Electrical Properties and Chemical Resistance of the Composites (1-x)Gd₂Zr₂O₇·xMgO in Li-Containing Chloride Melts

Abstract

Composites of (1-x)Gd₂Zr₂O₇·xMgO were prepared by mixing gadolinium zirconate with freshly precipitated Mg(OH)₂ followed by heat treatment at 1500 °C. Small concentrations of magnesium oxide dissolved in the complex oxide matrix of Gd₂Zr₂O₇. This led to decrease in the lattice parameters of the matrix phase and a complex redistribution of Gd and Zr over the A and B sublattices. According to the impedance spectroscopy results of the studied samples, for (1-x)Gd₂Zr₂O₇·xMgO (x = 0.05, 0.07, 0.10), the ionic conductivity was slightly higher than that for the undoped Gd₂Zr₂O₇. The share of dominant ion transport did not change upon doping with magnesium oxide. The composites showed chemical resistance in a lithium halide (LiCl) melt and interacted with LiCl-xLi₂O (x = 2 wt.%, 4 wt.%) melts at 650 °C with the formation of a Gd₂O₃ phase or a mixture of phases (Gd₂O₃, Li₂ZrO₃, ZrO₂, LiGdO₂, or LiGdCl₂) on the ceramic surface, respectively.

1. Introduction

The pyrochemical processing of spent nuclear fuel is an effective way to reduce nuclear waste. Processes such as anodic dissolution, mass transfer, and the electrodeposition of actinides in molten salts have been studied in detail by scientists around the world [1,2,3,4]. The electrochemical reduction of SNF to a metallic form is the main stage of the electrorefining process for LiCl-xLi₂O melts. As a result, actinide metal oxides are deposited on the cathode, while due to the proximity of electrode potentials, metallic lithium is formed on the cathode, which participates in the process of chemical reduction of actinide oxides. To obtain a high-quality product and avoid undesirable processes, such as the dissolution of the platinum anode at a low content of Li₂O in the melt, i.e., ~0.25–0.50 wt.%, the lack of solubility of plutonium and americium oxides at high contents of Li₂O, at ~3.4 wt.% and ~5.1 wt.%, respectively, or the degradation of structural materials, it is necessary to maintain a certain concentration of oxide ions in the working melt [5]. Therefore, devices for the continuous monitoring of the oxygen ion content in the working LiCl-xLi₂O melt environment are needed.

Most of the ion-selective membrane sensors proposed for fabrication are made of ceramic materials, which must be inert in LiCl-based melts containing Li₂O. Materials based on zirconium oxide used in oxide-halide melts are not stable enough to maintain stable operation for a long time. The formation of Li₂ZrO₃ is inevitable when using ZrO₂ in LiCl-xLi₂O melts, and with an increase in the content of Li₂O, the stoichiometry of lithium zirconate changes to Li₆Zr₂O₇ and Li₈ZrO₆ [6,7,8]. The coating of ceramics with a layer of lithium zirconate changes the conductivity from oxygen ion to cationic transport (by Li⁺-ions) [9,10,11,12,13,14]. The instability of electrolytes based on ZrO₂ in the corrosive environment of Li⁺-containing melts [15,16,17] has stimulated the search for alternative materials—that is, sensors of the activity of O²⁻ ions in melts.

One of the promising materials with high oxygen ionic conductivity is the Gd₂Zr₂O₇ solid electrolyte with a pyrochlore structure. It is known that the structure of gadolinium zirconate withstands high doses of radiation due to the ability of the structure to endure cationic disorder. Under the action of radiation, the structure can undergo a structural phase transition from pyrochlore to a structure of defective fluorite, without being subjected to amorphization. In this regard, Gd₂Zr₂O₇ is considered as a matrix for radioactive waste disposal [18,19,20,21,22]. Gadolinium zirconate is also used as a material for thermal barrier coatings [20,21,22], catalysts [23,24], solid electrolytes for SOFCs [25,26], and sensors [27].

The sensory activity of gadolinium zirconate has been insufficiently studied. However, there is evidence showing that when testing Ca⁺²-doped Gd₂Zr₂O₇ as a NO₂ sensor at 500 °C, the device showed the highest current value and the fastest response time, as well as the fastest recovery time. The results obtained were better than those of the commercial YSZ-based sensor [28].

In order to explore the possibility of using the material in the chloride melt as a sensor, it is necessary to study its chemical stability in the melt. For this purpose, it is necessary to determine the range of dopants that can ensure the acquisition of a high-density ceramic material with high, dominant oxygen ion conductivity. Previously, it was found that the use of lithium as a dopant for the Gd₂Zr₂O₇ phase makes it possible to increase the oxygen ion conductivity but does not make it possible to obtain high-density ceramics [29,30]. Therefore, the problem of searching for dopants for the Gd₂Zr₂O₇ phase remains relevant.

It has been established [31] that the homogeneous magnesium doping of Gd₂Zr₂O₇ (solid-solution Gd_2-xMg_xZr₂O_7-x/2) can prevent the ion exchange of high-density ceramics with LiCl-xLi₂O melts up to a Li₂O concentration of 4 wt.%. Thus, a small concentration of magnesium is assumed to be able to protect the material from interaction with the melt [32]. For this article, materials based on Gd₂Zr₂O₇ were chosen as the objects of study. In order to obtain high-density Gd₂Zr₂O₇ ceramics, a composite with MgO additive was used. Composites based on gadolinium zirconate (1-x)Gd₂Zr₂O₇·xMgO were investigated for their chemical resistance in oxide-chloride lithium melts for the first time. The structure and transport properties of (1-x)Gd₂Zr₂O₇·xMgO were also studied for the first time.

2. Materials and Methods

Composites of (1-x)Gd₂Zr₂O₇·xMgO were prepared by mixing gadolinium zirconate Gd₂Zr₂O₇ with magnesium hydroxide. Gadolinium zirconate was synthesized by mixing with the starting reagents Gd₂O₃ (99.998%, VEKTON, RF) and ZrO₂ (99.99%, REACHIM, RF). The oxide powders were ground together in a planetary ball mill (Fritsch Planetar Micro Mill Pulverisette 7, Fritsch GmbH, Federal Republic of Germany) with a rotation speed of 350 rpm in ethanol-exploiting zirconium balls for 8 h. After milling, the mixture was dried and annealed at 1200–1400 °C with steps of 100 °C. Freshly precipitated Mg(OH)₂ was obtained from MgCO₃ (99.99%, VEKTON, RF). The MgCO₃ was dissolved in dilute nitric acid a with pH ~2–3, and then the resulting Mg(NO₃)₂ was caused to interact with a solution of NH₄OH with the formation of a corresponding precipitate. The obtained Mg(OH)₂ was dried and mixed with Gd₂Zr₂O₇ in the amount required for the preparation of the samples of (1-x)Gd₂Zr₂O₇·xMgO (x = 0.05, 0.07, 0.10). The corresponding calculations were carried out taking into account TG studies of the obtained Mg(OH)₂. The resulting mixture was ground in the planetary ball mill at a rotation speed of 700 rpm for 2 h in ethanol. Then, the dried mixtures were pressed into cylindrical briquettes and treated at a temperature of 1500 °C for 12 h. The sintering of the cylindrical briquettes of undoped Gd₂Zr₂O₇ was carried out at a temperature of 1550 °C for 15 h.

A crystal structure analysis of the samples was conducted using an X-ray diffractometer with Cu-Kα radiation (scanning rate of 0.5 °/min, steps of 0.01°) (Rigaku MiniFlex600, Japan) and the SmartLab Studio II (Rigaku, Japan) software with the PDF-2 2019 database (ICDD, USA). Silicon powder (NIST640f) was used as the external standard. The FullProf program was used for the refinement of the parameters with the Rietveld method.

Scanning electron microscopy of the sintered samples was conducted with Phenom ProX (Phenom-World, Holland). Before analysis, the surfaces of the samples were polishing on a diamond wheel. The composition of the samples was determined using energy-dispersive X-ray spectroscopy with a silicon drift detector and Phenom Pro Sute software.

The average density of the ceramics was determined using the geometric method. The value of the crystallographic density of the samples was taken as the standard.

The measurements of ac electrical conductivity, as a function of temperature and oxygen partial pressure, were made using the impedance spectroscopy method (Elins Z-1000P spectrometer, Elins Ltd., Moscow, Russia) in air over a temperature range of 300–1000 °C with cooling of 1 °/min and a pO₂ range of 1 × 10^–18–0.21 atm. The surfaces of the cylindrical-shaped samples were covered with platinum paste. The electrodes were burned in air at a temperature of 900 °C for 3 h. The ZView-4.0 software (Scribner Associate Inc., Southern Pines, NC, USA) was used to treat the impedance data. The oxygen partial pressure measurements were performed using an 8 mol.% Y₂O₃-stabilized zirconia (YSZ) sensor and an oxygen pump.

The stability of the ceramic composite samples was assessed in a glove box (SPEKS GB, Russia) in an argon atmosphere with an impurity oxygen content less than 2 ppm and a water content less than 1 ppm. The salts LiCl, LiCl-2 wt.% Li₂O, and LiCl-4 wt.% Li₂O were placed in a glassy carbon crucible. The salts LiCl (CAS No: 7447-41-8, Leverton Clarke Ltd., Basingstoke, UK) and Li₂O (99.5 wt.%, Alfa Aesar, Haverhill, MA, USA) were used for the experiment. Before the experiment, the LiCl salt was subjected to additional purification by zone recrystallization. The salt in the crucible was heated in a shaft furnace at a temperature of 650 °C. The temperature in the working area was controlled using a Pt-Pt(Rh) thermocouple. The ceramics were preliminarily weighed on a balance and dried in an oven at 200 °C. They were then immersed in a melt of the corresponding salt. Exposure in the salt melts was carried out for 48 h. After holding in the melt, the samples were removed and washed in distilled water and ethanol. After drying the samples in an oven at 200 °C, the composites were analyzed by XRD and SEM.

The components of the melt were analyzed before the experiment and after 8, 16, 24, and 48 h by chemical analysis with inductively coupled plasma atomic emission spectroscopy (Optima 4300 DV Perkin Elmer spectrometer, USA).

3. Results

3.1. X-ray Analysis and SEM Characterization

The interpretation of the X-ray spectra (Figure 1) showed that the main phase in all the obtained composites was Gd₂Zr₂O₇ with a pyrochlore structure (space group Fd-3m). The 0.90Gd₂Zr₂O₇–0.10MgO composite contained a small amount of the second phase, which corresponded to the magnesium oxide. According to the XRD data, the samples with additives of MgO at the amounts of 5 wt.% and 7 wt.% were single-phase.

The lattice parameter decreased with the addition of MgO up to 7 wt.% (Figure 2a). However, the parameter remained unchanged over the concentration range of 7–10 wt.%.

The change in the lattice parameter indicates the formation of solid solution due to the dissolution of magnesium in the structure and the redistribution of cations between the A and B sublattices. The redistribution of cations with different ionic radii will affect the lattice parameter [33]. Figure 2b shows the X-ray pattern of the 0.95Gd₂Zr₂O₇–0.05MgO sample processed using the Rietveld method in the FullProf program. In Figure 2b, one can see the diffraction data, which were treated using a gadolinium and zirconium redistributed model (

Table 1. The refined coordinates of atoms of 0.95Gd₂Zr₂O₇–0.05MgO in a cubic unit cell (space group Fd-3m) in the Gd and Zr disorder model.

Atom	Wyckhoff Symbol	x/a	y/b	z/c	Occupancy
Gd	16c	0.000	0.000	0.000	1.837 (5)
Zr	16c	0.000	0.000	0.000	0.163 (5)
Zr	16d	0.500	0.500	0.500	1.837 (5)
Gd	16d	0.500	0.500	0.500	0.163 (5)
O1	48f	0.411	0.125	0.125	6.00 (5)
O2	8a	0.125	0.125	0.125	1.00 (0)

R_p = 2.56, R_wp = 3.72, R_exp = 2.39, R_b = 3.67, R_f = 4.59, Chi² = 2.41.

). However, the inversion degree between the A and B sublattices grew from 8.15% (0.95Gd₂Zr₂O₇–0.05MgO) to 12.7% (0.90Gd₂Zr₂O₇–0.10MgO) with the increase in the Mg content.

The size of the Mg²⁺ ion is smaller than the Gd³⁺ ion in eightfold coordination [33]; thus, the parameter may decrease. In this regard, we assume that part of the Mg²⁺ dopes the Gd₂Zr₂O₇ matrix and occupies the Gd³⁺ position (

Table 2. The refined coordinates of atoms of 0.95Gd₂Zr₂O₇–0.05MgO in a cubic unit cell (space group Fd-3m), with Gd and Mg existing in the A sublattice.

Atom	Wyckhoff Symbol	x/a	y/b	z/c	Occupancy
Gd	16c	0.000	0.000	0.000	1.880 (5)
Zr	16c	0.000	0.000	0.000	0.070 (5)
Mg	16c	0.000	0.000	0.000	0.050
Zr	16d	0.500	0.500	0.500	1.930 (5)
Gd	16d	0.500	0.500	0.500	0.070 (5)
O1	48f	0.412	0.125	0.125	5.56 (2)
O2	8a	0.125	0.125	0.125	1.00 (0)
O3	8b	0.375	0.375	0.375	0.44 (2)

R_p = 2.46, R_wp = 3.54, R_b = 3.61, R_f = 4.19, Chi² = 3.23.

). Noting the misunderstanding of the exact amount of magnesium oxide included in this calculation, we considered the limited case, where the minimum concentration of the composite additive was included in the Gd₂Zr₂O₇ structure. Accordingly, we can write the composition 0.95Gd₂Zr₂O₇–0.05MgO as a solid solution: Gd_1.88Mg_0.05Zr_0.07Zr_1.93Gd_0.07O₇. The introduction of magnesium is accompanied by an increase in the degree of order of 0.95Gd₂Zr₂O₇–0.05MgO by 11.4% compared to the undoped sample. In our previous work [31], it was shown that the introduction of magnesium into the structure could be achieved at an amount of 0.05 wt.%, and the calculations were conducted using the Rietveld method.

The calculations performed for the two different models of 0.95Gd₂Zr₂O₇–0.05MgO sample presented in Table 1 and Table 2 show close results regarding the obtained calculation errors. It is probable that the lattice parameter is affected by a complex process associated with the disordering of the cationic and anionic sublattices, as well as the presence of small amounts of magnesium. Thus, to describe the structure of the obtained composites, one must not be limited to the X-ray method alone.

The SEM images confirm that all the (1-x)Gd₂Zr₂O₇·xMgO samples contain impurities of MgO. Figure 3a–c shows densely sintered grains in the main phase, which tend to coalesce with each other with an increase in the concentration of MgO additive. Additionally, in all the SEM figures, there are dark spots of poorly crystallized (amorphous) phase. The EDS spectra show that such poorly crystallized spots belong to phase with an increased magnesium content (Figure 4a,b), and we believe that this is magnesium oxide. It is probable that due to amorphization, the low concentration of the composite additive MgO (x = 0.05 and x = 0.07) does not enable its identification in the composites by X-ray diffraction.

The relative densities of the obtained samples are all close to each other (

Table 3. The relative density of Gd₂Zr₂O₇ and (1-x)Gd₂Zr₂O₇·xMgO samples.

Chemical Formula	ρ, %
Gd2Zr2O7	96d24
0.95Gd2Zr2O7–0.05MgO	96.40
0.93Gd2Zr2O7–0.07MgO	96.83
0.90Gd2Zr2O7–0.10MgO	97.27

), and there is a slight tendency for the relative density to increase with the content of magnesium oxide. It should be emphasized that higher temperatures and holding times are required to obtain high-density ceramics of undoped Gd₂Zr₂O₇. Therefore, it can be concluded that the MgO additive contributes to an increase in the sinterability of the material, as shown in previous works [34,35].

3.2. Electrical Conductivities

Well-resolved impedance hodographs, obtained in the range from 300 °C to 1000 °C in air, are presented as two semicircles (Figure 5).

It was possible to separate the volume and grain boundary contributions to the total electrical conductivity. One semicircle, with a capacity of C = 10⁻¹¹ F·cm⁻¹, exceeds zero and is responsible for the bulk properties of the sample, and the second one, with a capacity of C = 10⁻⁹ F·cm⁻¹, is responsible for the properties of the grain boundaries. By extrapolating the first semicircle to the OX axis, the bulk electrical conductivity was determined. The value of the bulk electrical conductivity of (1-x)Gd₂Zr₂O₇-xMgO, shown in Figure 6, is slightly increased compared to that of the undoped Gd₂Zr₂O₇ with a pyrochlore structure over the entire temperature range. The electrical conductivities of all the doped samples are close to each other. This effect is probably due to acceptor doping and the formation of oxygen vacancies.

The results obtained here were compared with the literature data on the conductivity of the most highly conductive phases [36,37]. Figure 7 shows the bulk conductivities of the materials with a pyrochlore structure. The electrical conductivity of the highly conductive phase Sm_2-xCa_xZr₂O_7-x/2 (x = 0, 0.05, 0.10) [36] is comparable with the data that we obtained for the composite (1-x)Gd₂Zr₂O₇·xMgO. It can also be concluded (Figure 7a) that these samples have similar activation energies. Calcium doping of the gadolinium sublattice in Gd₂Zr₂O₇ (Gd_2-xCa_xZr₂O_7-x/2 (x = 0.10)) significantly reduced the conductivity of the material. It can be seen in [37] that all the obtained samples have different activation energies of Gd_2-xCa_xZr₂O_7-x/2, and the electrical conductivity of the obtained compositions sharply decreases in the region of medium temperatures (Figure 7b). Thus, the addition of magnesium oxide MgO to gadolinium zirconate Gd₂Zr₂O₇ did not lead to a decrease in conductivity; the values of ionic conductivity are comparable to those of the most highly conductive pyrochlores. At the same time, the addition of magnesium oxide made it possible to sinter the material at a lower temperature and obtain high-density ceramics. We believe that this method of obtaining high-density ceramics can be extended to other promising systems with a pyrochlore structure. Thus, for pyrochlores, magnesium oxide can be used both as an acceptor dopant and as a sintering additive.

The transport properties of the ceramics were studied in atmospheres with different oxygen partial pressures. The differentiation of the total conductivities into their ionic (σ_ion) and electronic (σ_el) contributions was carried out through the analysis of the isotherms logσ–logpO₂. The total electrical conductivity (σ_tot) obtained in air at the oxygen partial pressure (pO₂) = 0.21 atm is the sum of the ionic and electronic conductivities (1). The ionic electrical conductivity σ_ion does not depend on pO₂ (it is a plateau on the logσ–logpO₂ graph).

σ_tot = σ_el + σ_ion

(1)

Small additions of a dopant do not change the general nature of the dependence. The isotherms of conductivity are shown in Figure 8a–c, where the plateau of ionic conductivity can be observed over the entire temperature range. With an increase in the concentration of the heterogeneous additive to x = 0.10, the plots of electrical conductivity against pO₂ show a positive slope in the region of high oxygen partial pressures, which indicates the appearance of a small p-type contribution to the total conductivity of the sample.

Based on the data obtained, the temperature dependence of ionic conductivity can be plotted. As can be seen in Figure 9, the ionic conductivity of all the composites is higher than that for the undoped phase. It can be assumed that this is due to an increase in oxygen vacancies as a result of acceptor doping and a complex redistribution of cations over the A and B sublattices.

The ionic transport number can be calculated as:

$$t_{ion}=\frac{{\sigma }_{ion}}{{\sigma }_{tot}}$$

(2)

where σ_ion is the ionic conductivity from the plateau of the logσ−pO₂ dependencies between 10⁻¹⁸ atm and 10⁻⁴ atm, and σ_tot is the total conductivity at pO₂ = 0.21 atm. As we can see from the equation, the ionic conductivity and the total conductivity increase, respectively, while the share (t_ion) does not change.

The ion transport numbers of the (1-x)Gd₂Zr₂O₇·xMgO and undoped Gd₂Zr₂O₇ samples are close (Figure 10). The transport numbers tend to increase with decreasing temperature and are close to unity over the entire temperature range. Thus, the addition of magnesium oxide did not affect the dominant ionic nature of transport of the Gd₂Zr₂O₇ phase.

3.3. Chemical Resistance of LiCl and LiCl-Li₂O Melts

Table 4. Photographs of the obtained ceramic samples of (1-x)Gd₂Zr₂O₇·xMgO before and after exposure in LiCl-xLi₂O melts at T = 650 °C.

x in (1-x)Gd2Zr2O7·xMgO	Photo before the Experiment	Photo after 48 h of Exposure in LiCl	Photo after 48 h of Exposure in LiCl-2 wt.% Li2O	Photo after 48 h of Exposure in LiCl-4 wt.% Li2O
0.05
0.07
0.10

shows photographs of the (1-x)Gd₂Zr₂O₇·xMgO ceramic samples before and after interaction with LiCl-xLi₂O melts, where x = 0 wt%, 2 wt%, and 4 wt% at a temperature of 650 °C. As can be seen in the photo, with an increase in the concentration of lithium oxide in the melt, the degree of destruction of the samples increases. In particular, the 0.90Gd₂Zr₂O₇-0.10MgO sample in a LiCl melt did not undergo destruction. The same sample in the LiCl-2 wt.% Li₂O melt was partially destroyed and split into three parts, and in the LiCl-4 wt.% Li₂O melt, it was completely destroyed and divided into three large parts and many small parts, which also underwent dissolution, judging by the presence of a white coating on the surface of the electrolyte during the experiment and chemical analysis. High-density ceramic samples of Gd₂Zr₂O₇ were examined in our earlier work [27]; however, it was shown that Gd₂Zr₂O₇ interacted with the LiCl-2 wt.% Li₂O melt during high-temperature exposure.

Table 5. Change in the mass loss of ceramic samples of (1-x)Gd₂Zr₂O₇·xMgO before and after exposure in LiCl-xLi₂O melts at T = 650 °C.

Composition	Before	After LiCl	After LiCl-2 wt.% Li2O	After LiCl-4 wt.% Li2O
	%	%	%	%
0.95Gd2Zr2O7-0.05MgO	-	+0.28	−0.07	−3.4
0.93Gd2Zr2O7-0.07MgO	-	−0.01	+0.09	−2.89
0.90Gd2Zr2O7-0.10MgO	-	−0.32	−2.73	−11.46

shows the change in mass of the ceramics after the exposure of the ceramic samples of (1-x)Gd₂Zr₂O₇·xMgO in LiCl-xLi₂O melts at a temperature of 650 °C. After holding the samples in the LiCl melt, there was no significant change in the density of the ceramics and no mass loss. Despite the fact that the 0.93Gd₂Zr₂O₇-0.07MgO ceramics cracked, the mass loss during the study was 0.01%. It is likely that the presence of small defects in the form of cracks in 0.93Gd₂Zr₂O₇-0.07MgO significantly affected the strength characteristics of the material and led to destruction; thus, this fact should be taken into account in further studies. When 2 wt.% Li₂O was added to the melt, the 0.90Gd₂Zr₂O₇-0.10MgO ceramics were destroyed, as a result of which the mass loss was ~3%. In the melt with 4 wt.% Li₂O, there was a significant mass loss in all the ceramics under study, decreasing from ~3% to ~11.5%. The greatest loss occurred in the 0.90Gd₂Zr₂O₇-0.10MgO sample.

The X-ray phase analysis data presented in Figure 11a show that the 0.95Gd₂Zr₂O₇-0.05MgO sample interacted with the LiCl melt, since during the experiment, a phase of gadolinium oxide formed on the surface of the material. The samples (0.93Gd₂Zr₂O₇-0.07MgO, 0.90Gd₂Zr₂O₇-0.10MgO) did not chemically interact with the working medium during the 48 h of exposure. Figure 11b shows that in a more aggressive melt containing lithium oxide at an amount of 2 wt.%, exposure led to the interaction of all the ceramics with the melt, together with the formation of a gadolinium oxide phase.

The concentration of lithium oxide also significantly affects the rate of destruction of the material and the formation of end products through the chemical interaction of the material with the melt. Figure 12 shows the X-ray diffraction patterns of the composite surface after 48 h of interaction with the LiCl-4 wt.% Li₂O melt. The surface consists of a mixture of phases formed by the chemical interaction of gadolinium zirconate with the melt components. Such a process can be described by Equations (3)–(6):

$${\mathrm{Gd}}_2{\mathrm{Zr}}_2{\mathrm{O}}_7+2{\mathrm{Li}}_2\mathrm{O}=2{\mathrm{Li}}_2{\mathrm{ZrO}}_3+{\mathrm{Gd}}_2{\mathrm{O}}_3$$

(3)

$${\mathrm{Gd}}_2{\mathrm{O}}_3+{\mathrm{Li}}_2\mathrm{O}=2{\mathrm{LiGdO}}_2$$

(4)

$${\mathrm{LiGdO}}_2+4\mathrm{LiCl}=2{\mathrm{LiGdCl}}_4+2{\mathrm{Li}}_2\mathrm{O}$$

(5)

or

$${\mathrm{Gd}}_2{\mathrm{O}}_3+8\mathrm{LiCl}=2{\mathrm{LiGdCl}}_4+3{\mathrm{Li}}_2\mathrm{O}$$

(6)

The SEM data of the composites are confirmed by the XRD data. Figure 13a–c shows the surfaces of composites exposed in a LiCl melt at a temperature of 650 °C. It was found that the surface of 0.95Gd₂Zr₂O₇-0.05MgO had grains of different contrasts. The surfaces of the remaining compositions remained identical to the initial state. The surfaces of the composites exposed in LiCl-2 wt.% Li₂O were distinguished by grains with different morphologies and contrasts which were unevenly distributed over the ceramic surfaces (Figure 13d–f). The smallest grain growth and grain concentration were recorded for the composite 0.93Gd₂Zr₂O₇-0.07MgO. EDS spectra confirmed (Figure 14) that the dark surface grains corresponded to a phase with a high content of zirconium, and the light grains corresponded to a phase with a high content of gadolinium. With an increase in the content of lithium oxide to 4 wt.%, the morphology of the grains also changed (Figure 13g–i), and their enlargement occurred.

The chemical analysis of the (1-x)Gd₂Zr₂O₇·xMgO composite did not establish the presence of Gd or Zr ceramic components in the LiCl melt. With an increase in the concentration of lithium oxide in the melt, the concentrations of zirconium and gadolinium as components of the melt increased (Figure 15). Additionally, depending on the composition of the ceramics, the saturation time of the melt, with respect to gadolinium, varied. At x = 0.05, the saturation time for the LiCl-4 wt.% Li₂O melt was 8 h, and at x = 0.07 and x = 0.10, the saturation times were 24 and 30 h, respectively.

Thus, the results obtained showed that the addition of magnesium oxide did not lead to an increase in the chemical stability of the phase Gd₂Zr₂O₇ in the melts of LiCl-Li₂O. We believe that this is due to the morphology of the obtained composites when the magnesium oxide phase was localized in the form of individual grains. It is probably that the situation would be more favorable with the addition of magnesium oxide, encapsulating the grains of the main phase, Gd₂Zr₂O₇, and preventing interaction with the melt. Since the system under study was demonstrated to be capable of obtaining high oxygen ion conductivities and high-density ceramics, it is of interest for further research, for example, using higher concentrations of magnesium oxide or changing the method for obtaining a composite. This will be the topic of our further research.

It should be emphasized that there are few studies on the chemical resistance of ceramics in halide melts. The investigations are mainly carried out for simple oxides (doped with an acceptor impurity), for example: ZrO₂(Y₂O₃), ZrO₂(MgO), ZrO₂(Sc₂O₃), ZrO₂(CaO), CeO₂(Gd₂O₃) [38,39]. It was shown that the increase in the content of Li₂O led to the destruction of the samples. Therefore, investigations of Gd₂Zr₂O₇-based pyrochlores are the most promising.

4. Conclusions

Composites with the general formula (1-x)Gd₂Zr₂O₇·xMgO were obtained. Small concentrations of magnesium oxide were dissolved in the complex oxide matrix of Gd₂Zr₂O₇ with the formation of a solid solution, which manifested itself through a decrease in the lattice parameters of the matrix phase. This led to the formation of oxygen vacancies and a complex redistribution of the Gd and Zr over the A and B sublattices. The ionic conductivity of the studied samples of (1-x)Gd₂Zr₂O₇·xMgO (x = 0.05, 0.07, 0.10) was slightly higher than that of the undoped phase. The oxygen ionic conductivity reached 6.6 × 10⁻³ Ohm·cm at 800 °C and 8.1 × 10⁻⁴ Ohm·cm at 600 °C for (1-x)Gd₂Zr₂O₇·xMgO (x = 0.07). The transfer numbers, below 500 °C, reached 100%. The share of dominant ion transport did not change upon doping with magnesium oxide.

Composites based on gadolinium zirconate were investigated for their chemical resistance in oxide-chloride lithium melts for the first time. Composites of (1-x)Gd₂Zr₂O₇·xMgO (x = 0.07, 0.10) showed chemical resistance in lithium halide (LiCl) melts for 48 h at 650 °C. However, when Li₂O was introduced into the melt at the amounts of 2 wt.% and 4 wt.%, the composites chemically interacted with the melt component (Li₂O), with the formation of a Gd₂O₃ phase or a mixture of phases (Gd₂O₃, Li₂ZrO₃, ZrO₂, LiGdO₂, or LiGdCl₂) on the ceramic surface, respectively. This interaction was also confirmed by the chemical analysis of the melt after the experiments.

Additives of magnesium oxide had a significant effect on sintering, which ensured the production of high-density samples at lower temperatures and shorter heat treatment times.