Influence of hydration and dopant ionic radius on the elastic properties of BaZrO3

Solid State Ionics 344 (2020) 115130

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Influence of hydration and dopant ionic radius on the elastic properties of BaZrO3

T

Evgeniy Makagona, Rotraut Merkleb, Joachim Maierb, Igor Lubomirskya,

⁎

a b

Department of Materials and Interfaces, Weizmann Institute of Science, Rehovot, Israel Max Planck Institute for Solid State Research, Stuttgart, Germany

ARTICLE INFO

ABSTRACT

Keywords: Perovskites Proton-conducting oxides Acceptor doped BaZrO3 Elastic moduli Point defects Hydration

The influence of point defects (acceptor dopants, oxygen vacancies, protonic defects) on the macroscopic elastic properties of acceptor-doped BaZrO3 ceramics is investigated. Ultrasonic pulsed echo time of flight measurements are used to study the impact of dopant size, concentration and degree of hydration. Ceramics of BaXxZr1-xO3-x/2+δH2δ with X = Y, Sc and 0.05 ≤ x ≤ 0.2 were prepared by solid state reactive sintering with NiO sintering aid, added to achieve mechanical robustness sufficient for high-degree (58–71%) hydration without disintegration. Introduction of the dopants causes linear decrease in the Young's and shear moduli. By comparing the rate of decrease upon doping with Y3+ and Sc3+, the contribution of the lattice expansion was separated from the contribution of vacancy formation. Dissociative water incorporation also decreases the elastic moduli, however, for the case of Sc-doping the effect of hydration on the elastic moduli is much larger. All experimental data agree well with predictions by ab initio calculations.

1. Introduction Acceptor-doped proton conducting perovskites attract significant attention as they can be utilized in a wide range of applications from gas sensors to solid oxide fuel cells. Doped BaZrO3 (BZO) is particularly interesting as it combines high ionic conductivity and good chemical stability with respect to humidity and CO2 containing atmospheres (see e.g. [1]). BZO can sustain several types of point defects. By doping the lattice with an aliovalent ion, X3+ (X = Sc, In, Y, Dy, Eu, Gd, Sm etc.), which substitutes Zr4+, oxygen vacancies are created to compensate the dopant charge (defects are given in Kröger-Vink notation): x Y2 O3 + 2ZrZr + OOx

(1)

2YZr + VO•• + 2ZrO2 3+

B site doping of BZO with oversized acceptors such as Y was shown to cause lattice expansion which increases with dopant concentration and dopant ionic radius [2–6]. Since effective size of the formed oxygen vacancies is smaller than that of oxide ions [7], doping with undersized acceptors, such as Sc3+ slightly decreases the lattice parameter. Protonic point defects are formed in BZO by dissociative water incorporation according to [1]

H2 O + VO•• + OOx

2OHO•

(2)

whereby the H is attached to the lattice O by a short covalent bond to ⁎

form a hydroxide ion (OH•O ), and typically has two longer hydrogen bonds to neighboring hydroxide ions. As shown by thermogravimetric analysis (TGA), the degree of oxygen vacancy hydration is determined by dopant type and concentration as well as by water partial pressure and temperature (decreasing at high T) [2,3,8]. The rate of hydration is determined by ambipolar diffusion of VO•• and 2OH•O . The water chemical diffusion coefficient decreases from that of a proton diffusion coefficient (at low hydration) to the lower oxygen vacancy diffusion coefficient (at high hydration) [9]. Hydration leads to chemical expansion of the lattice, which was demonstrated by high temperature X-ray diffraction (XRD) [6,10–12] and dilatometry [1,13]. Lattice expansion of 1.2–3.3% was demonstrated in doped BZO and BaCeO3 (BCO) irrespective of variations in crystal symmetry (cubic vs. tetragonal in high dopant concentrations) and dopant type. The expansion of the lattice, driven by acceptor doping and subsequent hydration, can cause significant strain. While this is a critical factor influencing the long-term mechanical reliability of the material, only few studies deal with the influence of point defects on mechanical properties of perovskites. The effect of point defects on the mechanical properties of oxygen or proton conducting oxides is complex, as the strength of the bonds may be influenced by lattice expansion induced by oversized dopants but also more specifically by decrease of ion charges and formation of vacancies. One of the most studied oxygen conductors, CeO2, exhibits a

Corresponding author. E-mail address: [email protected] (I. Lubomirsky).

https://doi.org/10.1016/j.ssi.2019.115130 Received 13 March 2019; Received in revised form 24 July 2019; Accepted 25 October 2019 Available online 18 December 2019 0167-2738/ © 2019 Elsevier B.V. All rights reserved.

Solid State Ionics 344 (2020) 115130

E. Makagon, et al.

decrease in elastic moduli with increasing oxygen vacancy concentration, as measured by nanoindentation [14–16], high temperature oscillation [17] and flat punch [18]. When doped, the moduli decrease with increasing dopant concentration as shown by room temperature ultrasonic measurement of Ce1-xGdxO2-δ, the decrease is attributed to the combined effect of lattice expansion and defect formation [19]. In perovskites, hydration introduces protonic defects that influence the mechanical properties in addition to the oxygen vacancies. Ultrasonic [20] and oscillation [21] measurements determined Young's, shear and bulk modulus in Y doped BCO. No significant change in bulk modulus was observed, the shear modulus decreased slightly with increased dopant concentration and a slight increase in Young's modulus was demonstrated upon hydration. Nanoindentation measurements of Ba(Zr,Ce,Y)O3-δ showed changes in hardness and fracture toughness upon hydration, which mostly were attributed to hydration gradient induced strain [22]. We have recently investigated the mechanical properties of Y doped BZO to elucidate the impact of point defects on its shear, Young's and bulk moduli [23]. The elastic moduli were found to decrease linearly with increasing Y3+ concentration. Since in that study only relatively low hydration levels were accessible (~10%), no significant hydration dependence of the moduli was observed. The different contributions to the moduli decrease were separated using ab initio calculations (DFT), indicating that oxygen vacancies and protonic point defects are the dominant cause for bond weakening in Y doped BZO [23]. The present work significantly extends these investigations of elastic moduli in doped BaZrO3 in three points: (i) a different sintering procedure is used (“solid state reactive sintering” [24]) which yields samples that are mechanically more robust. (ii) Enabled by the improved sintering method, higher degrees of hydration are achieved without disintegration of the samples. The samples were able to withstand pure steam hydration, thus obtaining reliable conclusions on the hydration dependence of the moduli (effect of OH•O vs VO••). These data may be of crucial importance for future fuel cell design. (iii) To complement the data on Y3+ acceptor, which is much larger than Zr4+, samples doped with Sc3+, with a close size match to Zr4+, are investigated. This allows to experimentally distinguish different contributions to the decrease in elastic moduli (lattice expansion vs. decreased cation charge and corresponding VO•• formation).

degrees of hydration. Repeating the steam hydration procedure with the samples crushed particles yielded similar hydration degrees, indicating that the obtained hydration levels are not kinetically limited but represent equilibrium values. The water content of the as-prepared and hydrated samples was measured after the ultrasonic TOF measurements by TGA (Netzsch STA 449) of the crushed samples (sample weight 0.3–0.5 g, particle size ~300 μm, 60 mL/min dry N2, heating rate 5 K/min, buoyancy correction by measurement of inert Al2O3 sample). XRD of the hydrated samples (Fig. S1b,d) proves that the samples remain single-phase perovskite after the steam hydration. The surfaces of the pellets (sintered pellets for the nominally dry samples, and pellets after the steam hydration procedure for hydrated samples) were polished before XRD and ultrasonic measurements to remove surface impurity phases and obtain a smooth surface. XRD patterns of the ceramic pellets were acquired (TTrax Rigaku diffractometer in Bragg Brentano Θ/2Θ mode), and the lattice parameter of each sample was calculated (MDI Jade 2010 software [25]). Chemical composition of each sample was determined using energy dispersive spectroscopy (EDS) at 25 kV acceleration voltage (Bruker XFlash 6–60). Ultrasonic time of flight (TOF) echo, a well-established method for measuring sound velocity [19,20,23], was used to determine longitudinal (VL) and transversal (Vs) sound velocities with a USN 60L transducer (GE Inspection Technologies) coupled to the sample using high viscosity wax. Young's (E), shear (G) and bulk (B) moduli were deduced from the VL and Vs as [26]:

VL or T = G= E=G

d TOF vT2

3 VL2 VL2

E G 3 (3 G E )

B=

4 VT2 VT2

=

E 2G

1

(3)

where d is sample thickness, v is the Poisson's ratio and ρ is the metric density, which was determined by the Archimedes method with ethanol as the liquid medium. The values deduced from Eq. (3) were corrected for porosity with a dynamic model, which considers the scattering of sound waves by the pores [27]:

GD =

2. Experimental

BD =

Samples of Y- and Sc-doped BaZrO3 were prepared by a variation of the “solid state reactive sintering” method [24]. This method yields good densification and moderate grain growth (final grain size in the range of 5 μm) by formation of a transient Ba-(Y,Sc)-Ni-O liquid phase. For this method, respective amounts of BaCO3 (Alfa Aesar), ZrO2 (Tosoh TZ0), Y2O3 (Alfa Aesar), Sc2O3 (Alfa Aesar) were weighted and dry ball milled (Friatec zirconia mill with single ball of 5 cm diameter) for 1 h. The mixture was calcined for 6 h at 1100 °C in air, largely decomposing the BaCO3 but not completely forming the BaZrO3 perovskite phase. Then, 0.5 wt% NiO (Alfa Aesar) was added and the powder was milled for 24 h in a planetary mill (Fritsch Pulverisette 5, 200 rpm, ca. 30 g powder with 12 zirconia balls in a zirconia vial with 50 mL 2-propanol). After drying, the powders received a final 1 h dry ball milling, and were pressed isostatically into pellets. The pellets were sintered for 16 h at 1550 °C in air (heating rate 300 K/min, cooling rate 150 K/min) in Y stabilized ZrO2 (YSZ) crucibles covered with excess BaZrO3 powder (Sigma Aldrich) to prevent BaO loss. Hydration of 1–2 mm thick sample slices was carried out in a pure steam atmosphere (pH2O = 1 bar) in a quartz setup inserted with slight inclination into a tube furnace. The steam was generated by water slowly pumped with a peristaltic pump towards the hot zone of the setup where it evaporated; excess water condensed at the air cooled gas outlet. The temperature was lowered stepwise (6 h at 750 °C, 18 h at 700 °C, 24 h at 650 °C, 48 h at 600 °C, 96 h at 550 °C) to achieve high

F + F 2 4A C 2A 4GD B 4(1 p) GD 3p B

ED =

9 BD GD GD + 3 BD

(4)

where

A=

8(1

p) 3

C = 3(1 + p) B G F = (3

2p) B

8 + 4p 3

G

(5)

p is the relative porosity volume and the subscript “D” denotes the dynamic model correction. Three measurements were taken for each sample to ensure reproducibility. 3. Results and discussion The XRD patterns of both Y- and Sc-doped samples exhibited the dif– fraction peaks of pure perovskite phase (Pm 3m ), unindexed peaks were not observed (Fig. S1a and b). Fig. 1a shows the obtained lattice parameters as a function of dopant concentration for Y and Sc doped BZO. Lattice parameter increases linearly with increasing Y doping and decreases with Sc doping. Y3+ with ionic radius of 0.9 Å is significantly larger compared to Zr4+ sites (0.72 Å), hence lattice expansion occurs upon doping. Sc3+, on the other hand, with ionic radius of 0.745 Å is only slightly larger than Zr4+ [28]. The minor lattice expansion caused by the 2

Solid State Ionics 344 (2020) 115130

E. Makagon, et al.

b

Y Doped BaZrO3 as prepared : (1.31 ± 0.08)·10-3 Å/mol% Y Doped BaZrO3 hydrated : (1.9 ± 0.1)·10-3 Å/mol%

Y doped BaZrO3 as prepared

100

Y doped BaZrO3 hydrated

Sc Doped BaZrO3 as prepared : -(0.32 ± 0.01)·10-3 Å/mol%

Sc doped BaZrO3 as prepared

Sc Doped BaZrO3 hydrated : (0.17 ± 0.02)·10-3 Å/mol%

4.240

Levin et al. 2003

[OH]/[X] (%)

Lattice Parameter ( )

a 4.260

4.220

Sc doped BaZrO3 hydrated

80 60 40

4.200 20 4.180

0

5

10

15

20

0

25

5

10

mol% dopant

15

20

25

mol% dopant

Fig. 1. (a) Lattice parameters of BaXxZr1-xO3-x/2+δH2δ X = Y, Sc. Undoped lattice parameter was taken from Levin et al. [30]. (b) Degree of hydration for BaXxZr1-xO3-x/2+δH2δ X = Y, Sc. Y doped BaZrO3 as prepared: -(1.11 ± 0.09)% Y doped BaZrO3 as prepared by SPS [23]: -(0.96 ± 0.03)% Undoped BaZrO3: Yamanaka et al. 2003 [32]

260 240 220 200 0

5

10 15 mol% Y

20

Y doped BaZrO3 as prepared by SPS [23]: -(1.00 ± 0.04)% Undoped BaZrO3: Yamanaka et al. 2003 [32]

100 90 80 0

d0.28

c 176

Y Doped BaZrO3 as prepared: -(0.8 ± 0.2)% Y Doped BaZrO3 as prepared by SPS [23]:-(0.7 ± 0.3)%

Y Doped BaZrO3 hydrated: -(0.8 ± 0.1)%

5

10 15 mol% Y

20

25

Y Doped BaZrO3 as prepared: (0.5 ± 0.2)% Y Doped BaZrO3 Hydrated: (0.28 ± 0.09)% Y Doped BaZrO3 as prepared by SPS [23]: (0.2 ± 0.1)%

0.27

Poisson`s ratio

168

Y doped BaZrO3 hydrated : -(1.09 ± 0.08)%

110

70

25

Bulk Modulus (GPa)

180

Y doped BaZrO3 as prepared : -(1.17 ± 0.09)%

b 120

Y doped BaZrO3 hydrated: -(1.05 ± 0.08)%

Shear Modulus (GPa)

Young`s Modulus (GPa)

a 280

160

0.26

152

0.25

144

0.24

136 0

5

10 15 mol% Y

20

25

0.23

0

5

10 15 mol% Y

20

25

Fig. 2. Elastic moduli of BaYxZr1-xO3-x/2+δH2δ as a function of Y concentration. (a) Young's modulus E, (b) Shear modulus G, (c) Bulk modulus B, (d) Poisson's ratio ν. Dashed lines are linear fits to the data with compositional error; the slope gives the moduli relative change with respect to extrapolated undoped values. For comparison, the regression lines for nominally dry BaYxZr1-xO3-x/2+δ samples sintered by spark plasma sintering from [23] are indicated by a dashed grey line.

doping cannot compensate for the lattice contraction caused by the formation of oxygen vacancies, and an overall lattice contraction occurs Å [7,29]. The calculated slope of (1.31 ± 0.08) 10 3 mol % Y for the as prepared Y doped BZO is within the error margins of the experimental and DFT Å results [2,4]. The calculated slope of (0.32 ± 0.01) 10 3 mol % Sc for the as prepared Sc doped BZO is with good agreement with the DFT lattice parameters reported in ref. [23] (Fig. S2a). Extrapolating the lattice parameter of the as prepared samples to that of an undoped BZO yields 4.1935 ± 0.0006 Å closely matching the value reported in ref. [30]. A lattice expansion of 0.05–0.25% was observed for low and high dopant concentrations, respectively, in the hydrated state. Although both, oxygen vacancies and proton interstitials, are smaller than an oxygen ion, the contraction for the vacancy is significantly larger than

for the proton, hence oxygen vacancy filling causes net expansion [7]. Lattice expansion values are consistent with values reported in refs. [5, 12], and further support that the steam hydration leads to high degrees of hydration even for thick pellets. The lattice expansion due to hydration increases with increasing dopant concentration as expected. The dependence of lattice parameter on Y concentration is identical to that reported in ref. [23] for samples prepared without NiO sintering aid. The absolute values are slightly lower for the samples discussed in the current work, which were prepared with this sintering aid (Fig. S2b), but still within the range of literature data [2–4]. A more detailed discussion of the effects of the NiO sintering aid on BZO bulk defect chemistry can be found in [31]. NiO was used to lower BZO's sintering temperature and promote densification and grain growth in a 3

Solid State Ionics 344 (2020) 115130

E. Makagon, et al.

“reactive sintering” approach [24] which involves the transient formation of a Ba-(Y,Sc)-Ni-O liquid phase. Adding 0.5 wt% of NiO to the sintering process improved the mechanical stability for hydration compared to previous protocols [23] and yields larger grains in the range of several micrometers (Fig. S3). Although no secondary phases were identified in the XRD measurements, their presence cannot completely be ruled out if they are present only in trace quantities. Achieving a fully hydrated state for dense, millimeter-thick samples is technically challenging. Higher temperature promotes faster proton and oxygen vacancy diffusion but reduces the maximal proton concentration and vice versa. To achieve the optimal balance between defect diffusion and concentration, hydration was carried out in pure steam (pH2O = 1 bar) instead of only humidified gas (20 mbar). This yields higher equilibrium proton concentrations at elevated temperatures where diffusion is still sufficiently fast. The use of the reactive sintering method significantly improved the mechanical robustness of the samples. Hydration levels of 58–71% of the nominal acceptor concentration were achieved (Fig. 1b) while keeping the sample intact. The presence of the NiO sintering aid decreases the effective acceptor concentration and thus the maximum proton uptake, while leaving the enthalpy and entropy of the hydration reaction unchanged [31]. The results of the unhydrated samples shown below demonstrate that NiO addition does not perceptibly change the elastic properties. The proton uptake of the hydrated samples is still sufficiently high to observe changes in the lattice structure (cf. the lattice parameter measurements (Fig. 1a)). The as prepared samples exhibited ~15% hydration, due to water uptake during the time interval between sintering and measurement, in which they are exposed to ambient air. The relative densities for Y and Sc doped BZO are presented in Fig. S4. The densities are > 94% of the theoretical one as calculated according to the lattice parameters. The density decreases with increasing dopant concentration as demonstrated previously [2,23]. While the

b120

Sc doped BaZrO3 as prepared: -(0.22 ± 0.04)% Sc doped BaZrO3 hydrated: -(0.4 ± 0.1)%

Shear Modulus (GPa)

Young`s Modulus (GPa)

a 280

lattice volume increases upon hydration, it cannot compensate for the weight gain from water uptake, and a slight increase in the relative density is observed. Surface porosity was observed in samples with relative densities lower than 100% (Fig. S3), confirming the reliability of the density measurements. All samples exhibited high quality multiple echoes for longitudinal and transverse waves (Fig. S5). Samples that produced less than ten clear echoes were excluded from further examination. 6 mol% Sc doped BZO sample was within 0.9% from theoretical density, hence no porosity correction was applied to it. The data deduced from Eq. (3) for the rest of the samples were subjected to porosity correction by Eqs. (4) and (5) [27]. The shear, Young's, bulk moduli and Poisson's ratios calculated from ultrasound TOF measurements are presented in Fig. 2 for Y doped BZO and in Fig. 3 for Sc doped BZO. Both shear and Young's moduli decrease linearly with increasing dopant concentration with an average slope of −1.14 ± 0.09% per mol% Y (Fig. 2a–b) and –0.24 ± 0.04% per mol% Sc (Fig. 3a–b) for the samples in the as prepared state. In Y-doped samples, this decrease in moduli occurs due to weakening of the bonds by lattice expansion, the decrease in cation charge from +4 to +3 and oxygen vacancy formation that decreases the number of chemical bonds per unit cell. The absolute values as well as the slopes for Y doped BZO matches the published values in ref. [23] for samples that had been sintered by “spark plasma sintering” (without NiO addition), indicating that the use of 0.5 wt% NiO sintering aid had no significant effect on the elastic properties(Fig. 2, grey regression lines). The comparison of Y- and Sc-doped samples allows us to identify the contribution from lattice expansion. Y3+ is larger by ~16% than Zr4+ while Sc3+ is larger only by ~4% which makes the lattice parameter of Sc-doped BZO largely independent of the Sc concentration. Nevertheless, Sc-doped BZO exhibits a decrease of the moduli with acceptor concentration (caused by decreased cation charge and VO•• formation), but with a smaller slope. Under the assumption that the individual contributions are additive and transferrable to different

100

240 220 200

Bulk Modulus (GPa)

c176

Sc doped BaZrO3 hydrated : -(0.5 ± 0.1)%

110

260

180 0

Sc doped BaZrO3 as prepared : -(0.27 ± 0.03)%

5

10 15 mol% Sc

20

90 80 70 0

25

d0.28

Sc Doped BaZrO3 as prepared Sc doped BaZrO3 hydrated

5

10 15 mol% Sc

20

25

Sc Doped BaZrO3 as prepared: (0.25 ± 0.06)% Sc Doped BaZrO3 Hydrated: (0.63 ± 0.08)%

0.27

Poisson`s ratio

168 160

0.26

152

0.25

144

0.24

136 0

5

10 15 mol% Sc

20

25

0.23 0

5

10 15 mol% Sc

20

25

Fig. 3. Elastic moduli of BaScxZr1-xO3-x/2+δH2δ as a function of Y concentration. (a) Young's modulus E, (b) Shear modulus G, (c) Bulk modulus B, (d) Poisson's ratio ν. Dashed lines are linear fits to the data; the slope gives the moduli relative change with respect to dopant concentration change. 4

Solid State Ionics 344 (2020) 115130

E. Makagon, et al.

1.5

Y doped BaZrO3 as prepared: Experimental(

% per mol% dopant

Y doped BaZrO3 hydrated: Experimental(

1.0 0.5

Sc doped BaZrO3 as prepared: Experimental( Sc doped BaZrO3 hydrated: Experimental( Gd doped CeO2: Yavo et. al 2016 [19] (

), DFT( ), DFT( ), DFT( ), DFT(

clearly show a significantly smaller dopant concentration dependence of the moduli. The numerical value of the slope from experiment (−0.22 ± 0.04% and –0.4 ± 0.1% per mol% Sc for unhydrated / hydrated samples) is moderately smaller than from DFT (−0.58 ± 0.07% and –0.74 ± 0.07% per mol% Sc for unhydrated / hydrated samples). The decrease in elastic moduli due to acceptor-doping-induced vacancy formation was observed previously in CeO2 [14,15,19] with a dopant concentration dependence of 0.48% per mol% Gd [19]. This concentration dependence is lower than for Y doped BZO by more than a factor of two (Fig. 4). The difference may be attributed to two factors: (i) The mechanical properties of perovskites are more susceptible to oxygen vacancy formation than fluorites, extensively discussed in ref. [23]. (ii) While Y3+ is larger than Zr4+ in Y doped BZO by ~25%, Gd3+ is larger only by ~16% compared to Ce4+ in CGO [28]. This difference in relative sizes may cause a smaller lattice expansion effect in CGO, hence a smaller moduli change than in Y doped BZO. The good agreement of experimental and DFT results demonstrated here for BZO encourages the use of either of the methods for investigating the influence of point defects also in other material systems. While ultrasonic TOF measurements allow for high-precision moduli determination, they are very sensitive to porosity and (micro) cracks in the samples. Thus, in cases where the preparation of sufficiently large crack-free samples is difficult, DFT calculations represents an alternative, the validity of which was demonstrated here. DFT calculations come with challenges as well (large supercells, necessity to probe various defect configurations etc., cf. [23] and Supplementary there), nonetheless, they allow one to separate individual contributions to the moduli change. The elastic moduli are relevant not only for the mechanical properties of a material. They could indirectly affect also the electrical properties of grain boundaries. In many cases, grain boundaries of electrolytes are blocking for oxygen vacancies and protons because of space charge effects (positive grain boundary core charge and corresponding positive carrier depletion in adjacent space charge zones, see e.g. [33]. The positive core charge is often caused by structural conflicts (“overcrowded situation”) which leads to accumulation of positively charged oxygen vacancies in the core (cf. correlation of vacancy segregation energies with compressive strain [34]. A decrease of the elastic moduli for high dopant concentration may be beneficial, as it lowers the energetic cost of local structural distortions in the boundary core and could thus decrease oxygen vacancy segregation and positive core charge.

) ) ) )

)

0.0 -0.5 -1.0 -1.5

Young`s

Shear

Bulk

Poisson`s

Fig. 4. Comparison of the relative experimental change of the elastic moduli with dopant concentration for BaYxZr1-xO3-x/2+δH2δ and BaScxZr1-xO3-x/2+δH2δ with DFT results from ref. [23] and with Ce1-xGdxO2-x/2 data from ref. [19].

dopants, this indicates that for Y doping, the acceptor doping and vacancy formation are responsible for a modulus decrease by −0.24 ± 0.04% per mol% of the dopant, while the lattice expansion accounts for the larger part of 0.9 ± 0.09% per mol% of the dopant. Undoped BZO Young's and shear moduli were extrapolated from the regression lines in Fig. 2 and showed good agreement with previously published values [32]. Lattice hydration causes shear and Young's moduli decrease by 2–4% for Y doped BZO and by 6–10% for Sc doped BZO throughout the whole dopant concentration range. The fact that oxygen vacancy hydration causes a slight lattice weakening is at first glance counterintuitive, but was predicted by DFT calculations [23]. Hydration fills the oxygen vacancies, which restores the missing bonds, but on the other hand, oxide ions OOx are replaced by protonic point defects OHO• (hydroxide ions). This lowers the nominal charge of –2e of an oxide ion to –1e of a hydroxide ion, weakening the electrostatic bonding. The lattice expansion caused by hydration weakens the bonding further. For Y doped BZO the dependence of shear and Young's moduli on Y concentration does not change significantly upon hydration (−1.07 ± 0.08% per mol% Y). On the other hand, for Sc the dependence increases from −0.22 ± 0.04% per mol% Sc for the as-prepared to −0.45 ± 0.1% for the hydrated samples. While the absolute value of lattice expansion upon hydration is the same for both dopants, hydration has a stronger relative effect on BZO when it is doped with smaller size dopants. While keeping the trend demonstrated by shear and Young's modulus, the bulk modulus in Y doped BZO exhibits a slightly lower concentration dependence for both as-prepared and hydrated samples (−0.8 ± 0.2% per mol% Y) (Fig. 2c) as observed for Ni-free samples in [23]. In Sc doped BZO, no significant change in bulk modulus was observed with dopant concentration and hydration (Fig. 3c). Poisson's ratio increases linearly for both Y (Fig. 2d) and Sc (Fig. 3d), and the values of the hydrated samples are slightly larger than for the as prepared state. However, as bulk modulus and Poisson's ratio are rather susceptible to errors in shear and Young's moduli, no significant conclusions could be drawn from their behavior. All moduli changes are summarized in Fig. 4 and compared to DFT values [23]. The presented experimental results for the absolute moduli and Poisson ratio as well as for the dependence on dopant ion (Y or Sc) and concentration agree very well with the preceding DFT calculations [23]. Experimental and DFT results agree very well with the finding that hydration leads to a perceptible but small decrease of the modulus values. The DFT calculations allow us to trace this, at first glance, unexpected behavior back to the combination of chemical lattice expansion and the effect of the protonic defects, as was discussed in [23]. Also the agreement of the slope of the dopant dependence from experiments and DFT is good for Y doped BZO (slope of Young's modulus −1.11 ± 0.09% per mol% Y (exp.) vs −0.94 ± 0.07% for dry and –1.05 ± 0.08% (exp.) vs. –1.21 ± 0.04% for hydrated Y doped BZO). For Sc doped BZO, both experiment and DFT

4. Conclusions Ceramics of BaXxZr1-xO3-x/2+δH2δ with X = Y, Sc and 0.05 ≤ x ≤ 0.2 were prepared by solid state reactive sintering with NiO sintering aid, added to achieve mechanical robustness sufficient for high-degree (58–71%) hydration without disintegration. The elastic moduli decrease linearly with increasing dopant concentration, which is attributed to bond weakening induced by lattice expansion and point defect formation. Sc3+ doped samples, with a small size mismatch to Zr4+, were measured and compared to previously reported data on larger size mismatch dopant (Y3+). A strong decrease in the dopant concentration dependence was observed when changing from Y3+ to Sc3+, experimentally showing the effect of lattice expansion. Interestingly, the hydration of the oxygen vacancies causes a further slight decrease in elastic moduli. Protonic point defects tend to weaken the BeO bonding network because oxide ions are partially replaced by hydroxide ions and hydrogen bond formation induces additional lattice distortions. Hydration was shown to have a stronger relative effect on BZO when doped with smaller size dopants like Sc3+. All experimental results are in close agreement with previous DFT calculations [23]. The present work clearly demonstrates the significance of structural characteristics such as point defect type, size and concentration on the macroscopic mechanical properties of BaZrO3 perovskites. Although BaZrO3 was considered as a case study, these findings are expected to hold also for other point defect containing ABO3 perovskites and may be of crucial importance for future fuel cell design. 5

Solid State Ionics 344 (2020) 115130

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Acknowledgements

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