Stabilizing coexisting n-type electronic and oxide-ion conductivities in

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Stabilizing coexisting n-type electronic and oxide-ion conductivities in donor-doped Ba-In-based oxides under oxidizing conditions - Roles of oxygen disorder and electronic structure Yukio Cho, Masayuki Ogawa, Itaru Oikawa, Harry L. Tuller, and Hitoshi Takamura Chem. Mater., Just Accepted Manuscript • DOI: 10.1021/acs.chemmater.8b03818 • Publication Date (Web): 27 Mar 2019 Downloaded from http://pubs.acs.org on March 28, 2019

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Chemistry of Materials

Stabilizing coexisting n-type electronic and oxide-ion conductivities in donor-doped Ba-In-based oxides under oxidizing conditions - Roles of oxygen disorder and electronic structure Yukio Cho,† Masayuki Ogawa,† Itaru Oikawa,† Harry L. Tuller,∗,‡ and Hitoshi Takamura∗,† †Department of Materials Science, Graduate School of Engineering, Tohoku University, Sendai 980-8579, Japan ‡Department of Materials Science and Engineering, Massachusetts Institute of Technology, Cambridge, MA 20139, USA E-mail: [email protected]; [email protected] Abstract Surface oxygen reduction reaction (ORR) rates at n-type oxide-based Mixed IonicElectronic Conducting (MIEC) Solid Oxide Fuel Cell (SOFC) cathodes can be expected to be enhanced relative to that at p-type MIEC cathodes due to the greater availability of electrons at higher energies in the band structure needed for the charge transfer reaction. However, given the difficulty of achieving coexisting oxygen vacancies and electrons in conduction band under oxidizing cathode conditions, no stable n-type MIEC cathodes have been reported to date. In this study, a predominantly n-type MIEC

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conductivity is confirmed in a Ba-In-based oxide (BNIM) co-doped with Nd and Mn at high temperature and high PO2 as confirmed by the PO2 dependence of the electrical conductivity and negative Seebeck coefficients, combined with readily measurable oxide-ion transference numbers. This coexistence of n-type electronic and oxide-ion conductivities is discussed based on the electrical behavior of BNIM with different Mn levels, and is attributed to the significant change in the degree of anion-Frenkel ordering and the band structure associated with heavy donor doping of Ba2 In2 O5 . This novel n-type MIEC has the potential for enhancing the ORR at SOFC cathodes at reduced temperatures and thereby identifying new potential candidate cathode materials for next-generation SOFCs.

Introduction During the last several decades, the solid oxide fuel cell (SOFC) has received much attention as a candidate for next-generation alternative energy conversion devices due to its high conversion efficiency, fuel flexibility, and potentially emissions-free operation when hydrogen is the fuel. 1 Most SOFCs require a high operating temperature to ensure a high electrolyte oxide ionic conductivity. However, high operating temperatures also limit the selectivity of materials and result in the accelerated degradation of SOFC components. 1,2 Reducing the operating temperature of SOFCs while maintaining their high conversion efficiencies is therefore a key objective of recent research. 3,4 Unlike conventional electronically conducting electrodes (e.g. Pt or La1−x Srx MnO3−δ ), MIECs exhibit both ionic and electronic conductivity that contribute to improved cathode performance at intermediate temperatures. They do so by enlarging the reaction sites, which are limited to three phase boundaries for strictly electronic conductors, to the entire cathode surface for MIECs. 5 Nevertheless, elevated cathodic polarization resistances remain limiting at intermediate temperatures. 6 Recently, a study of SrTi1−x Fex O3−δ demonstrated that the surface exchange rate in this system is governed by the concentration of electrons in the conduction band rather than 2


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either the high levels of its majority p-type electronic and/or ionic conductivity. 7 This finding was explained by the authors as consistent with the limited exchange rate due to the charge transfer reaction associated with the Oxygen Reduction Reaction (ORR), as summarized in Fig. 1 and in Eq. (1) below: ′

q − − Oq− x,ad + e = Ox,ad

(1)

where Oq− x,ad is the energy state of oxygen intermediates adsorbed at the surface. This reaction benefits from having electrons at higher energy states to drive the reduction reaction to the right.

gas

(a)

electron Oq- x, ad

EC EG EA EF EV hole

(b)

electron

gas Oq- x, ad

EC EF ED EG EV

Figure 1: Schematic diagram of electronic structure for (a) p-type MIEC and (b) n-type MIEC.

Whereas the electrons in the conduction band are minority carriers in p-type MIECs such as La0.4 Sr0.6 Co0.2 Fe0.8 O3−δ (LSCF) 8 and Ba0.5 Sr0.5 Co0.8 Fe0.2 O3−δ (BSCF), 9 the concentration 3


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of electrons decreases exponentially with decreasing temperature and thereby could limit the electrode reaction rate (Fig. 1a). 10–12 On the other hand, an n-type MIEC would have a much higher concentration of electrons in the conduction band as the majority carrier and this, in turn, is expected to enhance the charge transfer reaction even in the intermediate temperature regime if the surface exchange mechanism remains the same for the n-type MIEC (Fig. 1b). Mixed n-type electronic and oxide ionic conductivities are, however, normally only observed under reducing conditions as, for example, in acceptor-doped CeO2 13,14 or donor doped SrFeO3−δ , 15 while the coexistence of both n-type conductivity and oxide ionic conductivity under oxidizing atmospheres is considered unlikely because the reduction reaction (see Eq. (2)) is driven to the right under oxidizing conditions. 1 •• VO + 2e− + O2 → O× O 2

(2)

Although a stable n-type MIEC in the highly-oxidized SOFC operating environment has not yet been reported, such an MIEC would show promise as a cathode material for SOFCs at intermediate temperatures as discussed above. In this work, we describe the approach taken to prepare an n-type MIEC by donor doping Ba2 In2 O5 . The brownmillerite structured Ba2 In2 O5 has already attracted considerable interest as a solid oxide electrolyte candidate. 16 The brownmillerite structure (A2 B2 O5 ) can be viewed as a highly oxygen-deficient perovskite compound (1/6 oxygen sites empty relatives to perovskite structure) with oxygen vacancies ordered in the [101] direction as shown in Fig. 2.

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order oxygen defects [101]

Ba2+ In3+

Ba2In2O5 brownmillerite structure

O2-

Figure 2: Crystal structure of brownmillerite structure Ba2 In2 O5 with [101] order oxygen defects.

Whereas brownmillerite Ba2 In2 O5 exhibits relatively low oxide ion conductivity because of the ordered oxygen defects at reduced temperatures, these oxygen defects become disordered, dramatically increasing oxide ionic conductivity above the phase transformation temperature to the perovskite structure that occurs at above 930◦ C. 17 Many studies have focused on decreasing the order-disorder phase transformation temperature of Ba2 In2 O5 by doping with aliovalent elements in order to maintain the high oxide ion conductivity at lower temperatures. Kakinuma et al. succeeded in stabilizing the disordered oxygen-deficient perovskite structure at room temperature by doping with La by substitution on the Ba site in Ba2 In2 O5 . 18 Moreover, much higher oxide ionic conductivity (σi = 0.12 S/cm−1 at 800◦ C) 19 and high power density with Pt and Ni electrode (0.51 W cm−1 at 800◦ C) 20 were obtained with (Ba0.3 Sr0.2 La0.5 )2 In2 O5.5 as the electrolyte in an SOFC by enhancing the concentration of mobile oxide ions and also optimizing the unit cell free volume by substituting Ba2+ with Sr2+ . The unique point for this materials is that Ba2 In2 O5 achieves high oxide ionic conductivity, contrary to normal expectations, e.g., in (La1−x Srx )(Ga1−y Mgy )O3−δ (LSGM), 21 by 5


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reducing the total number of oxygen vacancies in the material. This follows from the fact that because the undoped material has a very high oxygen vacancy concentration, they tend to order and thereby become more localized. Although the predominant electronic carriers in Ba2 In2 O5 have been identified as the holes, 17 the appropriate selection of donor elements and dopant levels is expected to provide Ba2 In2 O5 with the properties of a mixed n-type electronic and oxide ion conductor. In this research, a two-pronged procedure was adopted for the preparation of the novel n-type MIEC. First, we investigated the use of dopants with smaller radii and higher charge than Ba2+ to maintain the disordered structure and high free volume as in (Ba0.3 Sr0.2 La0.5 )2 In2 O5.5 . Then we added transition metal elements to create higher impurity energy levels within the band gap of Ba2 In2 O5 , and therefore enhanced n-type electronic conductivity. In an attempt to create an n-type MIEC, our group has investigated a number of combinations of dopant elements, including (Ba0.33 Sr0.17 La0.5 )2 (In0.4 Fe0.6 )2 O5+δ ; however, all of the resultant MIECs have been p-type. 22 Here, we introduce the composition Ba1−y Ndy In1−x Mnx O3−δ : we believe this is the first reported mixed n-type electronic and oxide ionic conductor stable in an air environment.

Experimental Section Sample Preparation Ba1−y Ndy In1−x Mnx O3−δ (BNIM) were synthesized by the Pechini method using stoichiometric amounts of Ba(NO3 )2 (99.9%, Wako Pure Chemical Ind.), Nd(NO3 )3 · 6H2 O (99.9%, Sigma-Aldrich Co.), In(NO3 )3 · xH2 O (99.9%, Kojundo Chemical Laboratory Co.), and Mn(NO3 )2 · 6H2 O (99.9%, Kojundo Chemical Laboratory Co.). The metal nitrates were dissolved in distilled water with a 6 ∼ 9 times molar ratio of citric acid (99%, Sigma-Aldrich Co.) and propylene glycol (98%, Wako Pure Chemical Ind.). After mixing the mixtures until dissolved, the mixtures were heated to 150◦ C for polyesterification, resulting in the 6


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formation of a polymer resin. After sequentially calcining the polymer resin in air at 400◦ C and 900◦ C, the BNIM powders were pressed into pellets and sintered at 1375◦ C for 3h.

Characterization The crystal structures of the BNIM powders and polished pellets were determined by Xray diffraction (XRD) using a D8 ADVANCE (Bruker) diffractometer. XRD patterns of the samples were collected by CuKα radiation with 2θ from 20-120◦ at increment rates 0.2◦ /s. The lattice constants of the samples were refined by WPPD (Whole Powder Pattern Decomposition) based on the Pawley method, using TOPAS 4 software. The microstructure was then observed and a compositional analysis of each sample was performed by SEM-EDS using a JSM6360LA electron microscope (JEOL).

Electrical conductivity The electrical conductivity of BNIM samples was measured by the conventional DC fourprobe method and for temperatures ranging from 300 to 900◦ C in air, and for oxygen partial pressure (PO2 ) ranging from 0.21 to 10−14 atm. The data points were taken when the resistance change was less than 0.3% within a 600 second period. The typical equilibrium duration was approximately 2 hours for each data point. The DC resistance was calculated by linearly fitting measured voltages corresponding to eight different sourcing currents using precision voltage and current source measure units (Keithley 2400). The PO2 was controlled by a laboratory constructed YSZ oxygen pump.

Oxygen permeation and EMF measurements Electrochemical oxygen permeation measurements were conducted with symmetric pellet sample to confirm the oxygen permeation rate and determine oxide ionic transference numbers of BNIM for temperatures ranging from 750 to 900◦ C. The measurement was made on

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polished 1 ∼ 1.5 mm-thick bulk pellets with a diameter of 13 mm. To perform EMF measurement simultaneously, the diameter of 8 mm Pt paste for both sides are used as current collectors. The Pt paste was baked at 900◦ C for 2 h prior to the measurement. The details of the setup were described in our previous study. 23 The equilibrium duration was set to be 60 min for a temperature change and 15 min for a PO2 gradient change. The oxygen permeation rates were calculated by eliminating the mechanical leakage of O2 (estimated as 1/4 of detected N2 ) from detected O2 . The mechanical leakage was approximately 300 ppm of N2 . The transference numbers were calculated from the Nernst open circuit EMFs generated across the polished pellets with a series of different PO2 gradients.

Thermoelectric power measurement The EMF induced across the BNIM samples under temperature gradients were measured following the electrical conductivity measurement by utilizing the outer 2 electrodes of the 4-probe conductivity bar for temperatures ranging from 300 to 900◦ C in air, and for oxygen partial pressure (PO2 ) ranging from 0.21 to 10−14 atm. The temperature gradients and EMFs across the sample bar were monitored with digital multimeters (Keithley 2700). Transient temperature gradients were induced across the samples by nearby nichrome heaters operating with DC voltage current source TR6143 (ADVANTEST).

Results Defect chemistry model of donor doped oxygen deficient perovskitetype Ba-In-based oxides Defect chemistry can be a powerful tool for understanding the electrical properties of ionic compounds with different chemistries and dopant types. Key to such analysis is predicting the appropriate model for the target materials of interest, and being able to fit predictions

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to experimental data relating to the PO2 dependence of the concentrations of the various key ionic and electronic defects. In the case of donor-doped perovskite type Ba-In-based oxide, similar to Sr(Ti, Fe)O3−δ 15 the most suitable framework to describe the defect structure +3 +4 +4 is BaIn+3 1−x Mx O2.5+x/2 rather than BaIn1−x Mx O3−δ , even though most samples show a

perovskite-type structure. Please see the discussion in Ref. 15 for a detailed explanation. For the brownmillerite structure Ba2 In2 O5 , the ordered oxygen defects can be regarded as quasi-interstitial anion sites V× i , that occupy one sixth of the nominal oxygen sites relative to the perovskite end member. While the stoichiometry of the oxide increases with donor doping × •• due to the incorporation of the oxygen on V× i , Vi and the ionized oxygen vacancy VO cannot

be treated as the same species in the defect model due to the different energetic interactions with oxide anion sublattice. The high concentrations of M+4 x added to the Ba-In-based oxide ′′

results in high concentrations of the interstitial oxygen defects Oi occupying the V× i sites. ′′

Since Oi in quasi-interstitial sites can be expected to be much higher concentrations than cation vacancies, 24 the creation of oxygen interstitial and vacancy defects can be described by the anion Frenkel defect pair generation, utilizing Kröger-Vink notation, as shown in Eq. 3. 25 ′′

′′

× •• •• O× O + Vi → VO + Oi ; Kaf = [VO ][Oi ]

(3)

•• where Kaf is the equilibrium constant, [O× O ] is the density of the oxygen sites, [VO ] is the ′′

density of the ionized oxygen vacancies, and [Oi ] is the density of the interstitial oxide ′′

•• ions. Whereas considerable variation was found in both [VO ] and [Oi ], the change of [O× O]

caused by the change in oxygen defect concentration was so slight that it can be assumed to be constant. Additionally, the oxygen reduction reaction and intrinsic electron-hole pair generation with their equilibrium constants are shown in Eq. 4, 5. 1 1 − •• •• 2 2 O× O → O2 + VO + 2e ; KR = PO2 [VO ]n 2

(4)

null → e− + h• ; Ki = np

(5)

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where n and p are the concentrations of electrons and holes, respectively. Another key condition that needs to be satisfied is the overall charge neutrally, given in Eq. 6. ′′

•• 2[VO ] + p + [D• ] = 2[Oi ] + n

(6)

where [D• ] is the concentration of ionized donors. By solving Equations 3, 4, 5 and 6 simultaneously, we can obtain solutions for the four defect species in terms of the equilibrium constants Kaf , KR , Ki , [D• ] and PO2 at a series of isotherms. Fig. 3 shows an example of the predicted dependence of the key defect concentrations calculated using the values for equilibrium constants listed in the figure captions, as a function of the oxygen partial pressure (PO2 ). Solutions for the defect concentrations in the different regime of PO2 in which the Brouwer approximation has been applied, are shown in Table 1 and are plotted in Fig. 3 as dashed lines.

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Figure 3: Simulated Brouwer diagram of oxygen-deficient oxide with Kaf = 1037 (cm−6 ), 1 KR = 1055 (cm−9 · atm− 2 ), Ki = 1041 (cm−6 ) and [D• ] = 1021 (cm−3 ). Solid lines represent calculated defect concentrations based on analytic solution while dotted line are defect concentrations derived assuming applicability of the Brouwer approximation as in Table 1.

Table 1: The predicted solutions of defect carrier concentration used the Brouwer approximations Charge neutrality n p •• [VO ]

[O′′i ]

Regime I •• n = 2[VO ] √ −1/6 3 2KR · PO2 Ki √ 3 2KR

·

KR /4

·

1/6 PO2

√

Ki [D• ] KR [D• ]2

·

Regime III ′′ • √2[Oi ] = [D ] KR [D• ] 2Kaf

[D• ]

1/6 PO2

√ −1/6 3 KR /4 · PO2 Kaf √ 3

Regime II n = [D• ]

−1/2 PO2

Kaf [D• ]2 KR

11

·

1/2 PO2


−1/4

· PO2

2Kaf Ki2 KR [D• ]

·

2Kaf [D• ] [D• ] 2

1/4 PO2

√

Regime IV 2[O′′i ] = p KR Ki 2Kaf

3

√ 3

√ 3

2Kaf Ki2 KR

2 4KR Kaf

√ 3

−1/6

· PO2

Ki2 Kaf Ki2 4KR

1/6

· P O2

−1/6

· PO2

1/6

· PO2


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Note that in the transitions between different Brouwer defect regimes, no simple single PO2 dependence is predicted for the defects based on the analytical solution. This is particularly evident for narrow defect regimes, such as Regime III in Fig. 3. According to the simulated Brouwer diagram of this model, there are 4 different PO2 regimes. Assuming the mobility of defects are independent of PO2 , the electronic conductivity is predicted to first decrease, pass through a minimum, and then increase with reducing PO2 . The PO2 at the point where the electronic carrier density passes through a minimum can be calculated. That is, at the point where the transition from n-type to p-type electronic conductivity occurs, the electronic carrier density can be determined by setting the solutions for n and p equal to each other in defect regime III. This calculated transition PO2 is given by: PO2 =

KR2 [D• ]2 4Kaf2 Ki 2

(7)

From this equation, the PO2 at the n to p transition moves to higher PO2 with increasing ionized donor dopant concentration. This indicates that n-type conductivity can be achieved under oxidizing conditions for materials with donor dopant concentrations of a certain range. This prediction will be tested experimentally in this study.

Synthesis and characterization of Ba0.9 Nd0.1 In1−x Mnx O3−δ In order to evaluate the charge of electronic carriers, polycrystalline single-phase samples were first prepared and investigated. Details of the preparation stage are explained in the experimental section, and the XRD patterns of polished Ba1−y Ndy InO3−δ (BNIO) and Ba0.9 Nd0.1 In1−x Mnx O3−δ (BNIM) following sintering at 1375 ◦ C for 3h are shown in Fig. 4. Compared to La3+ as an A-site donor, Nd3+ as donor element is able to stabilize the perovskite structure with less doping due to its smaller ionic radii as shown in Fig. 5. 18 Meanwhile, secondary phases appear when the Nd content is over y = 0.2.

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(a)

Ba1-yNdyInO3-δ

(b)

Ba0.9Nd0.1In1-xMnxO3-δ perovskite Mn rich phase

Intensity (a.u.)

perovskite NdInO3

Intensity (a.u.)

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


y = 0.3 y = 0.2

x = 0.4 x = 0.3 x = 0.2 x = 0.1

y = 0.1 20

30

40

50

60

2θ [Cu-Kα]

70

80

x = 0.0 90

20

30

40

50

60

70

80

2θ [Cu-Kα]

Figure 4: X-ray diffraction patterns of polished pellets after sintering at 1375◦ C for 3h (a) Ba1−y Ndy InO3−δ (b) Ba0.9 Nd0.1 In1−x Mnx O3−δ .

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90


(a)

(b) Brownmillerite

Perovskite

4.2

Brownmillerite

Lattice constant / Å

4.3

Lattice constant / Å

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

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4.25

Perovskite 4.20

4.1

4.15 0.00

0.10

0.20

0.30

0.0

Nd content / %

0.1

0.2

0.3

0.4

Mn content / %

Figure 5: Lattice constant of polished pellets after sintering at 1375◦ C for 3h (a) Ba1−y Ndy InO3−δ (y = 0 is referenced from 18 )(b) Ba0.9 Nd0.1 In1−x Mnx O3−δ .

For the y = 0.1 BNIO sample, X-ray diffraction patterns assigned as a single perovskite-type phase are observed for x = 0.1 to 0.3 while secondary phases appear when x exceeds 0.3, pointing to a solid solubility limit of Mn in BNIM of approximately x = 0.3. Backscattering SEM also shows good agreement with XRD patterns (Supporting Information I). The maximum amount of Mn used in this study is then set to x = 0.3. Even though this high dopant concentration can be generally regarded as substitution, doping instead of substitution is used throughout this article to emphasize the donor dopant aspect.

N-type mixed conduction of Ba0.9 Nd0.1 In1−x Mnx O3−δ The total DC electrical conductivity represents the sum of the contributions from all electronic and ionic charge carriers (ignoring any interfacial contributions). Since the oxide ionic conductivity does not exhibit a significant dependence on PO2 when ionic defects pre14


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dominate (see defect Regime III in Table 1), 26 the PO2 dependence of the total electrical conductivity is commonly determined by the PO2 dependence of the electronic conductivity. In this study, the conventional 4-probe measurement method is adopted for the single phase cubic perovskite structure Ba0.9 Nd0.1 In1−x Mnx O3−δ . Fig. 6 shows the PO2 dependence of the DC electrical conductivity measured at 800◦ C for Ba0.9 Nd0.1 In1−x Mnx O3−δ (x= 0.1, 0.2 and 0.3).

2

Ba0.9Nd0.1In1-xMnxO3-δ

-1

Electrical conductivity σ / S・cm

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


0.1

0.01

x = 0.3 9 8 7 6 5

x = 0.2

4

3

x = 0.1 2

9

-14

-12

-10

-8

-6

-4

-2

0

log(PO2 / atm) Figure 6: PO2 dependence of the DC electrical conductivity Ba0.9 Nd0.1 In1−x Mnx O3−δ (x= 0.1, 0.2 and 0.3) (BNIM) sintered bars.

at

800◦ C

for

Here, the three distinct PO2 regimes observed are largely consistent with the simulated model predictions for Regime I, II and III in Fig. 3. In the oxidizing region (PO2 = 10−1 ∼ 10−3 atm), the electrical conductivity of x = 0.2 and 0.3 increases with decreasing PO2 ; this is a characteristic behavior of n-type conduction. The near plateau region at intermediate PO2 corresponds to the expected transition from Regime III characterized by large oxygen interstitial compensation for the ionized donor to Regime II characterized by large electron compensation for the ionized donor, as predicted in the simulation model. The minimum conductivity for x 15



= 0.1 indicates the occurrence of an n- to p-type transition at approximately log PO2 = -2.5. For higher Mn donor doping levels, the absence of an n- to p-type transition is consistent with the expectations of Eq. 7, which predicts that the n- to p-type transition should shift to ever higher PO2 . Indeed, the data in Fig. 6 suggest that the n to p transition shifts to PO2 ≥ 0.21 atm at higher Mn doping levels. In the intermediate region (PO2 = 10−4 ∼ 10−10 atm), the conductivities of the 3 compositions are largely PO2 independent, with the magnitudes approximately proportional to Mn content. In the reducing region (PO2 = 10−10 ∼ 10−14 atm), the conductivities increase with decreasing PO2 as expected for n-type behavior. In summary, the total electrical conductivity of Ba0.9 Nd0.1 In1−x Mnx O3−δ is in good agreement with the simulated models, suggesting that BNIM exhibits majority n-type conductivity with heavy Mn doping. The conductivity minimum, reported for the first time for Ba2 In2 O5 based oxides in the high PO2 regime, 27 is strong evidence of a p- and n-type transition under oxidizing conditions enhanced by increasing Mn doping. Based on a detailed analysis of the PO2 and temperature dependence of this conductivity around the conductivity minimum for x = 0.1, the energy band gap (Eg = 2.15 eV) and the PO2 -dependent ionic transference number were calculated for x = 0.1 (Details in Supporting Information II). Fig. 7 shows the temperature dependence of the electrical conductivity in air of the BNIM specimens in the form of Arrhenius plots.

16


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1000

-1

Electrical conductivity log(σ / S・cm )

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


Temperature, T / ℃

800 700 600

-1

500

400

300

Ba0.9Nd0.1In1-xMnxO3-δ x = 0.3

-2

x = 0.1

-3

x = 0.2 -4

-5 0.8

1.0

1.2

1.4

1000/T / K

-1

1.6

1.8

Figure 7: Temperature dependence of the DC electrical conductivity in air for Ba0.9 Nd0.1 In1−x Mnx O3−δ (x= 0.1, 0.2 and 0.3) (BNIM) sintered bars.

There are two points of interests: the source of the two temperature regimes characterized by different activation energies (see Table 2), and the dependence of conductivity on x. At reduced temperature, maximum conductivity occurs at x = 0.1; the conductivity decreases at x= 0.2 and increases again at x = 0.3. In the higher temperature regime, while the x = 0.2 composition again provides the lowest conductivity, the switch at x = 0.1 and x = 0.3 means that the x= 0.3 composition now exhibits the highest conductivity, which is consistent with the observations at 800◦ C in Fig. 6. Keep in mind that in the higher temperature regime, the x = 0.1 specimen is p-type while the other two compositions are n-type - see Fig. 6. In the low temperature regime, however, this is not likely to be the case (see discussion below), so it is not surprising that the relative conductivities of the different specimens cross. The activation energies of all BNIM samples are observed to increase around 600◦ C as shown in Table 2.

17



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Table 2: Activation energy of total electrical conductivity in air for Ba0.9 Nd0.1 In1−x Mnx O3−δ (BNIM) sintered bar. Temperature regime x = 0.1 x = 0.2 x = 0.3 900 ∼ 600 ◦ C 0.84 eV 1.05 eV 0.92 eV 600 ∼ 300 ◦ C 0.42 eV 0.62 eV 0.77 eV

A similar increase in activation energy was reported for donor doped pyrochlore type oxides such as Nb and Mo doped Gd2 Ti2 O7 . 28,29 The lower activation energy in the low temperature regime where oxygen stoichiometry might be fixed may be attributed to activated electron or hole hopping between Mn neighbors and/or ionic conductivity (see discussion below). Since Ba2 In2 O5 is known to be an ionic conductor and with donor doping as in BNIM, mixed ionic and electronic conductivity is expected. As such, it is important to establish the relative contribution of the ionic conductivity to the total conductivity as given by the oxide ion transference number, ti . To evaluate the transference number as a function of PO2 , the EMF across an oxygen concentration cell-type sample were measured and the differential form of the Nernst equation, Eq. 8, was employed. 30 dE kT = ti (PO2 ) d(lnPO2 ) 4e

(8)

where k and e are the Boltzmann constant and elementary charge, respectively. The measured partial pressure dependence of the oxide ion transference number at 800◦ C is shown in Fig. 8.

18


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Oxide ion transference number ti

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0.8

x = 0.1

0.6

x = 0.2

0.16

0.14

0.02

x = 0.3

0.01

@800°C -4.2

-4.0

-3.8

-3.6

-3.4

-3.2

-3.0

log (pO2 / atm)

Figure 8: Oxygen partial pressure dependence of Ba0.9 Nd0.1 In1−x Mnx O3−δ oxide ion transference number at 800◦ C evaluated by EMF measurement.

While it is likely that the transference numbers of MIECs are somewhat influenced by electrochemical leakage (Details in Supporting Information III), the transference number of BNIM measured in this manner was found to change dramatically from approximately 0.7 down to 0.01 as the Mn doping level increased from x = 0.1 to 0.3, clear evidence of a change in conduction mechanism from predominantly oxide ion to electronic conductivity in BNIM. Additionally, the PO2 dependence of the transference number can also reveal whether a por n-type conduction mechanism dominates at a given temperature and PO2 range (Further discussion in Supporting Information III). Since the oxide ionic conductivity shows a rather weak PO2 dependence, the positive slopes in Fig. 8 for x = 0.1 and 0.2 point to ptype conduction, whereas the negative slope for x = 0.3 provides further evidence of n-type conduction in the n-type MIEC. Moreover, all the BNIM specimens are observed to be electrochemically permeable under certain PO2 gradients which is the characteristic mixed oxide ionic-electronic conductivity. For x = 0.3, the oxide ionic conductivity estimated from permeation rate is in good agreement with the measured ion transference number. (Supporting 19



Information III). Meanwhile, the proton concentration is found to be relatively small in x = 0.3 (Supporting Information IV), so we can conclude that the conductivity change is not affected by proton conduction.

Charge of electronic carriers in BNIM Although the conductivity and ion transference number provide strong evidence of a p- to n-type transition with increased Mn doping, the charge of the electronic carrier is not directly confirmed. In the semiconductor field, thermoelectric power measurements are commonly used to establish the charge of the majority electronic carrier. Fig. 9 shows the temperature dependence of the absolute Seebeck coefficient for x = 0.1 ∼ 0.3 samples, which excluded

-1

the contribution of connecting Pt wires. (Details in Supporting Information V). 31 600

Seebeck coefficient, S / µV・K

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

Page 20 of 33

x = 0.1

500 400 300 200

P-type conduction x = 0.2

100

50 0

x = 0.3

-50

N-type conduction

-100

300

400

500

600

700

800

Temperature, T / ℃

900

Figure 9: Temperature dependence of Seebeck coefficient in air for Ba0.9 Nd0.1 In1−x Mnx O3−δ . It is interesting to note that the absolute Seebeck coefficient changes from positive to negative values when the Mn doping level increases from x = 0.2 to 0.3 (The very small positive thermoelectric voltage across the x = 0.3 sample at high temperature falls within noise limits). This change in the sign of the charge carrier strongly suggests a transition from pto n-type electronic conduction, and agrees with the observed dependence of conductivity 20


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and ion transference number on increasing Mn doping level. Generally speaking, p and n in respective p- and n-type semiconductors increase with increasing temperature, leading to decreases in the Seebeck coefficient.This is consistent with the observations seen in Fig. 9. For MIECs, the interpretation of the Seebeck coefficient is more complicated because of the nonnegligible contributions of ionic conductivity to the total conductivity. 32 As confirmed in the last section, the oxide ionic conductivity is predominant in BNIM with x = 0.1. Due to the reduced contribution of oxide ionic conductivity at low temperatures (below 450◦ C), BNIM behaves in a manner more similar to a semiconductor, the absolute value of BNIMs ’Seebeck coefficient decreases as expected. Moreover, the estimated distance between the Fermi energy and the valence band for x = 0.1 (Details in Supporting Information IV) is confirmed to be in good agreement with the estimated one by using the conductivity minimum in Supporting information II. For higher temperature, the total Seebeck coefficient can be divided into three contributions: that from ions, electrons and the hole conduction, as shown in Eq. 9. 33

SMIEC =

1 1 qi qn qp 1 si [ti ( S0 (O2 ) − RlnPO2 − − ) − tn (sn − ) + tp (sp + )] F 4 4 2 2T T T

(9)

where F , tj , S0 (O2 ), sj and qj is the Faraday constant, transference number, standard molar entropy of oxygen, and partial molar entropy and heat of transfer, respectively. It is apparent from this equation that the Seebeck coefficient is negatively correlated with temperature as long as the transport coefficients do not change. Thus, the increase in the Seebeck coefficient at x = 0.1 around 600◦ C can be explained by the increasing contribution of the transport coefficient: that is, the ionic transference number increases as the oxide ionic conductivity becomes predominant. The PO2 dependence of the Seebeck coefficient for the three BNIM compositions shown in Fig. 10 can also be separated into three different distinctive regimes which correspond to the electrical conductivity explained in the last section.

21


-1

Seebeck coefficient, S / µV・K

Ba0.9Nd0.1In1-xMnxO3-δ 300

x = 0.1

200

x = 0.2

100

S = 0 μV/K

0

x = 0.3

-100

-14

-12

-10

-8

-6

-4

-2

0

log(PO2 / atm) Figure 10: PO2 dependence of the Seebeck coefficient at 800◦ C for Ba0.9 Nd0.1 In1−x Mnx O3−δ (x= 0.1, 0.2 and 0.3).

In the oxidizing region, the positive values of the Seebeck coefficients of BNIM with x = 0.1 and x = 0.2 at high PO2 decrease with decreasing PO2 , and ultimately switch to negative values when the oxygen partial pressure reaches approximately PO2 = 10−3 atm. At x = 0.3, smaller but always negative values for the Seebeck coefficient are observed. The inversion in sign for x = 0.1 and x = 0.2 with decreasing PO2 and the obvious trend from p towards ntype with increasing Mn content indicates that the switch-over from p-type to n-type can be controlled in BNIM by the appropriate choice of Mn content and PO2 at a given temperature. In the intermediate PO2 region, the Seebeck coefficient for all three compositions is negative with a near plateau (see Region II in Fig. 3). It can also be seen from Fig. 10 that the Seebeck coefficient is inversely proportional to the dopant concentration, as predicted by Eq. 10. 34 k N S = − (ln + A) e n

(10)

where k, e, N and A are the Boltzmann constant, elementary charge, conduction band

22

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density of states and transport coefficient, respectively. Assuming a PO2 -independent density of states and a near zero transport coefficient, as appropriate for polaron conductors, 32,35 the Seebeck coefficient is predicted to exhibit an inverse dependence on ln carrier concentration. Finally, in the reducing region, while the Seebeck coefficients of the three specimens remain negative, they all show a decreasing magnitude towards zero with decreasing PO2 , again consistent with the predictions that n increases with decreasing PO2 in Regime I of the defect diagram. Of particular note is that the crossover in the Seebeck coefficient at x = 0.1 from positive to negative in the oxidizing region with decreasing PO2 corresponds closely to the PO2 at which the electrical conductivity goes through its minimum value (shown in Fig. 6). However, PO2 for the Seebeck inversion at x = 0.1 is one order of magnitude lower than the conductivity minimum, and x = 0.2 also exhibits an inversion of the Seebeck coefficient even though there is no observed transition point for the conductivity as a function of PO2 . These inconsistencies can be explained by considering the contributions made by other major carriers to the Seebeck coefficient. Here we expect significant contributions from oxygen vacancies which are incorporated in Eq. 11. 35

ST otal =

Selec × σelec + Sion × σion σT otal

(11)

The Seebeck coefficient of the predominant oxide ionic conductors commonly shows negative contributions to S 35,36 : this would explain the shifts in the PO2 at which the Seebeck inversion of x = 0.1 and 0.2 occur. Also, the inversion of Seebeck coefficient at x = 0.2 suggests the n-type and p-type carriers have similar concentrations at this composition.

Discussion The important point of this study is to determine how both n-type electronic and oxide ionic conductivities can coexist in an oxidizing atmosphere in this particular material. Although 23



Page 24 of 33

there are some previous reports that Mn-based perovskite-type oxides such as CaMnO3−δ show n-type conductivity, 37 the defect model of CaMnO3−δ fails to explain the experimental data of BNIM. According to the defect model of CaMnO3−δ in the previous report, the n-type conductivity of CaMnO3−δ comes from the intrinsic Mn3+ , Mn4+ and Mn5+ distribution as shown in Eq. 12 and 13. 38 1 ′ × •• O2 (g) + VO + 2MnMn = O× O + 2MnMn 2 ′

• 2Mn× Mn = MnMn + MnMn

(12) (13)

The case of BNIM is considered to be different due to the extrinsic defect of Mn doping. The difference is consistent with the different magnitude of the Seebeck coefficient between CaMnO3−δ (-200 µV/K at 800◦ C) and BNIM (-few µV/K at 800◦ C), indicating the large difference in electronic carrier density. Applying the theory used to explain the properties of CaMnO3−δ , it remains difficult to explain the plateau that we observed in BNIM (Fig. 10) due to the continuous change of [Mn3+ ] and [Mn5+ ] as a function of PO2 . 38 Meanwhile, based on our defect model, one question, this is of particular interest, is why the PO2 that represents the predicted PO2 n-type to p-type transition point does not follow the expectations of Eq. 7 with increasing Mn content from x = 0.1 to x = 0.2. Eq. 7 predicts a shift in PO2 by a factor of 4, not the much larger observed values, i.e. 2 orders of magnitude higher than the oxidizing regime. This, of course, assumes that the equilibrium constants KR , Kaf , Ki are independent of the Mn concentration as part of the dilute solution approximation. With the high Mn contents examined here, it is not surprising that the dilute solution approximation is not satisfied and that the equilibrium constants change with increasing Mn content. Although we acknowledge that further investigation is necessary to clarify the reason for this, we tentatively conclude that, based on the shapes of the PO2 dependence of the electrical conductivity, the most likely source of the nonideality associated with the dilute solution approximation is the dependence of the intrinsic Frenkel

24


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term Kaf on Mn concentration (Details in Supporting Information VI). The near plateau region in Fig. 6 has been attributed to electrons compensating ionized donors. In conventional semiconductors, because the donor states lie close to the conduction band, they are fully ionized at elevated temperature and exhibit a negligibly small activation energy. This is also observed for well-known donor doped semiconducting oxides such as La doped SrTiO3 . 39,40 Table 3 shows the activation energies derived by examining the conductivity of our BNIM specimens as a function of reciprocal temperature at log (PO2 ) = -8.0 atm, within the plateau region. Table 3: Activation energy of total electrical conductivity at log at log(PO2 ) = -8.0 atm for Ba0.9 Nd0.1 In1−x Mnx O3−δ (BNIM) sintered bar. x = 0.1 x = 0.2 1.23 eV 1.13 eV

x = 0.3 0.88 eV

Surprisingly large activation energies (> 0.88 eV) are obtained with the activation energy negatively correlated to the level of Mn doping (Details in Supporting Information VII). There are a number of factors known to contribute to large activation energies. The most obvious is that because Mn is a deep donor, the concentration of ionized donors and thereby quasi-free electrons may be only a small fraction of the total Mn concentration. Furthermore, impurity bands introduced by doping are often narrow, and while they support extrinsic electronic transport, they do so only via a small polaron activated hopping process. These features are discussed in greater detail in Supporting Information VI. The reason for the negative correlation of the activation energies and Mn content is consistent with the expected widening of the Mn derived impurity band, 41 leading to both a decreased donor ionization energy and an enhancement in hopping percolation. 42 Another point of interest which emerges from the results is the unexpected suppression ′′

of that rate of increase of [Oi ] with donor doping. According to the theoretical defect ′′

model, [Oi ] is considered to be the majority negative carrier compensating the donor under 25



Page 26 of 33

′′

sufficiently oxidizing conditions (see Region III in Table 1). Even though the absolute [Oi ] concentration is predicted to increase with increases in Mn doping, the magnitude of Kaf ′′

determines the PO2 values of the ionic compensating regime, in which [Oi ] compensates the donor (see Eq. 7). If the Kaf is decreased, the ionic compensating regime can be shifted ′′

to considerably higher PO2 , leaving a much lower [Oi ] in air than would be expected if we assume a constant Kaf (Details in Supporting Information VIII). Overall, the increase of ′′

[Oi ] is suppressed with increasing donor content, resulting in the enhanced release of the electrons from the donor with increasing Mn doping at the expense of ionic compensation.

Conclusion In this study, mixed n-type electronic and oxide ionic conductivities are confirmed to coexist in Ba0.9 Nd0.1 In0.7 Mn0.3 O3−δ at high temperature and high PO2 . This important conclusion is based on the observed PO2 dependence of electrical conductivity in concert with predictions based on our donor doped defect model, the measured oxide ion transference numbers, electrochemical oxygen permeability and the negative sign of the Seebeck coefficient. Moreover, the transition from p-type to n-type conductivity in BNIM is consistent with the predictions made by our defect chemical model based on donor doped materials with extensive anion Frenkel disorder. The achievement of co-existing n-type electronic and oxide-ion conductivities in this material system can be attributed to the high intrinsic oxygen-deficiency and significant change of anion intrinsic disorder of the host brownmillerite structured Ba2 In2 O5 upon donor doping, a feature reported here for MIEC materials for the first time. This novel n-type MIEC suggests the potential for enhancing the oxygen surface reduction of SOFC cathodes at reduced temperatures and thereby identifying new candidate cathode materials for next-generation SOFCs. However, the electronic conductivity of BNIM is 2 to 4 orders lower than the popular SOFC cathode materials. While the exchange reaction may not be limited by the lower electronic conductivity, this shortage could limit the overall perfor-

26


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mance on SOFC due to the limited electronic charge collection. One possible solution could be adding a second highly electronically conducting phase to the MIEC electrode to serve as a current collector. Meanwhile, interface degradation due to interaction between BNIM and the solid electrolyte could also influence the cathode behavior. Such studies are planned for the near future.

Acknowledgement YC would appreciate the financial support from the Interdepartmental Doctoral Degree Program for Multi-Dimensional Materials Science Leaders in Tohoku University, one of the Leading Graduate School Programs run by the Ministry of Education, Culture, Sports, Science and Technology. HT would like to acknowledge the financial support provided by JSPS (18H03832). HLT was supported by grant DE SC0002633 funded by the U.S. Department of Energy, Office of Basic Science.

Supporting Information Available Synthesis and characterization of BNIM. Further discussion and analysis of electrical conductivity, transference number and Seebeck coefficient. Simulation of the change for the mass constant.

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Graphical TOC Entry

33


Stabilizing coexisting n-type electronic and oxide-ion conductivities in

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