of rare earth cations. However, conducting hydrogen/deuterium exchange experiments under switching H₂O/D₂O atmospheres is essential for future studies of the title compound as a proton conductor.

The structural study by X-ray diffraction and Raman spectroscopy methods showed that the maximum proton conductivity is exhibited by the solid solution (Ho_xTi_1−x)₄O_8−2x (x = 0.701), containing pyrochlore microdomains of optimal size in the fluorite matrix, or pure fluorite with the composition (Ho_xTi_1−x)₄O_8−2x (x = 0.75).

Author contributions

Nikolay Gorshkov – electrical conductivity measurements and interpretation of the rare-earth titanates, and writing of the manuscript; Egor Baldin – synthesis of the rare-earth titanates, XRD calculation, and writing of the manuscript; Dmitry Stolbov – SEM measurements and interpretation; Galina Vorobieva – TG measurements; Alexander Shatov – Raman spectral measurements; Anna Shlyakhtina – idea, synthesis of the rare-earth titanates, interpretation of Raman spectra, and writing of the manuscript.

Data availability

Data are available within the article or its ESI.†

The authors confirm that the data supporting the findings of this study are available within the article [and/or] its ESI.†

Conflicts of interest

There are no conflicts of interest to declare.

Acknowledgements

The work was supported partially by the subsidy from the Ministry of Education and Science allocated by the FRC CP RAS for the implementation of the state assignment no. 122040500071-0 and 122040500068-0.

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Footnote

† Electronic supplementary information (ESI) available. See DOI: https://doi.org/10.1039/d4dt01493f

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