Synthesis and in situ ion irradiation of A-site deficient zirconate perovskite ceramics

Abstract

The influence of A-site vacancies on zirconate perovskite structured ceramics of formula A_1−xLa_2x/3ZrO₃ (A = Ca, Sr, Ba) has been investigated using 400 keV Kr⁺in situ ion irradiation in a TEM with varied temperature from 100 K to 673 K. The solubility limit of La within the Ba_1−xLa_2x/3ZrO₃ system was 5.33 ml% (x = 0.08) using a solid-state synthesis route, with space group symmetry of Pm3̄m for x ≤ 0.08. The Ca_0.9La_0.067ZrO₃ and Sr_0.9La_0.067ZrO₃ samples were found to synthesise single-phase with the Pcmn and Pbnm symmetries of their respective end member structures. The addition of 6.67 ml% La (x = 0.1) in CaZrO₃ and SrZrO₃ was found to decrease the critical temperature of amorphisation, T_c, in comparison with the undoped structures. A similar decrease was found for the 5.33 ml% La doped BaZrO₃ structure. Namely, values were observed for doped samples of ≈404 K, ≈520 K and ≈621 K (A = Sr, Ca, Ba respectively), compared to ≈440 K, ≈623 K and ≈675 K for SrZrO₃, CaZrO₃ and BaZrO₃, respectively. This follows similar behaviour to that observed in titanate systems at low dopant levels, and these findings have been discussed in relation to both the presence of cation vacancies in the doped samples and the presence of the La cation. Void formation, growth and induced morphology changes have also been characterised, onset by ion irradiation.

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1. Introduction

Perovskite structured materials have been widely investigated for various material applications, including as fuel cells,^1–4 dielectric resonators,^5–7 photovoltaics,^8,9 light emitting diodes,¹⁰ lasers,^11,12 next generation nuclear fuels^13,14 and as high-temperature superconducting magnets.^15–18 Another potential application is in the immobilisation of nuclear waste streams as a host ceramic matrix. The namesake perovskite CaTiO₃ has been widely studied as a direct immobilisation matrix, including as a constituent of the SYNROC study,^19–24 while it is a commonly observed secondary actinide bearing phase when targeting zirconolite structured systems.^25–30

In the context of nuclear waste disposal, waste immobilisation matrices will be required to operate with a service lifetime on geological timescales (≈10⁶ years) due to the long half-lives of the isotopes in question. Over such timescales, doping a ceramic structure like perovskite with actinides, i.e. U, Pu, Am, will impart significant radiation damage to the crystalline matrix from alpha-decay processes. Key to ceramic performance in this context will be the damage required to amorphise the crystalline structure, a consequence of which could be increased leaching rates in contact with ground water should a disposal canister fail. In this context, previous studies of radiation damage in perovskites have focused on the titanate series,^31–47 while niobates and tantalates have also received attention.⁴⁸ Zirconate systems, namely CaZrO₃, SrZrO₃ and BaZrO₃, have received little attention, although the response of CaZrO₃ to swift heavy ions has been previously investigated.⁴⁹ The focus of the present study, therefore, is investigation of radiation damage effects within these zirconate systems using in situ ion irradiation techniques. Such techniques are frequently employed to calculate amorphisation doses over accelerated timescales.

Perovskite structured materials take the base structural formula ABO₃ and an array of elemental combinations have been synthesised and characterised. The aristotype perovskite structure is exhibited by SrTiO₃ in a cubic Pm3̄m arrangement, with Sr cations located at the unit cell corners, a body centred Ti cation and face-centred corner sharing O atoms, forming corner sharing TiO₆ octahedra throughout the structure. While SrTiO₃ exhibits the ideal ratio of A/B site cation radii (as governed by the Goldschmidt tolerance factor⁵⁰), distorted structures can form when this ratio changes through the tilting of BO₆ octahedra and, for example, anti-parallel A-site cation displacement. Within this study, the CaZrO₃, SrZrO₃ and BaZrO₃ compositions have been investigated, and these structures are shown graphically in Fig. 1. CaZrO₃ presents at room temperature in the orthorhombic Pcmn (a⁻a⁻c⁺ using Glazer notation⁵¹), with CaZrO₃ reported to remain orthorhombic <2023 K prior to a transition to cubic symmetry.^52,53 SrZrO₃ takes the Pbnm (a⁻a⁻c⁺) structure, with temperature induced phase transitions investigated by various groups.^54–56 Most recently Hasegawa et al.⁵⁷ revisited the SrZrO₃ system and reported transitions to Ibmm (a⁻a⁻c⁰) at 1042 K, I4/mcm (a⁰a⁰c⁻) at 1119 K and to Pm3̄m (a⁰a⁰a⁰) at 1375 K. BaZrO₃ has the aristotype perovskite structure of Pm3̄m (a⁰a⁰a⁰), with no temperature related phase-transitions reported within the literature from 10 K to 1273 K.^58–60 Both space groups Pcmn and Pbnm are non-stand settings of Pnma (space group 62), however, within this work we refer to the most commonly used settings within the literature.

Fig. 1 Structures of (a) CaZrO₃ (Pcmn; a⁻a⁻c⁺), (b) SrZrO₃ (Pbnm; a⁻a⁻c⁺) and (c) BaZrO₃ (Pm3̄m; a⁰a⁰a⁰), as viewed slightly tilted off the [100] cubic direction. Ca ions are presented as red spheres, Sr ions as blue spheres, Ba ions as white spheres and ZrO₆ octahedra in pink. Figures were produced using the CrystalMaker software package using structural information from Levin et al.⁷ Kennedy et al.⁵⁴ and Parida et al.⁵⁹

A further consequence of doping the perovskite ABO₃ structure is the likely induction of cation vacancy defects. Perovskite structures have been reported to exhibit a remarkable tolerance for such cation deficiency, and this has been hypothesised in a disposal context through doping experiments with La and Gd.⁶¹ When considering the multivalent nature of those isotopes suitable for ceramic immobilisation, understanding the effects of cation deficiency on radiation effects is of key interest. Titanate compositions investigated in this regard include the A-site deficient Sr_1−xLa_2x/3TiO₃ (ref. 62) and Ca_1−xLa_2x/3TiO₃ (ref. 63 and 64) systems. A-site cation vacancies are produced within these systems whereby a cation vacancy is induced for every two dopant La cations. In both of these systems, in situ ion irradiation in a transmission electron microscope (TEM) has shown a reduction in the critical temperature of amorphisation, T_c, in the 0.1 ≤ x ≤ 0.4 region. This value is defined as the temperature at which the damage rate equals the rate of recovery from damage, for a specific irradiation condition. A lower T_c value can therefore be interpreted as a sample having recovery characteristics that require less thermal energy to suppress amorphisation under irradiation. While this could be linked to a possible higher displacement energy for La, for ≥ 0.5 T_c was found to increase with respect to each end member, producing the opposite trend at higher La contents. It has been hypothesised that the increased A-site vacancy content in these systems increases recovery rates through increased defect migration and the increased likelihood for recombination events on the perovskite A-site. Therefore, as well as studying the end-member zirconate systems, this paper presents work to investigate whether such effects translate to zirconate perovskite systems in the context of nuclear waste disposal. To this purpose, samples of Ca_0.9La_0.067ZrO₃, Sr_0.9La_0.067ZrO₃ and Ba_0.92La_0.053ZrO₃ were also synthesised and investigated using in situ ion irradiation. Synthesis and irradiation of a Ba_0.9La_0.067ZrO₃ sample was intended, however, a study to find the solubility limit of La within BaZrO₃ determined this sample would not synthesise, as is presented within this paper.

2. Experimental

2.1 Sample synthesis

Compositions of A_1−xLa_2x/3ZrO₃ (A = Ca, Sr; x = 0 and 0.1) and Ba_1−xLa_2x/3ZrO₃ (x = 0, 0.02, 0.04, 0.06, 0.08 and 0.1) were synthesised via a solid-state route. The additional Ba_1−xLa_2x/3ZrO₃ samples were synthesised in order to find the solubility limit of La in this system. Stoichiometric quantities of CaCO₃, SrCO₃, BaCO₃, ZrO₂ (all 99.9%; Sigma-Aldrich), each dried at 180 °C, and La₂O₃ (99.9%; Sigma-Aldrich), dried at 800 °C, were weighed to stoichiometric amounts and milled in isopropanol using a planetary ball mill with silicon nitride milling media for 15 minutes at 400 rpm. After milling to homogeneity, powders were dried at 80 °C, sieved (212 μm; steel) and uniaxially pressed into 10 mm pellets to produce green bodies. Samples were then fired in zirconia crucibles at 1450 °C for 48 hours in air, re-ground using an agate pestle and mortar and re-sieved to produce fine, reacted powders.

2.2 X-ray diffraction

X-ray diffraction (XRD) was undertaken using a Panalytical X'pert³ powder diffractometer, operating at 45 mV, 40 mA with Cu Kα (λ = 1.5418 Å). The scan range was 10° to 90° with a step size of 0.02° and step time of 0.2 s. A 1/4° divergent slit was employed with 0.02 rad Soller slits. Powders were backloaded into sample holders with the rear surface of packed powders exposed to the X-ray beam, minimising preferential orientation. All data were calibrated using an external NIST 640e line position standard (ESI Fig. 1†). Lattice parameter refinement was undertaken using the Bruker TOPAS software package, with all errors reported as calculated by the refinement process. Such errors were likely underestimated due to the use of a laboratory-based X-ray source and this should be noted when referring to the lattice parameters.

2.3 Transmission Electron microscopy

TEM analysis was undertaken using a JEOL F200, operating at 200 keV, with images and electron diffraction patterns captured using a Gatan OneView. Electron energy loss spectroscopy (EELS) was undertaken using a Gatan Image Filter (GIF) Quantum SE. Samples were prepared for TEM via the powdered “crushed grain” route, with powder ground in isopropanol and pipetted onto a carbon holey lined Cu TEM grid.

2.4 In situ ion irradiation

In situ ion irradiation was undertaken at the University of Huddersfield's MIAMI (Microscopes and Ion Accelerators for Materials Investigations) facility. The MIAMI-2 system was utilised, which comprises a Danfysik 921A ion source coupled with a 350 kV National Electrostatics Corporation ion accelerator, with the ion beam entering the TEM column of a 300 kV Hitachi H-9500 TEM at an angle of 18.7° to the electron beam.⁶⁵ Samples were prepared via the crushed grain route as described previously, with Au grids used for temperatures of >300 K and Cu for ≤300 K. Prior to ion irradiation, the presence of each intended element in the target composition of the specific grains monitored was confirmed via EELS spectroscopy using a Gatan Image Filter (GIF) Quantum SE.

The ion beam setup employed was 400 keV Kr⁺ ions with a flux of 2 ± 0.2 × 10¹² ions per cm². Kr⁺ was employed for comparative purposes with the studies of Smith et al.⁶³ and Whittle et al.⁶⁴ Nuclear/electron stopping powers and ratios for each composition were estimated using the SRIM Monte Carlo code.⁶⁶ The “Ion Distribution and Quick Calculation of Damage” option was used with an average displacement energy of 50 eV set for each constituent element. Values calculated using SRIM are summarised in Table 1, including the determined theoretical densities of the powders used in the calculations and a ratio of electron to nuclear stopping powers (ENSP). Graphical SRIM curves are presented in ESI Fig. 2.† Under these conditions, stopping mechanisms were dominated by the nuclear component in all cases. Considering a TEM grain thickness of <100 nm, it was a requirement to induce displacement damage across the entire grain to observe full amorphisation. Peak displacement damage depth was estimated by SRIM to be 90 ± 10 nm in all cases, extending to a full damage range of ≈200 nm, ensuring the entire grain was subject to displacement damage. For each sample, the Kr retention to a thickness of <100 nm was estimated to be <5% by integrating the SRIM profiles for implantation depth.

Table 1 SRIM parameters and stopping powers^a

Sample	Density (g cm^−3)	(dE/dx)n (eV nm^−1)	(dE/dx)e (eV nm^−1)	ENSP
a Density (g cm^−3), energy loss to nuclear stopping (dE/dx)n (eV nm^−1), energy loss to electronic stopping (dE/dx)e (eV nm^−1) and electronic to nuclear stopping power ratio ENSP.
CaZrO3	4.62	1702	453.0	0.2662
Ca0.9La0.067ZrO3	4.73	1696	456.9	0.2694
SrZrO3	5.45	1736	490.2	0.2824
Sr0.9La0.067ZrO3	5.47	1719	488.1	0.2839
BaZrO3	6.23	1703	514.5	0.3021
Ba0.9La0.067ZrO3	6.16	1691	510.4	0.3018

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The electron beam valve was closed while ion irradiation was undertaken in order to avoid dual-beam synergistic effects, with the ion beam blocked off at intervals to allow for imaging and electron diffraction analysis. Several grains were monitored at each interval using bright-field (BF) imaging and selective area electron diffraction (SAED). A Gatan 652 heating holder was used for irradiations at elevated temperatures, and a Gatan 636 liquid nitrogen cooling stage was used for temperatures below ambient. An average of the time/fluence interval at which all Bragg diffraction peaks diminished and only amorphous diffuse scattering was observed within SAED patterns and the previous interval was taken as the critical amorphisation fluence, F_c, at a given temperature. This process for a SrZrO₃ grain is shown in Fig. 2. The error in F_c was taken as the difference between the two intervals in which F_c was observed to fall. An average of F_c from all monitored grains gave the final F_c value for a particular sample and temperature. An additional 10% error in the flux was applied to account for fluctuations in the ion beam. The ion beam flux was also measured before and after irradiation to further monitor beam stability.

Fig. 2 SAED patterns of a SrZrO₃ grain taken at intervals during irradiation to determine the critical fluence for amorphisation, F_c, at room temperature. Fluences are (a) 0F_c, (b) 0.35F_c, (c) 0.8F_c, (d) F_c. (a) is indexed to the [211] direction in Pbnm.

By measuring F_c as a function of temperature, the critical temperature of amorphisation, T_c, was determined. T_c is defined as the temperature at which the rate of damage induced equals the rate of recovery from damage under specific irradiation conditions. As previously implemented by various groups,^{37,62,63,67–76} values of T_c were calculated using a nonlinear least squares refinement of eqn (1) using several models that can be accessed elsewhere.^77,78

where F_c0 is the critical amorphisation fluence at 0 K, E_a is the activation energy for recovery from radiation damage, k_b is the Boltzmann constant (8.6173 × 10⁻⁵ eV K⁻¹) and T is temperature. F_c0 was extrapolated from the acquired data using eqn (1). The frequent underestimation of E_a has been reported using eqn (1) when compared to defect migration energies reported in oxides. Therefore, as for other studies, the relationship developed by Weber⁷⁷ was employed to determine both the thermal and irradiation assisted activation energies for recovery from damage, as outlined in eqn (2).where ϕ is the ion flux and ν the effective jump frequency. Values for irradiation-assisted (10¹² s⁻¹) and thermal (10⁹ s⁻¹) jump frequency were used based on previously reported estimations to allow for calculation of E_a values for both thermal (E_th) and irradiation assisted (E_irr).^37,62,63,76

3. Results and discussion

3.1 Sample synthesis

Initially, samples of target compositions CaZrO₃, Ca_0.9La_0.067ZrO₃, SrZrO₃, Sr_0.9La_0.067ZrO₃, BaZrO₃ and Ba_0.9La_0.067ZrO₃ were synthesised. XRD patterns of the Ca_1−xLa_2x/3ZrO₃, Sr_1−xLa_2x/3ZrO₃ and Ba_1−xLa_2x/3ZrO₃ (x = 0 and 0.1) samples are presented in Fig. 3a–c, respectively. Magnifications of the highest intensity perovskite reflections are included; (121) in Pcmn, (112) in Pbnm and (110) in the ideal cubic Pm3̄m. For the Ca_1−xLa_2x/3ZrO₃ and Sr_1−xLa_2x/3ZrO₃ systems, no additional peaks were identified other than those related to their end member counterparts, confirming the parent structure was maintained by the doped samples, i.e. Pcmn for CaZrO₃ and Ca_0.9La_0.067ZrO₃; Pbnm for SrZrO₃ and Sr_0.9La_0.067ZrO₃. La incorporation into the structures was confirmed through lattice parameter contraction/expansion, as can be observed through the shifts of diffraction peaks to lower 2θ values (higher d-spacing) for Ca_0.9La_0.067ZrO₃, and higher 2θ values (lower d-spacing) for Sr_0.9La_0.067ZrO₃. These shifts are readily observable in the magnified regions in Fig. 3a and b, and lattice parameters for all samples are reported in Table 2 with fitting errors. The expansion in lattice parameters for Ca_0.9La_0.067ZrO₃ was due to the smaller ionic radius of La (1.16 Å) when replacing Ca (1.12 Å) in 8-fold coordination, while Sr_0.9La_0.067ZrO₃ lattice parameters contract due to the larger Sr (1.26 Å) that is replaced, again in 8-fold coordination.

Fig. 3 XRD patterns of as labelled (a) CaZrO₃ and Ca_0.9La_0.067ZrO₃, (b) SrZrO₃ and Sr_0.9La_0.067ZrO₃ and (c) BaZrO₃ and Ba_0.9La_0.067ZrO₃. A magnification of the most intense reflection is presented alongside the full patterns. Patterns are normalised to the most intense reflection. Patterns are indexed to structure information from ICSD database codes 97 474 (CaZrO₃), 173 398 (SrZrO3) and 90 049 (BaZrO₃).

While lattice contraction was observed in the Ba_0.9La_0.067ZrO₃ sample as compared to BaZrO₃, a secondary phase of the pyrochlore-structured La₂Zr₂O₇ was detected (see Fig. 3c). This suggested the solubility of La within the Ba_1−xLa_2x/3ZrO₃ system was below x = 0.1 for this synthesis route. This was likely due to the large difference in size between Ba (1.61 Å) and La (1.36 Å) in the A-site 12-fold coordination state in Pm3̄m. The Goldschmidt tolerance factor,⁵⁰t, has been extensively used to predict the chemical compositions that will form as perovskites, however, to our knowledge no tolerance factor exists that can accurately predict the effects of A-site vacancies on the likelihood of perovskite formation. Considering an average cation radius based on either a fully occupied La site (i.e. Ba_0.9La_0.1ZrO₃) or the correct ratio (ignoring the presence of vacancies, as suggested by Ganguly et al.⁷⁹) reduces t below the ideal BaZrO₃ value of ≈1, but is still in the regime for perovskite formation. As Whittle et al.⁸⁰ report, vacancy defects can occupy more space than a fully-occupied cation lattice site, which is contrary to both these calculations. Without a tolerance factor that truly incorporates a specific vacancy-type defect, such calculations can never be considered accurate. Even so, the incorporation of a quantity of La within the target Ba_0.9La_0.067ZrO₃ formula was confirmed by the peak shift observed in Fig. 3c, confirming lattice expansion due to ionic substitution.

Samples of composition x = 0.02, 0.04, 0.06 and 0.08 were produced to identify the La solubility limit in Ba_1−xLa_2x/3ZrO₃. XRD patterns of the Ba_1−xLa_2x/3ZrO₃ system for x = 0, 0.02, 0.04, 0.06, 0.08 and 0.1 are presented in Fig. 4a, with magnifications of the 27° ≤ 2θ ≤ 32° region in Fig. 4b and the (110) reflection in Fig. 4c. A magnification of the 27° ≤ 2θ ≤ 32° region in which the (222) pyrochlore La₂Zr₂O₇ reflection was observed is presented in Fig. 4b. For x ≤ 0.08, no such reflection was observed, confirming the solubility limit of La in the region 0.08 ≤ x < 0.1. For x ≤ 0.08 the Pm3̄m structure was present, confirming the single-phase nature of these samples. A shift to higher 2θ values was observed incrementally with increasing La content (Fig. 4c) until x = 0.1, with calculated lattice parameters and unit cell volumes reported in Table 2. A linear decrease in unit cell volume was observed for the Ba_1−xLa_2x/3ZrO₃ system until the solubility limit was reached at x = 0.1. The contraction and expansion of the unit cells of the Sr_1−xLa_2x/3ZrO₃ and Ca_1−xLa_2x/3ZrO₃ systems was also observed, as is reported in Table 2.

Fig. 4 (a) Full XRD traces of the Ba_1−xLa_2x/3ZrO₃ samples (x = 0, 0.02, 0.04, 0.06, 0.08 and 0.1). Indexing is provided to the Pm3̄m structure with reflections indexed to La₂Zr₂O₇ indicated with a superscript *. Patterns are indexed based on structural information from ICSD database codes 90 049 (BaZrO₃) and 153 222 (La₂Zr₂O₇). The bottom trace is x = 0 as labelled, with each trace above increasing in increments of x = 0.02. (b) Magnification of the 27° ≤ 2θ ≤ 32° region indicating the (222) La₂Zr₂O₇ reflection present in the x = 0.1 sample and (c) magnification of the (110) reflection. The grey dashed line indicates the position of the (110) reflection of the x = 0 sample.

Table 2 Calculated lattice parameters of the Ca_1−xLa_2x/3ZrO₃, Sr_1−xLa_2x/3ZrO₃ and Ba_1−xLa_2x/3ZrO₃ systems^a

Sample	a (Å)	b (Å)	c (Å)	V (Å^3)	R wp (%)	R exp (%)
a Lattice parameters a, b and c (Å), unit cell volume V (Å^3) and refinement errors Rwp and Rexp (%).
Pcmn (a^−a^−c^+)
CaZrO3	5.58781 (8)	8.0123 (1)	5.75817 (8)	257.800 (6)	6.03	2.94
Ca0.9La0.067ZrO3	5.60273 (6)	8.02863 (9)	5.76078 (7)	259.133 (5)	6.04	3.00
Pbnm (a^−a^−c^+)
SrZrO3	5.79669 (8)	5.82136 (8)	8.2092 (1)	277.016 (6)	8.21	5.46
Sr0.9La0.067ZrO3	5.7955 (1)	5.8186 (1)	8.2081 (1)	276.791 (8)	7.71	5.60
Pm3̄m (a^0a^0a^0)
BaZrO3	4.19109 (4)	—	—	73.6177 (12)	9.99	6.32
Ba0.98La0.013ZrO3	4.19012 (4)	—	—	73.5665 (12)	7.54	5.55
Ba0.96La0.0267ZrO3	4.18963 (3)	—	—	73.4546 (1)	7.74	5.67
Ba0.94La0.04ZrO3	4.18603 (2)	—	—	73.3513 (1)	6.31	4.31
Ba0.92La0.053ZrO3	4.18421 (2)	—	—	73.2558 (12)	6.85	4.46

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Bisen et al.⁸¹ report that the solubility of La within the Ba_1−xLa_2x/3ZrO₃ (x = 0, 0.02, 0.04, 0.06, 0.08, 0.1) solid-solution extends to x = 0.1. They also report an extreme lattice contraction (6.04%, equivalent to 0.25 Å), which was not observed in our analysis. It is unlikely that such contraction could be induced by cation substitution, and if the lattice parameter reduction is taken as reported for x = 0 → 0.02 → 0.04 → 0.06 → 0.08 → 0.1, the associated reduction in the cubic parameter would be equivalent to 6.04% → 0.37% → 0.18% → 0.17% → 0.02%. Errors are not reported by the authors, and it is likely that the 0.02% reduction from x = 0.08 → 0.1 may be within error, and a solubility limit of La in the solid-solution had also been reached in this case.

Electron diffraction was further used to confirm the space group symmetry of each doped composition, with SAED patterns taken with the electron beam orientated down the [001] axis presented for the Ca_0.9La_0.067ZrO₃, Sr_0.9La_0.067ZrO₃ and Ba_0.08La_0.053ZrO₃ compositions presented in Fig. 5a, b and c, respectively. In each instance, the patterns from these doped samples were found to index to the counterpart parent end member, namely Pcmn for CaZrO₃, Pbnm for SrZrO₃ and Pm3̄m for BaZrO₃. The presence of La within powdered specimens of the Ca_0.9La_0.067ZrO₃, Sr_0.9La_0.067ZrO₃ and Ba_0.92La_0.053ZrO₃ samples was further confirmed using TEM/EELS. BF micrographs and corresponding EELS spectra are presented within ESI Fig. 3, 4 and 5.†

Fig. 5 Electron diffraction patterns taken down the [001] zone axis for (a) Ca_0.9La_0.067ZrO₃ indexed to Pcmn, (b) Sr_0.9La_0.067ZrO₃ indexed to Pbnm and (c) Ba_0.08La_0.053ZrO₃ indexed to Pm3̄m.

3.2 In situ ion irradiation

Values of critical amorphisation fluences for all samples are presented in Fig. 6a, with fits from the nonlinear least-squares refinements using eqn (1) also presented. Included within Fig. 6a are values for SrZrO₃ at 473 K, and BaZrO₃ at 673 K, at which amorphisation was not achieved, presented as grey data points. These values lie above the calculated T_c curves for the respective samples, suggesting these irradiations took place at temperatures that prevented amorphisation. Values for F_c0, T_c and E_a, calculated using eqn (1), and E_th and E_irr values, calculated using eqn (2), are reported in Table 3. Calculated values for T_c are also shown graphically in Fig. 6b.

Fig. 6 (a) Values of the critical fluence of amorphisation, F_c at given temperatures. Fits based on eqn (1) are also presented. The grey solid square refers to a SrZrO₃ fluence at 473 K and the grey solid circle to a BaZrO₃ fluence at 673 K; amorphisation was not achieved in either of these cases. (b) Values of the critical temperature of amorphisation, T_c for each sample plotted with atomic number. Each T_c value is located above the respective “Ca”, “Sr” and “Ba” markers referring to pristine and La doped CaZrO₃, SrZrO₃ and BaZrO₃ samples. Pristine samples are presented in black, and La doped in red.

Table 3 Fitting parameters for calculation of critical temperatures of amorphisation, T_c^a

Sample	From eqn (1)			From eqn (2)
	F c0 (× 10^14 ions per cm^2)	T c (K)	E a (eV)	E th (eV)	E irr (eV)
a Critical amorphisation fluence at 0 K Fc0 (× 10^14 ions per cm^2), critical temperature of amorphisation, Tc (K), activation energy Ea (eV), activation energy for thermally assisted recovery Eth (eV) and activation energy for irradiation assisted recovery Eirr (eV).
SrZrO3	6.4 (7)	440 (18)	0.05 (1)	1.01 (4)	1.27 (4)
Sr0.9La0.067ZrO3	7.2 (3)	404 (3)	0.042 (4)	0.926 (8)	1.167 (8)
CaZrO3	5.8 (9)	623 (11)	0.060 (8)	1.42 (3)	1.79 (3)
Ca0.9La0.067ZrO3	6.2 (8)	520 (8)	0.041 (5)	1.19 (2)	1.50 (2)
BaZrO3	4.1 (6)	675 (12)	0.053 (9)	1.52 (3)	1.92 (3)
Ba0.92La0.053ZrO3	4.7 (7)	561 (25)	0.042 (9)	1.27 (6)	1.60 (6)

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Firstly, the end member compositions have T_c values with the trend of SrZrO₃ (440 K) < CaZrO₃ (623 K) < BaZrO₃ (675 K). This was similar to the titanate compositions of CaTiO₃, SrTiO₃ and BaTiO₃ reported by Meldrum et al.⁴⁸ which showed a T_c relationship of SrTiO₃ (425 K) < CaTiO₃ (440 K) < BaTiO₃ (550 K). In all instances, the T_c values for the zirconate end members calculated in this study were appreciably higher than their A-site titanate equivalents. While Meldrum et al.⁴⁸ employed a 1 MeV Kr⁺ ion beam as opposed to the 400 keV Kr⁺ beam utilised in the present study, a significant difference in T_c under the two different conditions would not be expected considering the SRIM calculations discussed previously, although the temperature specific F_c values may be altered. Most comparisons within the present study are with those using 800 keV to 1 MeV Kr⁺ ion beams, and while a significant difference in calculated T_c values is not expected, the difference in the ion beam used should be noted. This includes the studies of A-site deficient systems previously reported.^62–64

The relationship between titanates and zirconates differs from those in the pyrochlore system, in which zirconates such as La₂Zr₂O₇ and Gd₂Zr₂O₇ show a remarkable tolerance for radiation damage compared to their titanate counterparts.^82–85 These relationships are governed by anti-site defect formation energies and phase transitions to the defect fluorite structure,^35,86 processes that are not replicated within perovskites. The calculated T_c values suggest the specific A-site cation plays a larger role than structure within perovskite systems. For the titanate systems, T_c values take the structural trend Pm3̄m < Pbnm < I4/mcm, while for the zirconates the trend is Pbnm = Pcmn < Pm3̄m, showing no relationship with relative symmetry, tolerance factor⁵⁰ or A/B site radii ratio. A recent study by Meena et al.⁸⁷ suggests a relative increase in the ionic-nature of bonding for SrTiO₃ > BaTiO₃ > CaTiO₃, which does not follow the T_c trends observed. With regard to displacement energy, E_d, a variety of values through both experimental and theoretical routes have been determined, but no conclusive values can be assigned to these systems. While values have been determined for CaZrO₃, the authors noted the reported cation values were appreciably lower than previously measured oxide samples,⁸⁸ preventing comparison between titanate and zirconate systems. It is thus hypothesised that the flexible nature of the Ti(III)/Ti(IV) B-site, in comparison with the highly-refractory Zr(IV), may promote defect mobility and an increased tolerance for oxygen displacements at higher temperatures. This could take the form of ATi_1−x⁴⁺Ti_x³⁺O_3−δ phases being induced during irradiation, whereas zirconium reduction is far less likely, and the resultant oxygen deficiency is harder to accommodate. This may also be the case for the as produced pristine materials, as has been observed for (Ba, Sr, Ca)TiO₃ perovskites.⁸⁹

For the La-doped samples, similar trends as with doped titanate perovskites are observed. For instance, a decrease in T_c between CaTiO₃ and Ca_0.9La_0.067TiO₃ is reported of ≈140 K,⁶³ while a drop of ≈100 K is reported here for the counterpart zirconates. Similarly, a reduction in T_c of ≈ 90 K is reported between SrTiO₃ and Sr_0.85La_0.1TiO₃ (ref. 62) and a decrease of ≈40 K within the zirconate counterparts is found in the present investigation (note that the titanate contains 3.33 mol% greater La content than within our study). In direct comparison, the titanates exhibit a lower T_c for both Ca and Sr based end members, with Ca_0.9La_0.067TiO₃ (<300 K) > Ca_0.9La_0.067ZrO₃ (520 K) and Sr_0.9La_0.067TiO₃ (308 K) > Sr_0.9La_0.067ZrO₃ (402 K). No structural changes were induced between the parent and doped samples in any of these cases. It is possible that La has an intrinsically higher E_d than each parent A-site constituent. Considering the observations reported regarding the Sr_1−xLa_2x/3TiO₃ and Ca_1−xLa_2x/3TiO₃ systems, in which La doping of x > 0.4 samples increased T_c with regard to the parent end member, it seems unlikely that an increased A-site displacement energy is the sole cause of the trends observed in these zirconates. Values for La within La₂Zr₂O₇ are reported to be similar to Ba, Sr and Ca in perovskite systems, namely E_d = 65 ± 20 eV,^90–92 providing further evidence against such a conclusion. As is reported in Table 3, the activation energies for both thermal and irradiation assisted recovery are reduced for all doped samples with relation to their respective end member counterparts. This does suggest that increased defect migration and ionic diffusion may be the cause of the reduction in T_c for doped samples, however direct measurement of such characteristics would be required to prove this hypothesis.

Several studies on the effects of doping and sintering characteristics on perovskite systems have been reported, and the resultant consequences on cation and oxygen deficiency.^93–96 Within studies of Sr_1−xLa_2x/3TiO₃ and Ca_1−xLa_2x/3TiO₃,^62,63 samples were sintered at 1573 K in air, while the bulk ion irradiation study showing increased room temperature resistance to amorphisation for 0.2 ≤ x ≤ 0.4 used pellets sintered at 1673 K.⁶⁴ Oxygen loss has been shown to rapidly increase with sintering temperature in the A-site deficient Sr_0.85La_0.1TiO₃ system, even at temperatures of 1473 K in flowing O₂.⁹⁷ This has the consequence of causing significant reductions in the activation energy, E_a, for bulk ionic conductivity, which can be readily linked to the concentration and mobility of oxygen vacancies within the system. Akin et al.⁹⁷ report a reduction in E_a from 0.97 eV for SrTiO₃ sintered at 1723 K in O₂ to 0.16 eV for Sr_0.85La_0.1TiO₃ sintered under the same conditions. This decreased energy requirement for oxygen vacancy mobility may lead to increased recovery rates, particularly at elevated temperatures, that lead to reductions in T_c. This does not take into consideration sintering temperatures. E_a values of 1.64 eV and 1.79 eV are reported in Sr_0.85La_0.1TiO₃ sintered at 1473 K under O₂ and N₂, respectively, and values of 0.16 eV at 1723 K in O₂. They report the same sample sintered at 1723 K in N₂ as being too conductive to measure even at 10 K, suggesting an even lower E_a value for this sintering condition. It may be that such effects regarding sintering atmosphere and temperature translate to radiation damage kinetics, but it should be noted the effects of sintering temperature are not as dramatic in SrTiO₃. While these are feasible explanations, such mechanisms have not been tested in the present study and it is not possible to transfer this theory to the zirconates studied here, other than through empirical observations of the similarities in T_c. Based on the T_c observations made, it is proposed that a similar mechanism leads to the reductions in T_c for doped samples, i.e. less oxygen deficiency within a zirconate system as compared to a titanate. Studying bulk pellets of A-site deficient samples sintered under various temperatures and atmospheres may elucidate this behaviour, and studies are now taking place in this regard.

Fig. 7 presents under-focus (≈500 nm) BFTEM micrographs of an irradiated grain of the Ca_0.9La_0.067ZrO₃ sample at 473 K (Fig. 7a–d) and 100 K (Fig. 7e–h). At a fluence of 0.3 of the critical fluence for amorphisation (0.3F_c; Fig. 7b) at 473 K, a high density of light circular contrast regions was observed. With increased fluence to 0.8F_c (Fig. 7c), these circular regions grew in size whilst decreasing in areal density, suggesting a migration and agglomeration of these features. On first observation, it could be argued this was the result of Kr bubble formation, with the elevated temperature promoting Kr migration and agglomeration into bubbles. Forming Kr bubbles in ceramic oxide samples normally requires the implantation of on the order of 10¹⁵ to 10¹⁶ ions per cm², and considering the <5% Kr retention (as predicted by SRIM) under these conditions it is unlikely this contrast was the result of Kr bubbles. Furthermore, previously observed bubble sizes in ceramic oxide systems are on the order of ≈2 nm, even at these elevated temperatures.^98–101 Therefore, it is likely these areas are void-type defects, i.e. a high concentration of irradiation induced vacancies, with the elevated temperatures allowing for void growth and agglomeration. It can be qualitatively observed that at 0.8F_c (Fig. 7c), larger voids were present and in higher density than at the amorphisation fluence, F_c (Fig. 7d). This process was likely driven by void migration and growth toward the grain edges, with the voids annihilating at the surfaces that act as a defect sink. While these observations evidence the microstructural build-up of defects under irradiation, such defects do not appear to have significantly affected the increased recovery rates onset by increased thermal energy, with an amorphisation fluence of F_c = 7.3 ± 0.8 × 10¹⁵ ions per cm² for this sample at 473 K.

Fig. 7 BF micrographs of Ca_0.9La_0.067ZrO₃ irradiated at 473 K to a fluence of (a) 0F_c, (b) 0.3F_c, (c) 0.8F_c and (d) F_c and at 100 K to fluences of (e) 0F_c, (f) 0.3F_c, (g) 0.8F_c and (h) F_c.

The morphology of the grain was also observed to change, with the grain edges smoothing with increasing fluence compared to the irregularly shaped pristine grain. This may have been due to material sputtering, however grain swelling has been reported in various ceramics under ion irradiation, and it may be that a combination of irradiation and temperature induced migration induced a “flow” of material that induced such a change in morphology. In contrast, at 100 K, no void formation was observed at any fluence (Fig. 7e–h), in a similar manner to void formation in steels that requires elevated temperatures. Indeed, it may be that the increased recovery rates at 473 K were required to induce a significant density of vacancy defects to allow for the formation of observable voids, or that at 100 K there is insignificant defect mobility to allow for the required agglomeration. Similarly, changes in grain morphology at this temperature were severely reduced, however this may have resulted from either reduced temperature preventing “flow” or reduced sputtering at a lower fluence. Such effects were not observed at room temperature or below in any samples, confirming this was a thermally driven process.

4. Conclusions

Solid-solutions of single-phase Ca_0.9La_0.067ZrO₃ and Sr_0.9La_0.067ZrO₃ were produced, and the solubility limit of La in BaZrO₃, through the compositional formula Ba_1−xLa_2x/3ZrO₃, was found at x = 0.08. Each of these samples maintained their parent end member structure, namely Pcmn, Pbnm and Pm3̄m for CaZrO₃, SrZrO₃ and BaZrO₃, respectively. Values of the critical temperature of amorphisation, T_c, were calculated of ≈440 K, ≈623 K and ≈675 K for SrZrO₃, CaZrO₃ and BaZrO₃, respectively, compared to ≈404 K, ≈520 K and ≈621 K for the La doped counterparts (Ba_0.08La_0.053ZrO₃ for the Ba sample). The response of the materials to ion irradiation were rationalised by linking to studies regarding ionic conductivity within doped perovskite titanates, and in particular oxygen vacancy mobility within such systems. Void formation, migration and agglomeration at elevated temperatures was observed, including changes in grain morphology, mechanisms that were not observed below room temperature.

Conflicts of interest

There are no conflicts of interest to declare.

Acknowledgements

SML, RC and CLC would like to acknowledge EPSRC for funding under grant number EP/N017374/1. GG would like to acknowledge funding under grant number EP/M028283/1. This research utilised the HADES/MIDAS facility¹⁰² at the University of Sheffield established with financial support from EPSRC and BEIS, under grant EP/T011424/1. We wish to acknowledge the Henry Royce Institute for Advanced Materials, funded through EPSRC grants EP/R00661X/1, EP/S019367/1, EP/P02470X/1 and EP/P025285/1, for JEOL JEM-F200 access at Royce@Sheffield.

Notes and references

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Footnote

† Electronic supplementary information (ESI) available: Bright-field TEM and EELS spectra for grains of doped perovskite samples. See DOI: 10.1039/d0ta05255h

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