Local structure adaptations and oxide ionic conductivity in the Type III

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Local structure adaptations and oxide ionic conductivity in the Type III stability region of (1-x)Bi2O3·xNb2O5 Julia Wind, Neeraj Sharma, Aleksey A Yaremchenko, Vladislav V. Kharton, Douglas A Blom, Thomas Vogt, and Chris D. Ling Chem. Mater., Just Accepted Manuscript • DOI: 10.1021/acs.chemmater.8b00846 • Publication Date (Web): 02 May 2018 Downloaded from http://pubs.acs.org on May 3, 2018

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Chemistry of Materials

Local structure adaptations and oxide oxide ionic ionic conductivity in the Type III stastability region of (1– (1–x)Bi2O3∙xNb2O5 Julia Winda,†, Neeraj Sharmab, Aleksey A. Yaremchenkoc, Vladislav V. Khartonc.d, Douglas A. Blome, Thomas Vogtf, and Chris D. Linga,* a

School of Chemistry, The University of Sydney, Sydney 2006, Australia

b

School of Chemistry, The University of New South Wales, Sydney 2052, Australia

c

Department of Materials and Ceramic Engineering, CICECO – Aveiro Institute of Materials, University of Aveiro, 3810193 Aveiro, Portugal d

Institute of Solid State Physics RAS, Chernogolovka 142432, Moscow Distr., Russia

e

NanoCenter & Department of Chemical Engineering, University of South Carolina, Columbia, SC 29208, USA

f

NanoCenter & Department of Chemistry and Biochemistry, University of South Carolina, Columbia, SC 29208, USA

†

Current Address: Department of Chemistry, University of Oslo, Blindern, P.O. Box 1033, 0315 Oslo, Norway

ABSTRACT: Starting from a previously published stoichiometric model for the commensurate Type III phase in the (1– x)Bi2O3·xNb2O5 system, Bi94Nb32O221 (x = 0.254), we have developed a crystal-chemical model of this phase across its solidsolution range 0.20 ≤ x ≤ 0.26. After using annular dark-field scanning transmission electron microscopy to identify the metal sites that support non-stoichiometry, we show that the maximum possible range of that non-stoichiometry is 0.198 ≤ x ≤ 0.262, perfectly consistent with the experimental result. Inter-site cation defects on these sites provide some local coordinative flexibility with respect to the surrounding oxygen sublattice, but not enough to create continuous fluorite-like channels like those found in the hightemperature incommensurate Type II phase. This explains the reduced oxide-ionic conductivity of Type III compared to Type II at all temperatures and compositions, regardless of which phase is thermodynamically stable under those conditions. The solidsolution model shows that oxygen disorder and vacancies are both reduced as x increases, which also explains why Type III becomes relatively more stable, and why oxide ionic conductivity decreases, as x increases.

Introduction The cubic δ phase of Bi2O3 shows the highest reported oxide ion conductivities of any known solid-state material (~1 S cm–1 at 750°C), largely due to the presence of 25% disordered vacancies in its fluorite-type oxygen sublattice.14 δ-Bi2O3 is only stable at temperatures between 750 and 825°C, which rules out its practical use as an electrolyte in solid-oxide fuel cells (SOFCs). However, doping with small concentrations of d0 transition metal cations such as V5+/Nb5+/Ta5+, Cr6+/Mo6+/W6+, or Tc7+/Re7+ has been shown to stabilize the average fluorite-type structure of δBi2O3 to room temperature.5-14 While this reduces ionic conductivity by reducing the total number of oxygen va-

cancies and increasing the ordering of the remaining oxygen sublattice, it serves to increase temperature stability and eliminates major structural transformations. Particularly when co-doped with rare-earth elements, such stabilized δ-Bi2O3 phases are of genuine interest for practical SOFC applications.15 The present study concerns itself with the Nb5+ doped system, i.e., (1-x)Bi2O3·xNb2O5. This system contains a variety of fluorite-type δ-Bi2O3-related superstructures depending on x, the relative conductivities of which can be rationalized in terms of the type of superstructure formed. The most intensively studied phase in the (1x)Bi2O3·xNb2O5 system is the solid-solution denoted Type

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II by Zhou et al.16 (0.062 ≤ x ≤ 0.25). Type II was first recognized by Miida and Tanaka17 as having a (3+3)-D incommensurately modulated structure which was further characterized by us8, 18 and, more recently, solved and refined against single crystal neutron diffraction data19 in the cubic superspace group Fm3m(α,α,α)q00(α,–α,–α)q00(– α,α,–α)000 (number 225.3.215.8 using the notation of Stokes, Campbell & van Smaalen20). Its complex but remarkably high-symmetry structure features displacive and compositional modulation of both anions and cations. The compositional modulation arises due to the ordering of Nb5+/Bi3+ and O2-/vacancies while the displacive modulation can be considered as a Coulombic consequence of the compositional modulation: O2- reacts to cationic charge by moving from its average fluorite-type position towards Nb5+, and away from Bi3+. At the Nb-rich end of the Type II solid solution, an intermediate phase referred to as Type III is found at lower temperatures.8, 16, 21-22 Type III has a commensurately modulated 3×3×7 superstructure of the δBi2O3 subcell with tetragonal I4m2 space-group symmetry, and lower oxide ionic conductivity.23 The published Type III model provides a reasonable fit to high-resolution synchrotron X-ray and neutron powder diffraction data; however, its fixed composition of Bi94Nb32O221 (x = 0.254) contradicts clear experimental evidence for a solid-solution range of 0.20 ≤ x ≤ 0.26.21 The exact range of this flexibility, and the crystal-chemical mechanisms that permit it, remain unclear and are addressed in this work. From a crystal-chemical perspective, both Type II and Type III can be understood in terms of the ordering of Nb dopant cations along F directions of the fluorite-type Bi sublattice, to form corner-connected chains of NbO6 octahedra. Where these chains meet, they form tetrahedral clusters of NbO6 octahedra reminiscent of the basic structural unit of pyrochlore-type, such that the Type II solidsolution can be thought of as representing a smooth transition from fluorite-type δ-Bi2O3 to a (hypothetical) pyrochlore-type Bi2Nb2O7 phase. In Type II, these ordered NbO6 chains delineate continuous, relatively disordered and oxygen-deficient δ-Bi2O3-like “channels” whose width depends on composition (Figure 1a), explaining its high thermal stability and oxide ionic conductivity (~0.1 S cm–1 at 1000°C).19 In contrast, Type III features non-continuous NbO6 chains organized into layers or isolated groupings (Figure 1b).18, 23 In Type III, oxide vacancies are all ordered in proximity to Nb cations, explaining its relatively lower oxide ionic conductivity at the same temperatures.19

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Figure 1: (a) Commensurate approximate (8×8×8 fluorite-type sub-unit cells) of the incommensurate Type II structure,19 and (b) the Type III structure23 in the (1–x)Bi2O3·xNb2O5 system, x ≈ 0.25. Bi atoms are green, NbOy polyhedra are blue, and O atoms are red. The Bi8 and Nb4 sites in the Type III structure (discussed in detail in the main text) are highlighted in slightly darker and lighter shades respectively.

Over the compositional range in which they coexist, the fact that the high-temperature Type II and low-temperature Type III phases have different Bi/Nb distributions as well as different O/vacancy distributions means that Type II can be easily quenched to low temperatures. Quenched metastable Type II samples will be referred to in the present work as Type II*. While oxide vacancy concentrations remain the same for Types II*, II and III of a given Bi/Nb ratio,24 the reconstructive Type II* → III → II phase transitions on reheating these quenched samples have obviously detrimental implications for their use in SOFCs. A number of kinetics-based studies22, 25 using conductivity and X-ray powder diffraction (XRD) of the Type II → Type III transition have set out to probe its structural subtleties. Type III phase was found to have a smaller unit cell volume compared to the Type II, and the Type III unit cell volume decreases as a function of x, implying that shrinkage occurs during conductivity measurements. Despite this shrinkage, it has been shown that Type III has higher electronic (σe-) and ionic (σion) conductivity than Type II at temperatures below 600°C for x = 0.25. Interestingly, some measurements suggest that the ionic and total conductivity converge near the temperature limits; this would indicate a completely ionically conduct-


ing material.22 The activation energies are 1.14 and 0.82 eV for Type II and Type III respectively, implying a lower energy barrier for conductivity in Type III at these temperatures. In both cases, the p-type σe- is shown to be the minor contributor to the total conductivity. However, when x = 0.25 at 820°C, σe- of the Type II phase increases by a factor of 3–4, consistent with a transformation from Type II to Type III. Related work on the (1-x)Bi2O3·xNb2O5 series where x = 0.05, 0.08, 0.10 and 0.1626 reports a gradual decrease in oxide ionic conductivity as a function of temperature as the Nb content is increased, but limited details concerning the structure were presented.

(a) Type III (150°C)

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(b)

In the study reported here, we set out to definitively establish the range of the Type III solid-solution by consolidating previous experimental reports and new data, and to develop a coherent crystal-chemical model explaining that range. We then used this crystal-chemical model to correlate changing structural features with physical properties, specifically, high-temperature oxide-ionic conductivity.

Experimental

Intensity (arb. units)

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Intensity (arb. units)

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Type II (923°C)

0

Polycrystalline samples of composition x = 0.20, 0.22, 0.24 and 0.26 in the series (1–x)Bi2O3·xNb2O5 were synthesized by ball-milling stoichiometric quantities of Bi2O3 (Sigma Aldrich 5N) and Nb2O5 (Sigma Aldrich 4N5) for 24 h. The samples were then placed in closed Pt crucibles and heated in cycles ranging from 700–880°C for 5–24 h. Phase purity at room temperature (Type II* or Type III) was confirmed by X-ray powder diffraction (XRD) using a Panalytical X'Pert Pro diffractometer with a Cu Kα source and by neutron powder diffraction (NPD) using the instrument D20 at the Institut Laue-Langevin (ILL), Grenoble, France,27 with a take-off angle of 120° and neutron wavelength λ = 1.88 Å. Figure 2 shows a Rietveld fit to the NPD data for the Type III phase at x = 0.26, using the structural model of Ling and Johnson.23

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Figure 2: Rietveld fit to λ = 1.88 Å NPD data for the Type III phase of (1-x)Bi2O3∙xNb2O5 where x = 0.26, using the structural model of Ling and Johnson.23 Red crosses are experimental data, black line is calculated fit, purple line below is the difference. The cell is tetragonal I-4m2, a = 11.5732(3), c = 38.6903(15) Å. Refinement statistics: Rp = 3.77%, wRp = 4.84%, χ2 = 1.37 for 30 parameters.

Aberration-corrected scanning transmission electron microscopy (AC-STEM) data in the Z-contrast high-angle annular dark-field (HAADF) imaging modes were collected from powdered samples. Dry powders were lightly ground in an agate mortar and pestle before loading onto holey carbon-coated Cu grids. A JEOL JEM2100F with a CEOS aberration corrector for the electron probe was used at 200 kV to image the sample with a 24 mrad convergence angle. Images were acquired using a Fischione Model 3000 detector with a camera length such that the detector spanned 75–178 mrad. The scanning acquisition was synchronized to the 60 Hz AC electrical power to minimize 60 Hz noise in the images, and a pixel dwell time of 15.8 µs was used. The total conductivity (σ) in air was studied by AC impedance spectroscopy (HP4284A precision LCR meter, 20 Hz – 1 MHz, Agilent Technologies) using dense bar-shaped samples with porous Pt electrodes. Dilatometric studies were performed in air using an alumina Linseis L70/2001 dilatometer at a constant heating rate of 3 K min-1.



Results and Discussion Relating the Type III solid-solution range to its crystal structure requires first of all knowledge of Bi/Nb mixed site occupancies. Unfortunately, the complexity and pseudosymmetry of the Type III structure (50 crystallographically independent atoms in the unit cell of the highest-symmetry ordered model23) makes it impossible to unambiguously determine them by Rietveld-refinement against XRD or even neutron powder diffraction data. We therefore used AC-HAADF-STEM as the primary structural characterization technique in this study.

(a)

(b)

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Figure 3a shows a = F zone axis HAADF STEM image of Type III (1–x)Bi2O3·xNb2O5, x = 0.26. Figure 3b is a portion of the micrograph in Figure 3a along with a structural model projection of the cations from the published structure.23 The labels refer to the crystallographic atomic sites in the published model. In the zone axis projection of the Type III model structure, there are four classes of atomic cation columns, fully Bi, fully Nb, 1:2 Nb:Bi and 2:1 Bi:Nb. We therefore expect to see differences in the HAADF signal due to the composition differences of these columns. Bi7/Bi8

Bi3/Bi7 Bi5/Bi9

Nb3/Bi4

Nb1/Bi10 Nb4/Bi1

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Bi1/Bi2

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Bi2/Bi6

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Bi4/Bi10 Nb5/Bi5

Bi3/Bi7

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Bi2/Bi6

Bi4/Bi10

1 nm

Bi3/Bi7 Nb2/Nb5

Nb1/Nb3 Nb2/Bi9

Bi3/Bi7

Bi7/Bi8

Figure 3: (a) HAADF-STEM images along the (a) = F zone axis of Type III (1–x)Bi2O3·xNb2O5, x = 0.26. (b) A portion of this micrograph along with a projection of the cations in the same direction, from the published structure,23 with labels referring to the crystallographic sites in that model.

It is evident from the micrograph that there are intensity variations for different columns of atoms in the unit cell. To quantify the HAADF signals for the different types of atomic columns in the zone axis image, we used the integrated signal over the projected columns which has been shown to be relatively insensitive to small focus variations.28-29 To account for the experimentally observed thickness gradient in the micrograph, we subtracted a background signal selected from the nearby electron probe positions not associated with the atomic columns.30 For each type of atomic column in the projected unit cell, an area with a radius of 7.5 pixels (131.26 pm) was selected and the signal integrated. Figure 4 shows the mean HAADF background-subtracted integrated signal for each of the 14 different atomic columns in this projection of the unit cell. Error bars are 95% confidence limits on the mean value. There are broadly 4 different signal levels observed in Figure 4. The brightest atomic columns are not surprisingly columns consisting of only Bi containing sites; Bi3/Bi7, Bi5/Bi9, Bi1/Bi2 and Bi4/Bi10. The least intense atomic columns correspond to

the two examples of atomic columns in the projection that are fully Nb; Nb2/Nb5 and Nb1/Nb3.

Figure 4: Mean HAADF background-subtracted integrated signal for each of the 14 unique atomic columns in the = F projection of the unit cell of Type III (1–x)Bi2O3·xNb2O5, x = 0.26.. Error bars are 95% confidence limits on the mean value.


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Because the HAADF STEM image was not collected on an absolute scale, for convenience Table 1 compares Zcontrast ADF intensities in the = F image (normalized by the mean integrated intensity observed for the most intense atomic column, Bi3/Bi7) to the expected values for the fully ordered model for Type III Bi94Nb32O221 (x = 0.254).23 A number of discrepancies are apparent. Firstly, we note that in the crystallographic model derived from powder diffraction four columns are significantly offset in projection: Nb2/Nb5 (by 32pm) Bi2/Bi6 (34 pm), Nb4/Bi6 (25 pm), and Nb5/Bi5 (29 pm). This should make them appear dimmer than expected from a purely Z-contrast perspective due to electron channeling (a non-centrosymmetric intensity distribution along the channel direction).31 The Nb2/Nb5 column has an intensity consistent with the other fully Nb column, Nb1/Nb3, which itself has an offset in projection of 10.5 pm. The Bi2/Bi6 columns are in fact dimmer than most of the other fully Bi atomic columns. Mixed Nb-Bi column Nb4/Bi6 is also dimmer than the majority of the other mixed Nb-Bi columns, although the model structure implies that its composition is more Nb-rich than the other mixed columns and the atoms are offset in this projection, both factors which would tend to suppress the HAADF signal. Finally, the Nb5/Bi5 column exhibits a HAADF signal which is dimmer than other columns with the same composition according to the model structure. Clearly, electron channeling in this orientation plays a role in the observed HAADF intensity. The data at this point are not sufficient to definitively decouple the effects of atomic column composition and channeling for atomic columns which have significant offset. Table 1: Atom columns in zone axis projection to the fully ordered model for Type III Bi94Nb32O221 (x = 0.254)23, fractional Nb composition, normalized ADF intensities from Figure 2(a), offsets of the atomic sites in the model perpendicular to , and agreement between measured HAADF intensities and the model composition. Atomic Column

Model Frac(Nb)

Bi7/Bi8 Bi3/Bi7 Bi5/Bi9 Nb2/Nb5 Nb3/Bi4 Nb1/Bi10 Nb4/Bi1 Bi2/Bi6 Nb4/Bi6 Bi1/Bi2 Bi4/Bi10 Nb1/Nb3 Nb5/Bi5

0 0 0 1 1/3 1/3 1/3 0 2/3 0 0 1 1/3

Normalized Integrated Intensity 0.79±0.03 1.00±0.02 0.91±0.03 0.45±0.02 0.78±0.02 0.82±0.02 0.83±0.02 0.78±0.03 0.62±0.02 0.92±0.02 0.96±0.02 0.43±0.02 0.63±0.02

a offset (pm) 0 -2.65 1.27 0 0 -2.53 5.07 0 0 9.10 -8.30 0 0

c offset (pm) 2.66 2.66 4.13 32.35 -4.97 8.56 0.42 33.55 25.06 8.06 6.90 10.49 28.53

Agreement with model NO Yes Yes Yes Yes Yes Yes NO Yes Yes Yes Yes NO

Nb2/Bi9

1/3

0.80±0.02

-1.73

0.31

Yes

We then note that the Bi7/Bi8 column (expected to be 100% Bi), which is not offset, is also too dim. Given that the Bi3/Bi7 and Bi5/Bi9 columns which have similar amounts of offset to the Bi7/Bi8 column are both equally brighter than the Bi7/Bi8 column, this suggests that the Bi8 site has some mixed occupancy by Nb. If Bi8 was 100% Bi (fully ordered model), the cation stoichiometry would be Bi94Nb32 (x = 0.254); whereas if Bi8 were 100% Nb it would be Bi92Nb34 (x = 0.270). This is inconsistent with the experimental Type III solid-solution range 0.20 ≤ x ≤ 0.26, indicating that some Nb sites from the ordered model must simultaneously support Bi occupancy. Noting that Nb2/Nb5 and Nb1/Nb3 are by far the dimmest columns in Table 1, Nb4 is the site most likely to support additional Bi occupancy. If Nb4 was 100% Bi, the Bi-rich end of the solid solution becomes Bi102Nb24 (x = 0.190). AC-HAADF-STEM results concerning cation occupancy are therefore consistent with a solid-solution range 0.190 ≤ x ≤ 0.270. The fact that the experimental range is more restricted than this must then be due to the oxygen sublattice. In considering this, we note that every Nb5+ which replaces a Bi3+ increases the total O stoichiometry. The fully ordered model for Bi94Nb32O221 includes one O vacancy per formula unit to meet this requirement. However, every Nb5+ also increases the number of O vacancies required to achieve local NbO6/NbO4 coordination rather than the average 8-fold cubic coordination of Bi in fluoritetype δ-Bi2O3. Bi8 has cubic coordination and its symmetry is consistent with tetrahedral coordination, so converting a Bi atom on this site into Nb increases the required number of O vacancies by 8 per formula unit, while also increasing the O stoichiometry by 2. This could only be accommodated up to the point of filling the one oxygen vacancy per formula unit in the fully ordered model of Bi94Nb32O221, i.e., Bi93Nb33O222 or x = 0.262. Higher Nb content on the Bi8 site would require losing Nb from another site (i.e., Nb4). Converting a Nb atom on the Nb4 site into a Bi atom decreases the required number of O vacancies by 2.5 per formula unit, while also decreasing the O stoichiometry by 2. This gives a lower end to the solid-solution range at Bi102Nb24O213 (x = 0.190). However, the Bi atoms replacing 4-coordinate Nb would be significantly under-bonded, especially with the concomitant reduction in O stoichiometry. Placing Nb on the Bi8 site would increase the O stoichiometry and relieve this under-bonding, up to Bi101Nb25O214 or x = 0.198.


Chemistry of Materials This gives us a crystal-chemical model of Bi substitution on the Nb4 site, partially compensated for by a smaller degree of Nb substitution on the Bi8. This model predicts a solid solution range 0.198 ≤ x ≤ 0.262, perfectly consistent with the experimental range of 0.20 ≤ x ≤ 0.26.21 It also predicts the increased stability of Type III relative to Type II at higher x, observed experimentally:21 increasing x leads to replacement of the last remaining coordinatively flexible (partially disordered) Bi(III) site, Bi8, with coordinatively rigid Nb(V); while also filling in the last remaining oxygen vacancies. Type III is thus most ordered, and therefore most enthalpically favoured, at higher x. Another significant feature of this model is that the compositional flexibility of the Type III phase is limited to intersite cation defects between only two sites, Bi8 and Nb4. These sites are highlighted in Figure 1b. From inspection of the structure, it is clear that while these cation defects will provide some local coordinative flexibility in the surrounding oxygen sublattice, they will not create continuous fluorite-like channels comparable to those found in the Type II phase.19 The model therefore predicts reduced oxide-ionic conductivity of Type III compared to Type II at all temperatures and compositions, regardless of which phase is more thermodynamically stable under those conditions.

getic barriers imposed by increasing Nb content and decreasing O vacancies. Castro et al.32 report a lower activation energy of 85 kJ mol-1 for x = 0.25, but Takahasi and Iwahara3 report activation energies of 115 kJ mol-1 and 135 kJ mol-1 for x = 0.22 and x = 0.30 respectively. Our observations follow the same trend as the latter measurements, namely, increasing Nb content slightly increases activation energies. The result of Castro et al.32 appears to be an outlier, possibly related to differences in sample preparation and sample history, noting that the compositions in their study (like ours) involve two phases with significantly different conductivities (see below). At lower temperatures, a difference in conductivity between the Type II and Type III phases of composition x = 0.25 was also observed by Wang et al.22 0 (Bi2O3)1-x(Nb2O5)x -1

log σ (S×cm-1)

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-2 x = 0.20 x = 0.22 x = 0.24 x = 0.26

-3

This is consistent with reduced oxide-ionic conductivity of Type III compared to Type II at all temperatures and compositions, regardless of which phase is more thermodynamically stable under those conditions. Conductivity measurements taken on cooling Type II samples at a range of compositions are shown in Figure 5. Table 2 presents densities and activation energies for conduction. As the Nb content is increased, the oxide ion conductivity is reduced and the activation energy exhibits a slight increase. Thus, increasing Nb content requires more energy for conduction. Two activation energy regions were found for the more Nb-rich samples. Higher temperature regions were characterized by higher activation energies. Conductivity measurements follow the trends described by Yaremchenko et al.26 for x = 0.05–0.16 (wholly Type II region in the phase diagram, see Figure 9 of Pirnat et al.21); i.e., as x increases (Nb content increases), conductivity decreases. The activation energies for samples 0.20 ≤ x ≤ 0.26 range from 118–157 kJ mol-1 (Table 2), higher than observed for 0.05 ≤ x ≤ 0.16 for which the range was 58– 121 kJ mol-1.26 This can be ascribed to the change in Nb and O vacancy content, whereby increasing Nb content reduces the O vacancies by charge compensation. More energy is required to drive conduction and overcome ener-

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Figure 5: Temperature dependence of the total conductivity of (1– x)Bi2O3·xNb2O5 samples in air.

Table 2: Activation energy for total conductivity of Type II (1– x)Bi2O3·xNb2O5 samples. Density Composition T (K) EA (kJ mol-1) (g cm-3) x = 0.20 x = 0.22 x = 0.24

7.25 7.12 7.31

x = 0.26

8.31

910–1220 910–1220 1110–1250 910–1110 1090–1220 910–1090

118 ± 1 127 ± 1 155 ± 1 129 ± 1 157 ± 2 130 ± 4

To determine the differences in conductivity between the Type II/II* and Type III phases upon heating, AC impedance measurements were taken on heating an x = 0.24 Type II* sample to 860°C. The sample was held at this temperature for 12 h to allow the Type II* → III transition to occur,


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then taken to 980°C, followed by cooling. Figure 6 shows the results, highlighting a reduction in conductivity in the Type III region between 860 and 940°C. These show that on cooling, sufficient time was not allowed for the transition back to Type III.

-0.8

log σ (S×cm-1)

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(Bi2O3)0.76(Nb2O5)0.24

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Type III phase

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Figure 6: Temperature dependence of the total conductivity of x = 0.24 in air, measured in heating and cooling regimes. The sample was equilibrated at each temperature for 0.5 h and at the Type II → III transition temperature for 12 h.

These observations were verified by an ex situ series of experiments with pellets sintered at appropriate temperatures (Figure 7). The Type II* → Type III transition is demonstrated by the splitting of the cubic Type II* peak and its re-formation, showing the presence of the Type III phase between 860 and 940°C. The Type III phase shows lower conductivity relative to the Type II phase, although their activation energies are similar ~155 ± 5 kJ mol-1 (see Table 2).

After conductivity measurement 945°C for 24 h 945°C for 1 h 900°C for 1 h 865°C for 12 h 820°C for 12 h Type II at 1000°C

55

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Type III

Figure 7: Portions of ex situ XRD patterns of an x = 0.24 sample subjected to the same heating cycles as for the conductivity measurements. The grey pattern is from the sample after conductivity measurements. At 820°C, small signs of the Type III phase are seen but the sample is predominantly Type II. At 865°C the sample is all Type III and at 945°C it transforms back to Type II.

There is a noticeable reduction in conductivity on conversion from Type II* to Type III at high temperatures (Figure 6), reinforcing the notion of a reconstructive phase transition. In the high temperature regime, activation energies for Type II and Type III are virtually identical, ~ 155 kJ mol-1. If progressive disordering on heating was occurring, we would see a change in the activation energy; the failure to observe a change implies that the Type II* → III transition leads to blocking of mobile ions or vacancies. This is consistent with the presence of continuous, coordinatively flexible bismuth-rich fluorite-like channels along F directions in the Type II structure, passing between ordered NbO6 octahedral chains.19 There are no such continuous channels in Type III. Wang et al.22, 25 conflict with our results in that they report Type III as having higher conductivity than Type II. However, closer examination of their observations (Figure 12 of Wang et al.22) show the conductivity of both the Type II and Type III phases converging at approximately 830°C, a trend we also observe (Figure 6). We measure total conductivity at higher temperatures for the Type III and Type II phases, while Wang et al. use the Wagner polarization method to extract electrical conductivity from the total conductivity measurements and find that at temperatures below 830°C this component is higher in Type III than Type II. The activation energy we obtain for the Type II phases of composition x = 0.24 and 0.26 below 830°C is around 130 kJ mol-1, close to the previously reported value of 110 kJ mol-1 for x = 0.25 in this temperature region by Wang et al.22 Interestingly, their activation energy for the Type III phase is smaller, ~80 kJ mol-1, similar to Castro et al.’s32 findings, suggesting that Castro et al. measured a Type III rather than a Type II phase. Thus, we conclude that at lower temperatures (< 830°C) the activation energies of the Type III phase are lower (80–85 kJ mol-1) than for Type II (110–130 kJ mol-1). A final consideration in our experiment was the potentially small volume changes of the well-sintered pellets with temperature. Dilatometric measurements on the Type II* x = 0.24 sample showed some volume contraction at temperatures above 800°C (Figure 8), which suggests either further densification of the pelletized sample or a volume contraction due to the Type II* → III transition. However,


Chemistry of Materials the ex situ XRD patterns in Figure 7 show that the phases present under the same heating regime are Type II and Type III.

1.0 0.5

∆L/L0, %

1 2 3 4 5 6 7 8 9 10 11 12 13 14 15 16 17 18 19 20 21 22 23 24 25 26 27 28 29 30 31 32 33 34 35 36 37 38 39 40 41 42 43 44 45 46 47 48 49 50 51 52 53 54 55 56 57 58 59 60

0.0 -0.5 -1.0 -1.5

(Bi2O3)0.76(Nb2O5)0.24

0

200

400

600

800

1000

T, °C Figure 8: Dilatometric measurements of the x = 0.24 Type II* sample taken on heating and cooling.

Sample preparation and densification are important issues in direct conductivity measurements, thus it was important to undertake these dilatometry and ex situ XRD measurements to verify our observations in conductivity. The minimum density of the samples was 84% relative to the theoretical density of x = 0.25 compound of 8.48 g cm-3.23 If densities are not sufficiently high, the conductivity is a measure of pore and sample conductivity. The results of previous work measuring conductivities of the Type II/III phases may need to be critically analyzed in light of such density measurements.

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Bi94Nb32O221 (x = 0.254).23 Increasing x involves replacing the coordinatively flexible (partially disordered) Bi3+ site, Bi8, with fully ordered Nb5+; while also filling in the last remaining vacancies in the fluorite-type oxygen sublattice, due to the increased average cation valence. Type III is thus most ordered, and therefore most enthalpically favoured, at higher x. The model predicts a solid-solution range 0.198 ≤ x ≤ 0.262, in perfect agreement with the previously determined experimental value 0.20 ≤ x ≤ 0.26.21 We then conducted unique high-temperature measurements to track ionic conductivity through the Type II* → Type III → Type II transitions, confirming a relative drop in total conductivity in the Type III regime at all compositions. This contradicts some previous reports.22, 25 Even though the Type III phase has higher electronic conductivity22, its total conduction is below that of the (predominantly ionic conducting) Type II phase. All of these factors coincide to have detrimental effects on oxide-ionic conductivity as x increases. Not only does the Type III phase become less conductive due to reduced coordinative flexibility and oxygen vacancies, but its residual phase fraction should increase relative to the more conductive Type II phase (noting that the first-order Type II* → Type III → Type II transitions have a long induction time22), which we confirmed by ex situ XRD measurements of our own samples. There are no competing factors at play in this case, suggesting that further attempts to optimise the performance of Nb-stabilized δ-Bi2O3 (and, most likely, other transition metal-doped versions) should unambiguously focus on the Bi-rich end of the Type II region.

Finally, we note in Figure 8 that a small positive change in volume appears to occur above 600°C. This could indicate a small degree of reduction, given the very high oxide ion diffusion of these materials, even in such a high density sample.

AUTHOR INFORMATION

Conclusions

The manuscript was written through contributions of all authors. All authors have given approval to the final version of the manuscript.

HAADF-STEM of samples in the (1–x)Bi2O3·xNb2O5 system shows that the low-temperature commensurate Type III phase becomes more well-ordered, and more stable relative to the high-temperature incommensurate Type II phase, as x increases. This is consistent with previous phase diagram studies based on X-ray and neutron powder diffraction.8,21 Quantitative examination of Z-contrast in the images leads us to a crystal-chemical model that explains the reason for this. It shows that the compositional flexibility of the Type III phase is limited to inter-site cation defects between only two sites – Bi8 and Nb4 – in the near-stoichiometric model

Corresponding Authors * Chris D. Ling, Email address: [email protected]

Author Contributions

Funding Sources This work was supported by the Australian Research Council (Discovery Project scheme). J.W. acknowledges financial support from the Australian Institute of Nuclear Science and Engineering (Postgraduate Research Awards scheme). V.Kh. acknowledges financial support from the Russian Science Foundation (project 17-79-30071) and Russian Ministry of Education and Science (project 14.B25.31.0018). A.Y. acknowledges financial support


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from the FCT, Portugal (projects IF/01072/2013/CP1162/CT0001 and CICECO - Aveiro Institute of Materials POCI-01-0145-FEDER-007679 (FCT ref. UID/CTM/50011/2013)). D.B. and T.V. thank the VRR of the University of South Carolina for support operating the JEOL 2100F.

ABBREVIATIONS AC-STEM, aberration-corrected scanning scanning transmission electron microscopy; HAADF, high angle annular dark-field; NPD, neutron powder diffraction; XRD, X-ray powder diffraction

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30. Klenov, D. O.; Stemmer, S., Contributions to the contrast in experimental high-angle annular dark-field images. Ultramicroscopy 2006, 106, 889-901. 31. Voyles, P.; Muller, D.; Kirkland, E., DepthDependent Imaging of Individual Dopant Atoms in Silicon. Microsc. Microanal. 2004, 10, 291-300. 32. Castro, A.; Aguado, E.; Rojo, J. M.; Herrero, P.; Enjalbert, R.; Galy, J., The New Oxygen-Deficient Fluorite Bi3NbO7: Synthesis, Electrical Behavior and Structural Approach. Mater. Res. Bull. 1998, 33, 31-41.


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TOC Figure: High-resolution electron microscopy image and crystal structure of Type III (1–x)Bi2O3·xNb2O5.

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Local structure adaptations and oxide ionic conductivity in the Type III

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