Crystal Structure and Transport Properties of Oxygen-Deficient

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Crystal Structure and Transport Properties of Oxygen-Deficient Perovskite Sr0.9Y0.1CoO3-# Tianrang Yang, Victoria Mattick, Yan Chen, Ke An, Dong Ma, and Kevin Huang ACS Appl. Energy Mater., Just Accepted Manuscript • DOI: 10.1021/acsaem.7b00275 • Publication Date (Web): 29 Jan 2018 Downloaded from http://pubs.acs.org on February 1, 2018

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ACS Applied Energy Materials

Crystal Structure and Transport Properties of Oxygen-Deficient Perovskite Sr0.9Y0.1CoO3-δδ Tianrang Yanga, Victoria F. Mattickb, Yan Chenc, Ke Anc, Dong Mac and Kevin Huanga* a

Department of Mechanical Engineering, University of South Carolina, Columbia, SC 29208 b

Department of Chemical Engineering, University of South Carolina, Columbia, SC 29208 c

Neutron Scattering Division, Oak Ridge National Laboratory, Oak Ridge, TN 37831 *Corresponding author: [email protected]

Abstract The present work reports a systematic study on temperature-dependent local crystal structure, oxygen stoichiometry and electrical/electrochemical properties of an oxygen-deficient Sr0.9Y0.1CoO3-δ (SYC10) perovskite using variable temperature neutron diffraction (VTND), thermal gravimetric analysis (TGA) and electrical/electrochemical methods, respectively. The VTND reveals that the crystal symmetry of SYC10 remains P4/mmm tetragonal up to 900 oC. The tetragonal symmetry reflects the net effects of temperature and oxygen stoichiometry on crystal symmetry. The observed p-type electronic conductivity behavior originates from the charge ordering between the two distinctive Co sites. The partial oxide-ion conductivity and diffusivity obtained from oxygen permeation measurements are 2.3 × 10-2 S cm-1 and 7.98 × 10-8 cm2/s at 800 oC in air, respectively. The electrochemical oxygen reduction reaction kinetics of the SYC10 cathode is primarily limited by the charge-transfer process at low temperatures (600650 oC) and oxide-ion migration from the cathode into the electrolyte at high temperatures (700800 oC).

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Keywords: perovskite; oxygen stoichiometry; charge-ordering; electronic conductivity; oxideion conductivity;


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I.

Introduction

A key to the success of intermediate-temperature solid oxide fuel cells (IT-SOFCs) is to discover electrocatalytically active oxides to promote the limiting oxygen reduction reaction (ORR) at the cathode. It has been demonstrated in the past that perovskite-related and oxygen-deficient complex oxides are a promising class of cathode materials for IT-SOFCs [1-8]. Among them, BaCoO3-δ (BCO) and SrCoO3-δ (SCO) based simple perovskite oxides attract the most interest because of their high intrinsic catalytic activity in regards to ORR. However, due to their overly high Goldschmidt tolerance factor (>1.0), pure BCO and SCO are not stable under high temperatures and ambient pressure conditions, making their use impractical. Previous studies have shown that a proper doping of the A- and/or B-cation in BCO and SCO can stabilize the perovskite structure when subjected to elevated temperatures, which has become a major driver for active studies on these two classes of materials in recent years [9-20]. Another challenge for BCO and SCO-based perovskites to be used as practical cathodes is their much higher thermal expansion coefficient (TEC, >20 ppm K-1) when compared to the electrolyte (~10 ppm K-1). A direct use of these two materials in SOFC as a bulk cathode is, therefore, inadequate. To solve this issue, the ORR-active, but TEC-high BCO- and SCO-based materials, are usually utilized as a cathode in the form of nanoparticles (NPs) impregnated into a TEC-compatible scaffold [21-23]. In this design, TEC of the cathode is determined by the scaffold, not the attached NPs. Another critical issue particularly associated with BCO-based perovskites is their high chemical propensity to react with CO2 in air, resulting in a gradual formation of BaCO3 and long-term performance decay. In comparison, SCO-based perovskites



are less reactive with CO2, thus they have a better potential to be considered a practical IT-SOFC cathode [24, 25]. In recent years, there have been many reports on SCO-based oxygen-deficient perovskites in the literature [13-20, 26-28]. A general conclusion from these prior studies is that the thermal stability of these materials can be significantly enhanced through donor-doping (dopant with higher oxidation state than the host cation) on either A- and B-cation. Among them, Sr0.9Y0.1CoO3-δ (SYC10) [26, 27] has been reported as an IT-SOFC cathode material with low polarization resistance (Rp) [27]. From the room temperature structural analysis using X-ray diffraction (XRD), both tetragonal (P4/mmm, #123) [26] and cubic (Pm-3m, #221) structures have been observed to exist in SYC10 [27, 28]. Since SYC10 operates at elevated temperatures when employed as a SOFC cathode, the knowledge of the local structure and oxygen stoichiometry at elevated temperatures would be more meaningful than that under ambient temperature to interpret the ORR activity. However, the conventional XRD techniques, whether at room temperature or at higher temperatures, cannot provide accurate information on the oxygen positions/occupancies due to a much weaker scattering contribution of oxygen to the total Bragg scattering than that of cations. In this study, we report a systematic study on temperature-dependent local crystal structure and oxygen occupancy (stoichiometry) in SYC10 determined by variable temperature neutron diffraction (VTND). To complement VTND results, conventional thermogravimetric analysis (TGA) and iodometric titration for oxygen nonstoichiometry, oxygen permeation for partial oxide-ion conductivity and impedance spectroscopy for polarization resistance are also carried out. With these fundamental data in hand, we attempt to correlate the crystal structure and oxygen nonstoichiometry with transport properties such as partial oxide-ion conductivity, 4 ACS Paragon Plus Environment

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oxygen-related diffusivity and polarization resistance to further elucidate the ORR mechanism in SYC10.

II.

Experimental Procedure

2.1 Sample synthesis The SYC10 was prepared by a conventional solid-state reaction method. Stoichiometric amounts of SrCO3 (≥99.9%, Aldrich), Co3O4 (99.7%, Alfa Aesar) and Y2O3 (99.9%, Alfa Aesar) were first ball-milled in ethanol for 3 h. After drying, the powder was then pressed into pellets, followed by calcination at 1000 oC for 12 h with a 3 oC min-1 heating and cooling rate. The pellets were then broken up and ball-milled for 3 h, and calcined again at 1050 oC for another 12 h. Another 3 h ball-milling was followed before making the final pellets or rectangular bars. The final pellets, or rectangular bars, embedded in SYC10 powder were sintered at 1230 oC for 10 h to ensure phase purity and homogeneity. 2.2 Variable temperature neutron diffraction Variable temperature neutron diffraction (VTND) experiments were conducted on the sintered pellets (~17 mm in diameter and ~10 mm long) situated in an MgO crucible using VULCAN, the time-of-flight Engineering Diffractometer at the Spallation Neutron Source in Oak Ridge National Laboratory [29, 30]. The temperature was varied from the base furnace temperature (70 o

C) to 900 oC, while the atmosphere remained in ambient air. The VTND data were collected

continuously during both heating and cooling with a heating rate of 10 oC min-1 below 300 oC and 5 oC min-1 between 300 and 900 oC, and a constant cooling rate of 10 oC min-1. During heating, the samples were held at 300 oC for 3 h and for 2 h at three other temperatures, i.e., 500, 5 ACS Paragon Plus Environment


700 and 900 oC. Only data collected during the last 1-h with total neutron counts of 4×106 at each temperature were used for the structural refinement. The beam gauge volume was determined to be 5×10×5 mm3 by the incident slits and receiving collimators. RT patterns were collected ex-situ from a 5 g powder sample. The Rietveld refinements were performed with a GSAS program using the EXPGUI interface [31, 32]. Typically, a refinement starts with trials and errors using the model-free Le Bail method [33]. The background and peak profiles were modeled with the shifted Chebyshev polynomial function and Pseudo-Voigt function with Finger-Cox-Jephcoat asymmetry correction, respectively. The DIFA (one of the three parameters defining relationship between measured time-of-flight and d-spacing) [31], unit cell, background and profile parameters from Le Bail fitting were used as the initial inputs into the Rietveld refinement and kept fixed during the first stage. Scaling factors, atomic positions, isotropic thermal displacement parameters and absorption coefficients were then refined separately. Following that, all the fixed parameters were released and refined together. Anisotropic thermal displacement factors were refined depending upon allowance by the data. Oxygen occupancies were the outcomes of the final run. The DIFA and absorption coefficients of in-situ ND were considered temperature-independent. The crystal structure was constructed with Vesta 3.0 [34]. 2.3 Thermogravimetric analysis TGA was performed with a NETZSCH STA 448 TGA/DSC thermal analyzer. Approximately 50 mg powder was placed in an alumina crucible under a flow of 60 ml min-1 air. The sample was held at a reference state of 70 oC for 2 h to remove the absorbed H2O, followed by the same temperature profiles used for VTND runs. With the initial RT oxygen stoichiometry determined


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from iodine titration or ND refinement, the oxygen stoichiometry as a function of temperature can then be mapped according to weight changes. 2.4 Iodometric titration In addition to ND refinement, the initial oxygen stoichiometry (3-δ) at RT was also determined by iodometric titration. The procedure is described as follows. Ultrapure N2 is bubbled into 30 ml of diluted 1M HCl solution for 1 h to eliminate the dissolved O2 and Cl2. Then 50 mg powder was dissolved in the above mentioned solution under a N2 atmosphere, followed by adding 10 ml of 0.2M KI solution to reduce all Co3+/4+ to Co2+. Next, 0.01 M Na2S2O3 is added dropwise to titrate the released elementary iodine by the reduction reaction. A few drops of saturated starch solution was used to indicate the endpoint of titration. With a known Na2S2O3 concentration, the average Co-ion oxidation state of the sample can then be calculated based on the following charge balance:

Co + nI = Co + I

(1)

I + 2S O = 2I + S O

(2)

To ensure accuracy, the concentration of Na2S2O3 solution needs to be calibrated. To do this, 25 mg K2Cr2O7 powder is first dissolved in the 1M HCl solution mixed with an excessive 0.2 M KI 3+ solution. The Cr O by KI. Following the same is expected to be fully reduced to Cr

procedure described above, the produced elementary iodine is titrated using the Na2S2O3 solution. The concentration of Na2S2O3 can then be determined from the charge balance between reaction (2) and (3). Cr O + 3I + 7H O + 6I + 14H = 2Cr

2.5 Electrical conductivity measurement 7 ACS Paragon Plus Environment

(3)


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Bar samples (5mm×2mm×15mm) were used for total conductivity (σ) measurements by the standard four-probe method. Both T- and Po2-dependent σ were measured from 800 oC to 500 oC with an interval of 100 oC and Po2 of 1, 0.5, 0.2, 0.1, 0.05 atm. Since it takes time for O2 to fully equilibrate with SYC10, we adopted the following wait-time for each temperature to ensure the accuracy of the conductivity: 3 h at 800 oC, 5 h at 700 oC, 10 h at 600 oC, and 18 h at 500 oC. Clearly, a lower temperature requires longer wait-time to reach a true gas/solid equilibrium, which also infers that this material has a low oxygen surface exchange rate at low temperatures. The resistance data were collected in the form of V-I curve with the E-I testing module available in the CorrWare program operated on a Solartron 1470 multichannel potentiostat. 2.6 Partial oxide-ion conductivity determination The partial oxide-ion conductivity in SYC10 was determined by the classical Wagner equation using the measured oxygen permeation flux data: 567869:6 ;66? -34

!

J = − % "# $

-34

t ' t ( σ*+*,- d[lnP ]

(4)

where J denotes the oxygen permeation flux in mol cm-2 s-1; R is gas constant 8.314 J mol-1 K-1; F is Faraday constant 96,485 C mol-1; L is membrane thickness in cm; t ' and t ( are the ionic and electronic transport number, respectively. For a MIEC with a dominant σ( , t ( ≈1, eq. (4) is then simplified into !

567869:6 ;66? -34

J = − % "# $

-34

σ' d[lnP ]

(5)

Since σ' varies with Po2 by a power law relationship: σ' = σ+' P 8 ACS Paragon Plus Environment

(6)

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substitution of eq. (6) into eq. (5) followed by integration gives: D(EF(,*(

A

J = BCP

A=−

H((IJK

G −P

L

(7)

NO ' ! 4

(8)

"# $

D(EF(,*(

Here A and n values are constants that can be obtained by fitting experimental J vs P H((IJK

at a fixed P

using eq. (7). The partial oxygen-ion conductivity σ' can then be calculated

from eq. (6) with σ+' value obtained from eq. (8). The actual permeation cell setup schematic was given in Fig. S1 of SI. Briefly, a dense SYC10 membrane with a thickness of 1.15 mm was first sealed onto a supporting alumina tube by a commercial silver paste (Shanghai Research Institute of Synthetic Resins). Air flowing at 100 sccm was fed to one surface of the membrane, while the other surface was swept with ultra-pure Ar at variable flow rates. To vary Po2 at the permeate side, the sweeping Ar flow rate was changed from 5 to 200 sccm. All gas flow rates were controlled by mass flow controllers (SmartTrak 50 series). Since the cell needs to be first sealed at 800 oC, all measurements were made from 800 oC down to 650 at 25 oC interval. The equilibration time at a specific temperature varied from 2 h at 800 oC to 12 h at 650 oC. The concentrations of O2 and N2 in the effluent were analyzed by an online gas chromatography (Agilent 490). The final flux densities of oxygen (J ) were calculated by the following equation: J =

P4

"P4 PQ

S

×A

(9)

where C and CT are the measured concentrations of O2 and N2, respectively; Q is the flow rate of sweeping Ar gas in sccm; A is the effective area of the sample (0.745 cm2). Any detectable N2 9 ACS Paragon Plus Environment


is corrected in eq. (9). The detected N2 flux is very small, typically 500 oC do O2-sites show vacancies. In comparison, major oxygen vacancies (V+∙∙) are populated at O3site within the temperature range measured. A further plot of c-axis coordination of 2h Sr (ideally 0.25) and 2g O (ideally 0.75) shown in Figure 4 (b) reveals variations of the structural 13 ACS Paragon Plus Environment


distortion with temperature. Overall, the structural distortion is resulted from two competing effects: higher temperature promoting symmetry and oxygen stoichiometry variations promoting distortion.

Figure 4. (a) Oxygen occupancies variation with temperature; (b) c-axis coordination of Sr and O2 determined from VTND. The deviation of Sr and O2 from their ideal position (Sr ½, ½, ¼, O2 0, 0, ¾) indicates the distortion level.

3.2 Lattice constants vs temperature The lattice constants a and c of the SYC10’s tetragonal structure are plotted against temperature in Figure 5a and b, respectively. It appears that there is a deviation from the linear trending at 500 oC, above which a greater increase of a and c with temperature is observed. We believe that this increased lattice expansion reflects an additional contribution from chemical oxygen loss at O2 site above 500 oC as shown in Fig. 4a. Since O2 is the bridging oxygen between the two distinct CoO6 octahedra layers, its deficiency could create dimensional changes along both a and c axis as observed in Figure 5a and b. The TGA curve shown in Figure 6a indeed indicates a greater weight loss above this temperature. In addition, the decrease of average Co-ions 14 ACS Paragon Plus Environment

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oxidation state as shown in Figure 6b also contributes to the curvature of lattice constant a and c vs T. The fitted linear thermal expansion coefficients (LTECs, α) along a-axis (Figure 5c) and caxis (Figure 5d) are listed in Table 2 for both low and high temperature regions. The average α is calculated by (2αa+αc)/3.

Figure 5. Variations of lattice parameters and thermal expansion coefficients with temperature in SYC10. (a) and (c) along a-axis; (b) and (d) along c-axis. Table 2. The fitted linear thermal expansion coefficients α along a-axis and c-axis. o o o α×10-6/K-1 RT-500 C 500 C-900 C

a-axis

19.5

34.8

c-axis Average

19.5 19.5

22.1 30.6



3.3 Oxygen stoichiometry vs temperature The oxygen stoichiometry (3-δ) at RT determined by iodometric titration is 2.580, a value which matches well with 2.543 obtained from VTND’s Rietveld refinement. The temperature dependence of 3-δ with the initial value from iodometric titration can be further calculated from the TGA results of Figure 6a; the results are shown in Figure 6b and compared with those obtained from VTND. The two curves generally follow the same trend, but the deviation between the two methods becomes larger above 700 oC. A possible reason for the enlarged deviation is the increased uncertainties due to the increased background noise in VTND refinement. It is interesting to note that both TGA and VTND confirm an oxygen gain occurring from RT to 300 oC. We believe that this oxygen gain is resulted from gaseous O2 occupying the “frozen oxygen vacancies” trapped by the slow SYC10/O2 exchange kinetics during cooling; the compacted samples used for enhancing cation diffusion and phase formation could be the reason for the latter.

Figure 6. (a) TGA curves measured in air; (b) oxygen stoichiometry (3-δ) and Co-ion oxidation state from both ND and TGA. The initial value for TGA-derived (3-δ) was taken from iodometric titration.


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3.4 Electrical conductivity vs temperature (T) and partial pressure of oxygen (Po2) The variations of electrical conductivity (σ) of SYC10 with T and Po2 are shown in Figure 7. In general, at a given temperature, the conductivity increases with Po2, suggesting a p-type electron hole conduction mechanism. However, at a given Po2, the conductivity decreases with temperature. This infers that increasing temperature substantially decreases hole-concentration (or oxidation state of Co-ions). Clearly, oxygen loss is one major reasons for the lowering of hole-concentration, thus electronic conductivity. The average Co-ion oxidation state obtained from TGA and VTND as shown in Figure 6b is generally lower than +3.0, particularly between 500 to 900 oC. What causes the observed p-type electronic conduction behavior at such a low oxidation state of Co-ion is an intriguing scientific question that should be addressed. With the VTND data, we are able to perform bond valence sums (BVS) analysis for the oxidation state of Co-ions within their local coordinating environment. In theory, the BVS of a cation i is the sum of the individual bond valences vJX defined by [37]: vJX = exp

EO E=\ ]

(10)

where j is the surrounding cation; rJX is the observed bond lengths between i and j; r+ is the ideal bond length (1.70 Å for Co3+-O [38]) and b is an empirical constant (0.37 Å). The rJX values are taken from VTND refinement results listed in Table S1. The calculated BVS results are shown in Table 3, where a remarkable difference in BVS between Co1 and Co2 is observed from RT to 900 oC, suggesting a charge ordering at these two Co-sites, i.e. a mixed oxidation state of Co3+/4+ on the Co1-site and Co2+/3+ on the Co2-site. The determined oxidation state of Co-ions from oxygen stoichiometry reflects the average value of Co1 and Co2. 17 ACS Paragon Plus Environment


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Figure 7. (a) Temperature dependence of σ at different Po2; (b) Arrhenius plots of σ. Table 3. Bond valence sums of Co-ions. 300 oC 500 oC 700 oC 900 oC

Co-ion

RT

Co1

3.82

3.75

3.62

3.57

3.37

Co2

2.65

2.55

2.55

2.41

2.44

With the charge ordering mechanism, it is possible to postulate that the SYC10 can exhibit either p-type electron hole conduction via Co1 or n-type excess electron conduction via Co2, depending on actual T and Po2 conditions. In general, hopping of electrons between two neighboring Co-ions is mediated by the conduction band formed by the orbital overlap between Co-3d and O-2p commonly known as a double exchange mechanism [39]. The stronger the overlap, the broader the band and therefore the easier for electrons to hop. However, the presence of oxygen vacancies at O3-sites in coordination with Co2 (see Figure 3) distorts the crystal field and weakens the overlap between Co2-3d and O-2p, resulting in a narrower conduction-band formation and promoting n-type electronic conduction with a higher energy barrier [40]. 18 ACS Paragon Plus Environment

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• Meanwhile, the positively charged donor dopant ^_` can also trap negatively charged n-type

electrons, making the total electronic conductivity in SYC10 dominated by p-type electron holes. It is to be noted that the portion of n-type electronic conduction will increase with a narrowed band, increased oxygen vacancy concentration and activation energy as T increases and Po2 decreases. Figure 7b indeed shows such a trending. Finally, the level of activation energy shown in Figure 7b seems to suggest that electron holes conduction in SYC10 gradually transitions from a delocalized to a localized state as Po2 is decreased at a fixed T. This should also be the case as T is increased at a given Po2.

3.5 Partial oxide-ion conductivity and oxygen self-diffusion coefficient Partial oxide-ion conductivity (σ' ) of a mixed electron and oxide-ion conductor (MIEC) plays a critical role in oxygen electrocatalysis for reversible SOFCs. Therefore, σ' is often used to evaluate whether an oxide is a suitable cathode material for SOFCs. However, determination of σ' is a challenging task due to the domination of electronic conductivity σ( in these materials. To separate σ' from σ( , a number of experimental methods have been proposed in the past [41-44], among which oxygen permeation is deemed the most reliable and reproducible method [45]. Therefore, in this study we adopt the oxygen permeation method to determine σ' . Figure 8 shows the measured oxygen flux density of SYC10 vs Po2 at different temperatures, along with the modeled curves using eq. (7). The detailed model values are listed in Table S2. The oxygen permeation flux is nearly 1-2 orders of magnitude lower than those of similar systems SrCo0.9Nb0.1O3-δ [13,46-52, 64] and SrCo0.9Ta0.1O3-δ [19,53,54] under a similar testing condition. For example, at 800 oC, a 1-mm thick SrCo0.9Nb0.1O3-δ [48] and 1.36-mm thick SrCo0.9Ta0.1O3-δ [53] can yield an oxygen flux density of 1.5×10-6 and 6×10-7 [53] mol cm-2 s-1, 19 ACS Paragon Plus Environment


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respectively, versus 8.9×10-8 mol cm-2 s-1 for a 1.15-mm thick SYC10 of this study. Clearly, the cubic structure of B-site doped SCO [35, 54] favors a higher partial oxide-ion conductivity than the tetragonal structure A-site doped counterpart. The σ' of SYC10 at Po2=0.21 atm obtained from the oxygen flux is shown as a function of temperature in Figure 9 and analytically expressed by: ab' =

.d×"ef g

× exp

"."he.ei jk lg

(11)

The activation energy Ea=1.16 ± 0.05 eV is close to that reported for oxide-ion conductors with an oxygen vacancy-hopping mechanism [55].

D(EF(,*(

Figure 8. Oxygen flux density vs P

from 650 to 800 oC.


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Figure 9. Partial oxide-ion conductivity σ' vs temperature at P = 0.21 atm. The oxygen vacancy diffusion coefficient Dn can be further calculated using the Nernst–Einstein equation [56], in which the oxide-ion conductivity from oxygen permeation, oxygen nonstoichiometry from TGA and unit cell volume from VTND refinement are used: σ' =

o pq rq

st =

lg

[ku∙∙ ] kv

(12)

(13)

where Cn is the oxygen vacancy concentration; VF is the molar volume of the oxide, and [V+∙∙ ] equals oxygen non-stoichiometry δ assuming that all oxygen vacancies are mobile at elevated temperatures. The oxide-ion transport in perovskite oxides is generally considered to occur via a vacancy mechanism. Therefore, combining eqs. (12) and (13) gives wt =

N4' lgkv o x


(14)


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The oxygen-ion self-diffusion coefficient D+ is related to Dn by wy =

pq

pu

x

wt = x wt

(15)

where sy is the concentration of oxide-ion. Combining eqs. (14) and (15) leads to wy =

N4' lgkv o

3 − z

(16)

The calculations of VF and δ can be found in the SI. The temperature-dependent Dn and D+ , along with Cn and C+ in a temperature range of 650 to 800 oC is shown in Figure 10. The decreases of C+ with T arises from both oxygen loss and lattice expansion (larger molar volume). However, oxygen loss plays a more pronounced role since Cn increases monotonically with T. Both Dn and D+ shown in Figure 11 follow the Arrhenius relationship with the following analytical expressions: wt = 7.14 × 10 × exp

"."he.ei jk

(17)

"."dhe.ei jk

(18)

wy = 2.42 × 10 × exp

lg

lg

The activation energy Ea for D+ is 1.18 eV, which follows the trend suggested by previous reports that Ea varies within 3.2 - 1.5 eV with x varying from 0 to 0.5, respectively, for La1xSrxCoO3-δ

systems [57, 58]. The D+ values at 800 oC obtained in this work are also compared

with others in Table 4. For a similar A-site doping sample Sr0.9Ce0.1CoO3-δ, the obtained D+ by this work is comparable, e.g., 7.98×10-8 vs 6 ×10-8 cm2 s-1. However, it is 1-2 orders of magnitude lower than other cubic structured perovskites listed. This observation also supports the later conclusion (Section 3.6) that the rate-limiting step for ORR at SYC10 is oxide-ion 22 ACS Paragon Plus Environment

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migration from cathode into electrolyte at 700-800 oC because of relatively lower oxide-ion conductivity and diffusivity.

Figure 10. (a) Dn , Cn and (b) D+ , C+ dependence on temperature.

Figure 11. Arrhenius plot of Dn and D+ . Table 4. Comparison of D+ of this work with other perovskite cathodes in the literature. Composition

Do (cm2 s-1) at 800 oC 23 ACS Paragon Plus Environment

Ref.


La0.35Sr0.65Co0.7Fe0.3O3-δ (LSCF) SrFe0.33Co0.67O3-δ (SFC) Ba0.5Sr0.5Co0.8Fe0.2O3-δ (BSCF) Sr0.9Ce0.1CoO3-δ (SCC)

9.0×10-7

[59]

5.2×10-6

[60]

2.98×10-6

[61]

6.0×10-8

[62]

7.98× ×10

Sr0.9Y0.1CoO3-δ (SYC10)

Page 24 of 36

-8

This work

3.6 Cathode polarization behavior The P -dependence of the area-specific polarization resistance (Rp) was measured to understand the polarization mechanisms. The relation between Rp and P in an oxygen reduction reaction (ORR) generally follows [63, 64]: R D ∝ P F

(19)

with m values giving the information on the rate-limiting steps: (1) m=1, O g → O, ,I+ED*J+ , oxygen gas adsorbs on the cathode surface. (2) m=0.5, O, ,I+ED*J+ → 2O,I+ED*J+ , adsorbed oxygen gas dissociates into oxygen atoms. (3) m=0.25, O,I+ED*J+ + 2e + V+∙∙ → O× + , adsorbed oxygen atoms is reduced into oxide-ion by electrons and vacancies into lattice oxygen, which is also known as charge-transfer process. (4) m=0, O (-(*E+I( → O(-(*E+-*( , oxide-ion migrates from cathode into electrolyte.

The Nyquist plots measured at 650 oC from P = 0.01 to 1 atm are shown in Figure 12 as an example; the plots at other temperatures are similar, but with a different resistance range. The plots consist of two semicircles, one at high-frequency (Rp, high) and another at low-frequency (Rp, low).

Figure 13a shows a very weak dependence of Rp, high on P , i.e. m=0.0008-0.03, suggesting 24 ACS Paragon Plus Environment

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the high-frequency arcs correspond well to step (4). For the low-frequency arc, Figure 13b, the m values increase with temperature from 0.36 at 600 oC to 0.97 at 800 oC, suggesting that the ratelimiting step changes from charge-transfer (step (3)) to surface dissociation/adsorption (step (1) and (2)). To calculate the activation energy of SYC10 as a cathode, the EIS of symmetrical cells was measured as a function of temperature in air; the results are shown in Figure 14. The Rp is 0.24 Ω cm2 at 700 oC, which is almost twice the reported values of SrCo0.9Nb0.1O3-δ and SrCo0.9Ta0.1O3-δ [54]. This observation appears to be consistent with the lower oxygen permeation flux of SYC10 than those B-site doped SCO counterparts mentioned above. It also suggests that partial oxideion conductivity indeed plays an important role in determining RP.

Figure 12. Impedance spectra of SYC10 symmetrical cell under different P at 650 oC.



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Figure 13. P -dependence of polarization resistances. (a) high-frequency resistance and (b) low-frequency resistance. The low-frequency process has a higher activation energy than the high-frequency one (2.33 eV vs 1.16 eV), which suggests that the former may be related to a surface limiting process. The overall Rp is mostly dominated by Rp, low at 600-650 oC, but by Rp, high at 700-800 oC. Based on the mechanisms outlined above, it appears that, at low temperatures, ORR kinetics is limited by a slower charge-transfer (dissociative adsorption and electron transfer) process, and at high temperatures

the

limiting

ORR

step

transitions

to

cathode/electrolyte interface.


oxide-ion

transfer

across

the

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Figure 14. Arrhenius plots of Rp, Rp, high and Rp, low at P = 0.21 atm from 600 oC to 800 oC.

IV.

Conclusions

In summary, the crystal structure of Sr0.9Y0.1CoO3-δ (SYC10) is characterized by variable temperature neutron diffraction to be P4/mmm tetragonal from RT to 900 oC. The degree of tetragonal distortion is affected by two competing processes: high temperatures promoting a more symmetric structure and oxygen intake/loss promoting lattice distortion. The SYC10 structure contains alternately arranged oxygen-deficient Co2 layers coordinating with O2 and O3, where V+∙∙ exists and randomly distributes, and almost oxygen-saturated Co1 layer coordinating with O1 and O2. The average oxidation state of Co-ions in SYC10 is very close to or smaller than +3.0, while ptype conduction is observed in a temperature range of 500-800 oC and Po2 = 0.05 - 1 atm. The BVS calculations indicate a mixed oxidation state of Co3+/4+ on Co1-site and Co2+/3+ on Co2-site. 27 ACS Paragon Plus Environment


Hence, the charge ordering could render SYC10 to exhibit either p- or n-type electronic conduction, depending on the actual T and Po2 condition. The partial oxide-ion conductivity and diffusivity of the tetragonal structured SYC10 measured from oxygen permeation are 2.3 × 10-2 S cm-1 and 7.98 × 10-8 cm2/s at 800 oC, respectively, which is 1-2 orders of magnitude lower than cubic structured LSCF, SFC and BSCF. The lower oxide-ion conductivity appears to be the reason for the higher Rp exhibited by SYC10 in comparison to Nb- and Ta-doped SCO. The P -dependence of polarization resistances shows that ORR kinetics is primarily limited by surface-controlled charge-transfer process at low temperatures (600-650 oC) and oxide-ion migration from cathode into electrolyte at high temperatures (700-800 oC). ASSOCIATED CONTENT: Supporting Information Available: **[Oxygen permeation flux test setup schematic and additional data.]**

V.

Acknowledgements

This work was funded by the Advanced Research Projects Agency-Energy (ARPA-E), U.S. Department of Energy, under award number DE-AR0000492 and Office of Fossil Energy, U.S. Department of Energy, under award number DE-FE-0023317, and National Science Foundation, under award number of CBET-1464112. Neutron scattering was carried out at the Spallation Neutron Source (SNS) at Oak Ridge National Laboratory, which is one of user facilities sponsored by the Scientific User Facilities Division, Office of Basic Energy Sciences, U.S. 28 ACS Paragon Plus Environment

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Department of Energy. The authors thank Mr. M.J. Frost from SNS for the technical support of the neutron experiment. We also thank Libin Lei for helping TGA measurements.

NOTICE OF COPYRIGHT This manuscript has been authored by UT-Battelle, LLC under Contract No. DE-AC0500OR22725 with the U.S. Department of Energy. The United States Government retains and the publisher, by accepting the article for publication, acknowledges that the United States Government retains a non-exclusive, paid-up, irrevocable, worldwide license to publish or reproduce the published form of this manuscript, or allow others to do so, for United States Government purposes. The Department of Energy will provide public access to these results of federally

sponsored

research

in

accordance

with

the

DOE

Public

Access

Plan

(http://energy.gov/downloads/doe-public-access-plan).

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