Improvement of Ionic Conductivity in A-site Lithium Doped Sodium Bismuth Titanate

Duke P.C. Shih, Ainara Aguadero, Stephen J. Skinner*

Department of Materials, Imperial College London, Prince Consort Road, London SW7 2BP, UK

* Corresponding author.

Email address: s.skinner@imperial.ac.uk (Stephen J. Skinner).

Abstract

Oxide-ion conductors play a significant role in various applications such as solid oxide fuel cells (SOFCs), oxygen separation membranes and sensors. Recently, high ionic conductivity (~1 x 10-4 S cm-1 at 600 oC) was found in sodium bismuth titanate (NBT), which originates from oxygen vacancies compensating the introduced Bi-deficiency. By providing pathways with low diffusion barriers, the highly polarizable Bi3+

ions with 6s2 lone pair electrons and weak Bi-O bonds are also beneficial for the migration of oxygen ions. Here we report the influence of lithium doping on the electrical properties of NBT. The optimal doping level of 4 at% Li on the Bi-site improves the ionic conductivity by one order of magnitude to ~ 7 x 10 -3 S/cm at 600 oC without changing the conduction mechanism, which could be attributed to an increase in the oxygen vacancy concentration based on an acceptor doping mechanism. A further increase in Li content does not improve the total conductivity. Oxygen diffusion data were acquired by the Isotope Exchange Depth Profile (IEDP) method in combination with Secondary Ion-Mass Spectrometry (SIMS). The oxygen self-diffusion coefficients (e.g. 7.04 x 10-9 cm2 s-1 at 600 oC) are in excellent agreement with the values derived from impedance spectroscopy data, suggesting that the oxygen ions are the main charge carriers in the system. Furthermore, a degradation test was performed for 100 hours under a variety of atmospheres, showing only a slight decrease in conductivity in both air and oxygen atmospheres attributed to the loss of material from the A-site. Comparison with other oxide-ion conductors indicates that Li-doped NBT materials are promising candidates for intermediate temperature SOFC applications.

Keywords: Oxide-ion conductors, Sodium bismuth titanate, Doping, Non-stoichiometry, Oxygen diffusion

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1. Introduction

Oxide-ion conductors are widely applied in important technological devices such as solid oxide fuel cells (SOFCs), oxygen sensors, oxygen pumps, oxygen separation membranes, etc. [1-3]. Research in oxide-ion conductors has attracted considerable attention due to the desire to develop new clean energy sources [4].

Recently, a new class of oxide ion conductors based on the perovskite-type sodium bismuth titanate (Na0.5Bi0.5TiO3, NBT) has been investigated [5, 6]. High levels of oxide ion conduction, originating from Bi deficiency and from oxygen vacancies generated during the material processing, was described according to the following Kröger-Vink equation:

2 BiBiX +3OO

X →2VBi' ' +3VO

∙ ∙+Bi2 O3.

In addition, the highly polarizable Bi3+ ions with their 6s2 lone-pair electrons and weak Bi-O bonds were postulated to aid the migration of oxygen ions by providing pathways with low diffusion barriers [5, 7-9]. The ionic transport number in NBT was determined to be ti > 0.9 at 600-700 oC by electromotive force (EMF) measurements [5]. A modelling study showed that a migration path for oxygen vacancies with minimum energy barriers exists at the saddle points between two A-sites and a Ti-site when both A-sites are occupied by Bi ions [7]. However, there is no experimental evidence to support this result, and the rate-limiting step in oxygen ion migration in un-doped NBT is presumed to be the Na-Bi-Ti saddle points.

The further improvement of ionic conductivity can be achieved by acceptor doping. By doping 2 at% Sr2+ from group IIA on the Bi-site, the bulk conductivity can be enhanced up to 5 x 10-3 S cm-1 at 500 oC [10]. Experimentally, the highest total conductivity can be observed by doping Mg2+ on the Ti-site, enhancing bulk conductivity to 3.98 x 10-3 S cm-1 at 500 oC or further to 7.28 x 10-3 S cm-1 at 600 oC [5].

He et al. [7] employed first-principle calculations and molecular dynamics to demonstrate a doping strategy for NBT materials. Using their model, it was suggested that the ionic conductivity of NBT-based materials could be further improved by partially substituting dopants on the A-site. The computational results indicated that acceptor-type Bi-site dopants, such as Mg2+ for Ti4+, increase the oxygen migration barriers by binding with oxygen vacancies, so A-site acceptor doping is more desirable to achieve high conductivity. According to the calculations, doping 4% Na or K on the Bi-site can lead to higher bulk conductivity than Mg-doped NBT at the same oxygen vacancy concentration. The enhancement of conductivity is suggested to be due to the disordered A-site sub-lattice. The disordered A-site sub-lattice can adopt different local atomistic configurations to accommodate the electrostatic and strain field of the dopants, improves

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the phase stability of the doped materials, and reduces the undesired binding with B-site dopants [7]. So far, there has been no experimental evidence to prove the effect of Na and K dopants.

Here we report the preparation of Li-doped NBT materials, in which Li+ replaces Bi3+

ions with the creation of oxygen vacancies, and the effect that this has on the oxide-ion conduction was investigated. The results are discussed from the perspective of the ionic size mismatch between the dopant and the host element. The ionic size of Li+ is smaller than that of Bi3+, potentially leading to a decrease in the unit cell volume and an increase in the activation barrier for oxygen mobility. The possibility of Li mobility due to Li+

doping was also studied at room temperature given its similar structure to a superior perovskite Li-ion conductor, La2/3-xLi3xTiO3, where the Li+ ions migrate via vacant A-sites.

2. Experimental Section

Powders with nominal starting compositions of Na0.5Bi0.5-xLixTiO3-δ (for x = 0.02, 0.04 and 0.06) were synthesized with a modified Pechini citrate-nitrate synthesis route [11]. Stoichiometric masses of NaNO3 (Alfa Aesar, 99.999%), Bi2O3 (Alfa Aesar, 99.9%), LiNO3 (Alfa Aesar, 99.9%), and C10H14O5Ti (Alfa Aesar), were mixed with citric acid (10 mass%, 0.542 M, ~200 mL, Sigma Aldrich) and an addition of 10 mL of nitric acid (68%, VWR Chemicals) to dissolve the solids. The mixture was dehydrated by stirring on a hotplate until the solution started to form a yellow gel. This gel was then cooled naturally to ambient temperature. Once cool the gel was combusted in a decomposition furnace at 650 oC for 12 hours to remove nitrate and organic groups. The resultant product was ground into a fine powder with a mortar and pestle, compacted into pellets in a uniaxial steel die and then isostatically pressed for 30 seconds at 300 MPa. The pellets were finally sintered in air at 1050 oC for 2 and 4 at% Li-doped NBT and 1100 oC for 6 at % Li-doped NBT and covered by calcined powders of the same composition to minimize Na-, Bi- and Li-loss during the firing process. The density of the sintered pellets was measured using the Archimedes' method.

X-ray powder diffraction (XRD) using a Bruker D2 Phaser diffractometer with Cu Kα

radiation (λ = 1.5418 Å) was used to investigate the phase evolution of the sintered and crushed samples. Ceramic microstructure and chemical composition were studied by using a combination of scanning electron microscopy (SEM, JEOL 6400, JEOL Ltd, Tokyo, Japan) and energy dispersive X-ray spectroscopy (EDX, Oxford Link ISIS, Oxford Instruments Ltd, Oxfordshire, UK). Ceramic samples for SEM analysis were polished, or polished and thermally etched at 1000 oC for 1 hour, before being coated with carbon. EDX analysis was obtained with an average of 15-20 randomly selected areas. The compositional depth profile of the polished pellets was characterized by means of time-of-flight secondary ion mass spectrometry (ION-TOF GmbH, Münster,

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Germany) using O2+ primary ions and a low-energy Cs+ sputter gun to remove the surface

materials, allowing depth profiles to be generated on acceptable time scales. The optional electron gun was used to compensate for charging effects. The intensities of secondary ions, 7Li+, 23Na+, 48Ti+, and 209Bi+ were collected as a function of sputtering time. The depth of sputter crater was subsequently measured by using a surface profiler (Zygo new View 200 3D interferometer).

AC Impedance spectroscopy (IS) was performed using a Solartron 1260 frequency response analyzer (Solartron Analytical, UK) with an AC amplitude of 100 mV, over a frequency range from 10 mHz to 10 MHz. Prior to measurements, Au or Pt paste electrodes were coated onto the polished surfaces of the samples and fired at 900 oC for 2 hours in air. Impedance data were first collected at room temperature and continued from 300 to 750 oC with an interval of 25 oC during heating and cooling. The data were normalized to sample geometry (thickness/area of pellet). High frequency instrumental-related (impedance analyzer, lead, and sample rig) inductance effects were corrected by performing a short circuit measurement. To analyze the impedance data, an equivalent circuit of 3 resistor-constant phase elements (R-CPE) connected in series was constructed. The fits were considered valid only when the χ2 value was below 10-3 for the full data range demonstrating the quality of the fitting. The error of each parameter was also minimized.

A combination of oxygen isotope labeling and line-scan dynamic secondary ion mass spectrometry (ION-TOF GmbH, Münster, Germany) was utilized to determine the oxygen tracer diffusion coefficient, D*. Oxygen isotope exchanges were performed in the temperature range of 450-700 oC, at ~200 mbar oxygen pressure, and with an exchange time of 2 hours. The addition of a platinum coating, sintered at 1000 oC for 2 hours, facilitated surface oxygen exchange process. Each of the samples was equilibrated in 16O2 for a period of 10 times the exchange time, as discussed in ref [12-14] prior to the exchange.

3. Results & Discussion

3.1 Phase Purity

XRD patterns of all powder samples ground from sintered pellets are shown in Fig.1 (a). The refinements show that all of the samples could be indexed in the rhombohedral space group, R3c, so no detectable impurity phase was introduced with the addition of the Li dopant (Fig.1 (b) and Supporting Information S1). No extra peaks were observed within the limit of resolution of the X-ray diffractometer, and the lattice parameters and cell volume for each sample are extracted from Rietveld refinements (using the Fullprof software suite [15]) in Table 1. The Li-doped NBT materials have a smaller unit cell

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volume than NBT due to the smaller ion size of Li+ (~1.20 Å) when compared with radii of Na+ and Bi3+ (~1.39 Å [16]). The ionic size of Li+ and Bi3+ in XII-coordination are not available. The value for Bi3+ is set to equal Na+ (1.39 Å), because the mismatch between Na+ and Bi3+ is negligible in NBT [17]. The r(Li+) is calculated using the linear relationship between the ionic radius and the coordination number. The ionic radius for XII-coordination of Li+ is extrapolated using the slope above VI-coordination in Li+ [18]. As the doping level increases, the lattice parameter of the c-axis shrinks and causes a further reduction in unit cell volume. The reduction in the cell volume may increase the activation barrier for oxygen mobility. In short, Li was successfully introduced into the lattice of the NBT material without causing any significant lattice distortion.

Fig 1 (a) XRD patterns of NBT and Li-doped NBT after sintering; (b) XRD pattern and Rietveld refinement of 2 at% Li-doped NBT.

Table 1. Refined lattice parameters of Li-doped NBT ceramics compared with literature data for the parent phase. Rwp, RP and χ 2 are the refinement reliability factors.

Lattice Parameters a (Å) c (Å) V (Å3) RWP, RP, χ2

NBT [19] 5.4882(3) 13.502(1) 352.21(2)2% Li-NBT 5.4883(4) 13.484(1) 351.74(3) 5.86, 4.55, 2.064% Li-NBT 5.4886(3) 13.480(1) 351.68(5) 6.74, 5.12, 2.276% Li-NBT 5.4880(1) 13.479(2) 351.58(4) 6.45, 4.46, 1.67

3.2 Microstructure and Compositional Analysis

The density of all pellets was determined to be > 95% of the theoretical values calculated from the diffraction data. As shown in Fig 2 SEM images of the thermally etched samples shows smaller grains in the Li-doped NBT materials than in the un-doped

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sample. NBT ceramics exhibit grain sizes of ~ 10-20 μm. For the Li-doped samples, the grain size was found to be in the range of ~ 5-10 μm. It is noted that Li-doping improves the sinterability of the ceramics. With 6 at% Li-doping (nominal compositional Na0.5Bi0.46Li0.04TiO2.81), the density of the ceramics was found to reach ~98% (NBT~95%) under the same material processing.

Fig 2. SEM micrographs of thermally-etched surfaces showing the reduction in grain size with increasing Li content.

Samples were analyzed by EDX, and the cation ratios, normalized to 100%, are shown in Table 2. For stoichiometric NBT, the theoretical percentage of Na, Bi and Ti is 25.0 at%, 25.0 at%, and 50.0 at%, respectively. The measured composition of NBT was close to the theoretical cation ratio within instrument resolution and standard deviations. For the Li-doped samples, the composition of Li signal could not be detected by the EDX technique . Since Li content is undetectable, the measured composition of all cations is higher than their actual values. However, changing the Li doping level leads to an evident

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variation in the concentration of Bi. The composition of Bi is inversely proportional to the Li doping level, while the composition of Na remains relatively stable. This suggests that increasing Li doping compensates for the decrease in the Bi composition and overall A-site composition remains almost identical.

Table 2. Analyzed composition of un-doped and Li-doped NBT ceramics determined by EDX.

Na (at%) Bi (at%) Ti (at%) Li (at%)NBT 25.0 (±0.4) 25.2 (±0.2) 49.8 (±0.4)

N/A2% Li-NBT 25.8 (±0.7) 21.2 (±0.2) 53.0 (±0.4)4% Li-NBT 25.2 (±0.1) 20.9 (±0.4) 53.9 (±0.2)6% Li-NBT 25.4 (±1.2) 20.3 (±0.4) 54.3 (±0.4)

To further verify the Li doping in the samples, SIMS was employed to generate compositional depth profiles. Fig. 3 shows a depth profile and Li ion mapping for Li-doped NBT materials. Li ion mapping indicates that the Li content is evenly distributed in the ceramics. In the depth profiles for Li-doped samples, the intensities of all elements are stable in the whole measurement range, indicating a homogenous microstructure. The order of intensity counts in each sample has a similar trend (23Na+ > 48Ti+ > 209Bi+) while the intensity of Li was detected in the Li-doped samples, suggesting the presence of Li content. EDX analysis on the polished surface (without thermal etch) and compositional depth profiling confirms the Li doping in the materials.

Fig 3. a) SIMS depth profiles for 6 at% Li-doped NBT and b) Li ion mapping normalized to total intensity counts.

3.3 Electrical Properties

At room temperature, the Li-doped NBT materials are highly resistive, so the electrical performance of Li-doped NBT ceramics was focused on the high temperature

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region (> 300 oC). Fig. 4 shows the impedance spectrum of the 2 at% Li-doped NBT sample measured at 600 oC as an example. There are two arcs observed within the measured frequency range. The resistivity of the intermediate frequency arc is ~ 0.15 kΩ cm. The capacitance (C) associated with the arc is ~ 10 -9 F, indicating this arc is associated with the grain boundary. The arc in the low frequency region had a capacitance of 10-6 F, representing the response from the electrode. The arc corresponding to the bulk response is outside of the frequency range measured. With the ohmic contribution taken into account, the bulk response is represented by the high frequency intercept.

Fig 4. Impedance spectrum of the 2 at% Li-doped NBT ceramic measured at 600 oC in air. The inset figure shows the high-frequency data on an expanded scale.

At 600 oC, NBT ceramics exhibit large grain boundary resistance (~8000 Ω∙cm), and Li-doping can minimize the grain boundary contribution to the total resistivity as shown in Table 3. As Li doping increases, both bulk and grain boundary conductivities continue to improve until they reach the optimal doping level at 4 at%. A-site Li-doping enhances both σbulk and σgrain boundary by more than one order of magnitude compared to un-doped NBT. To achieve high oxide-ion conductivity, the material must have high levels of oxygen vacancies to facilitate charge carriers with high oxygen ion mobility. When a Li+

ion is introduced to replace Bi3+ in NBT, oxygen vacancies are generated according to the following Kröger-Vink equation:

Li 2O →2LiBi'' +2VO

∙∙ +OOX

Acceptor doping is believed to be the major source of oxygen vacancies. Compared to Mg2+ doping, Li+ doping would double the creation of oxygen vacancies at the same

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doping level. It is believed that Li+ doping increases the oxygen vacancy concentration and thus improves both the bulk and grain boundary conductivity.

Table 3. Bulk and grain boundary conductivity of Li-doped and un-doped NBT.

σbulk (S cm-1) σgrain boundary (S cm-1)NBT 1.25 x 10-3 1.15 x 10-4

2% Li-NBT 6.97 x 10-3 6.65 x 10-3

4% Li-NBT 1.77 x 10-2 1.26 x 10-2

6% Li-NBT 1.31 x 10-2 4.88 x 10-3

As depicted in Fig. 5 (a), an Arrhenius plot of the temperature dependence of the total conductivity, σtotal where σtotal = 1 / ρtotal, demonstrates that Li-doping also decreases the activation energy for conduction. The reduced lattice volume of Li-doped ceramics does not inhibit the migration of oxygen ions. However, a change in activation energy is observed at around 500 oC. The crystallographic phase transition temperatures of NBT remain unclear. It is believed that NBT undergoes a sequence of phase transitions from rhombohedral (R3c) at room temperature to a tetragonal phase (P4bm) at 250oC and then to a cubic phase (Pm3m) at 540oC [20]. However, the coexisting rhombohedral (R) and tetragonal (T) phases do not transform into a single tetragonal phase until ∼400 °C, and the material finally undergoes a tetragonal to cubic (C) transition at ∼520 °C [9]. Therefore, this change in activation energy could be attributed to the phase transition from tetragonal to cubic or the existence of two phases, but further evidence such as volume fractions of T and C phases in NBT and Li-doped NBT from X-ray diffraction or TEM are required.

A comparison of total conductivity for 4 at% Li-doped NBT and other known oxide ion conductors is given in Fig. 5 (b). The Li-doped NBT exhibits superior ionic conductivity in both the low and intermediate temperature ranges and is comparable to that of the best oxygen-ion conducting electrolytes. In comparison with other bismuth based conductors, the conductivity of Li-doped NBT is lower at high temperatures, but it does not show the drawbacks of high corrosion activity and low mechanical strength.

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Fig 5. (a) Arrhenius plot of total conductivity for un-doped and Li-doped NBT materials and (b) total conductivity of Li-doped, Mg-doped [5] and un-doped NBT in comparison with other oxide-ion conductors: CGO (Ce0.9Gd0.1O1.95) [21]; YSZ (Zr0.92Y0.088O1.96) [22]; LSGM (La0.9Sr0.1Ga0.9Mg0.1O2.9) [23]; 20ESB (BiO1.5)0.8(ErO1.5)0.2 [21] and 8D4WSB (BiO1.5)0.88(DyO1.5)0.08(WO3)0.04 [21].

3.4 18O Tracer Diffusion Measurements

The improvement of oxide ion conduction in 4 at% Li-doped NBT was confirmed by 18O tracer diffusion measurements using isotopic exchange and secondary ion mass spectrometry (SIMS), as shown in Fig 6 (a). The surface exchange coefficient (k*) is invalid in these measurements due to the previously discussed application of Pt to the sample surfaces to facilitate oxygen reduction. At 600 oC, the tracer diffusion coefficient (D*) was found to be 7.04 x 10-9 cm2 s-1 and is one order of magnitude higher than that of NBT (5.24 x 10-10 cm2 s-1) at 608 oC [5]. The ionic conductivity calculated from the diffusion data using the Nernst-Einstein equation, along with the total conductivity obtained from AC impedance spectroscopy is given in an Arrhenius plot in Fig 6 (b). The calculated oxide ion conductivity for 4% Li-doped NBT at 600 oC is 4.28 x 10-3 S cm-1, which is about half of the total conductivity ~7.67 x 10-3 S cm-1 obtained from the impedance data. In general, the calculated oxide ion conductivities are lower than the impedance data, which could be due to several factors. First, the difference between the calculated oxide ion conductivities and the impedance data might suggest that Li+ ion is another major charge carrier in the system. Further evidence such as the diffusion profiles from 6Li isotope exchange experiments is required to verify Li+ ion mobility at high temperatures. Moreover, due to Li doping, oxide-ion conduction may no longer occur through a vacancy-hopping mechanism and switch to an interstitialcy or cooperative type mechanism involving the concerted knock-on motion of interstitial and lattice oxide ions, leading to a lower oxygen diffusion rate. Last, the value of oxygen vacancy concentration used in the Nernst equation calculation might be different from the actual composition

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since the real composition of Li-doped NBT remains undetermined. Nevertheless, there is reasonable agreement between the calculated activation energy for the total conductivity determined from the IS data (Ea = 0.69 eV) and that of the oxygen diffusion coefficient (Ea = 0.53 eV) at temperatures above 450 oC, implying that at higher temperatures oxygen ion conduction dominates. In short, oxygen ions are the main charge carrier in the IT-SOFC operating temperature regime in conducting samples of Li-doped NBT, with additional charge carriers existing in the low temperature range.

Fig 6. (a) 18O diffusion profile of 4 at% Li-doped NBT after exchange at 600 oC for 2 hours and (b) Comparison of the total conductivity of 4 at% Li-doped ceramics obtained from diffusion data and AC impedance spectroscopy.

3.5 Degradation Test

In the degradation test, the changes in electrical performance were monitored for 100 hours at 600 oC in both flowing air and oxygen atmospheres. The grain boundary conductivity was observed to be relatively independent of the oxygen partial pressure, whereas the bulk conductivity suffered a significant reduction (Supporting Information Table S1). Overall, the total conductivity decreased by ~10% over the limited T-pO2

range investigated (Fig 7). SEM images of the cross-section and surface of 4 at% Li-doped NBT after annealing in the respective atmospheres are shown in Fig 8. The surface of 4 at% Li-NBT after annealing becomes porous, whereas the cross-section remains dense with only few pores. It is believed the porous surface is caused by the volatilization of elements (Li, Na and Bi) from the A-site during the testing process, initiating the degradation of the materials.

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Fig 7. Total conductivity versus time of 4 at% Li-doped NBT at 600 oC in different atmospheres.

Fig 8. SEM micrographs of Li-doped NBT samples after annealing at 600 oC in air and oxygen: cross-section (top) and surface (bottom)

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4. Conclusions

Partial substitution of Bi3+ with Li+ on the A-site of NBT improves the total conductivity by over one order of magnitude without changing the conduction mechanism. The enhancement of conductivity is attributed to an increase in the oxygen vacancy concentration and was confirmed by an increase in the oxygen self-diffusion coefficient. Among the doping levels, 4 at% Li was most effective in increasing electrical performance. Further enhancement may be restricted due to increasing association of the oxygen vacancies and dopant cations into complex defects of low mobility. Although first-principle calculations suggest that A-site acceptor doping is more promising, the results from this work do not support the superiority of A-site acceptor doping. This could be due to its unfavorable ionic size mismatch with the host Bi3+ [24] .

Compared to other oxide-ion conductors, Li-doped NBT shows competitive ionic conductivity above 600 oC, and initial results show that Li-doped NBT is reasonably stable under both air and oxygen atmospheres at 600 oC. Therefore, Li-doped NBT materials are promising for intermediate temperature SOFC applications. Further work would be needed to minimize the material loss from the A-site elements and to identify the other charge carriers in the system.

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