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Mixed ionic electronic conducting perovskite anode for direct carbon fuel cells
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Mixed ionic electronic conducting perovskite anode for direct carbon fuel cells A. Kulkarni a, F.T. Ciacchi a, S. Giddey a, C. Munnings a, S.P.S. Badwal a,*, J.A. Kimpton b, D. Fini a a b
CSIRO Energy Technology, Private Bag 33, Clayton South, Victoria 3169, Australia Australian Synchrotron, 800 Blackburn Road, Clayton 3168, Victoria, Australia
article info
abstract
Article history:
The conversion of carbonaceous materials to electricity in a Direct Carbon Fuel Cell (DCFC)
Received 9 August 2012
offers the most efficient process with theoretical electric efficiency close to 100%. One of
Received in revised form
the key issues for fuel cells is the continuous availability of the fuel at the triple phase
21 September 2012
boundaries between fuel, electrode and electrolyte. While this can be easily achieved with
Accepted 23 September 2012
the use of a porous fuel electrode (anode) in the case of gaseous fuels, there are serious
Available online 18 October 2012
challenges for the delivery of solid fuels to the triple junctions. In this paper, a novel concept of using mixed ionic electronic conductors (MIEC) as anode materials for DCFCs
Keywords:
has been discussed. The lanthanum strontium cobalt ferrite, La0.6Sr0.4Co0.2Fe0.8O3d (LSCF)
Fuel cell
was chosen as the first generation anode material due to its well known high mixed ionic
DCFC
and electronic conductivities in air. This material has been investigated in detail with
Carbon
respect to its conductivity, phase and microstructural stability in DCFC operating envi-
MIEC anode
ronments. When used both as the anode and cathode in a DCFC, power densities in excess
LSCF
of 50 mW/cm2 were obtained at 804 C in electrolyte supported small button cells with solid
Synchrotron XRD
carbon as the fuel. The concept of using the same anode and cathode material has also been evaluated in electrolyte supported thick wall tubular cells where power densities around 25 mW/cm2 were obtained with carbon fuel at 820 C in the presence of helium as the purging gas. The concept of using a mixed ionic electronic conducting anode for a solid fuel, to extend the reaction zone for carbon oxidation from anode/electrolyte interface to anode/solid fuel interface, has been demonstrated. Copyright ª 2012, Hydrogen Energy Publications, LLC. Published by Elsevier Ltd. All rights reserved.
1.
Introduction
Fuel cell technology offers an efficient way to generate electricity via electrochemical oxidation of a fuel. Hydrogen is widely considered as a primary fuel for various fuel cells. However, the lack of hydrogen infrastructure and energy intensive hydrogen production processes are a barrier to wide scale commercialisation of hydrogen fuel cells. The direct
carbon fuel cell (DCFC), although in its early stages of development, is a unique concept which utilises abundantly available solid carbonaceous fuels such as coal for electricity generation. With near 100% thermodynamic efficiency of the carbon oxidation reaction, an electrical efficiency of up to 70% has been projected for practical DCFC systems [1]. This will have the beneficial effect of substantially reducing the greenhouse gas emissions and also extending the life of coal
* Corresponding author. E-mail address: [email protected] (S.P.S. Badwal). 0360-3199/$ e see front matter Copyright ª 2012, Hydrogen Energy Publications, LLC. Published by Elsevier Ltd. All rights reserved. http://dx.doi.org/10.1016/j.ijhydene.2012.09.141
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reserves. Moreover, the DCFC exhaust could potentially contain an almost pure stream of carbon dioxide ready for capture and subsequent sequestration leading to substantial cost and energy savings. Coal is the most widely used fuel for electricity generation accounting for nearly 40% of electricity generation worldwide and is well poised to retain this position for many more decades to come. Furthermore, carbonaceous fuels generated from biomass sources can be used with near net zero carbon foot print. Various DCFC designs have been proposed by research groups and companies across the globe. The detailed description of various DCFC designs, fuels used and technology status have been presented in various reviews [1,2]. Briefly, three types of DCFC systems are under consideration, differentiated by the type of electrolyte used: solid oxide, molten carbonate or molten hydroxide. Depending on the fuel feed method to the anode, the solid oxide electrolyte based designs are further classified into those that either use molten metal or molten carbonate as the media, and those that use a fluidised bed to deliver carbon to the anode. The designs based on molten media systems can lead to corrosion of various cell components such as electrode, electrolyte, current collectors, and containment materials in addition to other technical issues such as a low degree of wetting of electrolyte and fuel by the molten media. The designs based on conventional SOFC are comparatively robust as no corrosive media is involved. However, the major technical issue is the difficulty in getting the solid fuel to the anode reaction zone or triple phase boundaries (TPBs) through the porous anodes. In the case of the gaseous fuel, fuel can be easily delivered through the porous structure of the anode to reach the fuel/anode/electrolyte triple junction reaction sites for its oxidation with migrating oxygen-ions and electron transfer to the anode. However, for a solid fuel, continuous fuel delivery to reaction sites is a serious challenge. There are literature reports which indicate power densities in the range of 40e300 mW/cm2 are possible by using a solid carbon pellet as the consumable anode or the traditional Ni-YSZ anode with solid carbon fuels [1,3e5]. However, power densities in the upper range are attributed to the generation of CO in-situ in the anode chamber via the reverse Boudouard reaction and subsequent oxidation of CO at TPBs. While the in-situ gasification driven DCFCs are under investigation, the direct electrochemical oxidation of carbon is always sought, as it is the most efficient way of converting carbon into electricity, given that the thermodynamic efficiency for direct electrochemical oxidation is near 100% as compared to 70% in the case of CO oxidation. The overall chemical and electrochemical reactions in direct and indirect carbon oxidation in DCFCs are given by: Overall (direct): C þ O2 / CO2
has been proposed. The use of a MIEC material with sufficient ionic and electronic conducting properties would facilitate the direct electrochemical reaction on the surface of the anode itself as shown in Fig. 1 without the requirement of diffusion of carbon through the porous electrode to TPBs at the electrode/electrolyte interface. While facilitating the direct electrochemical oxidation, the MIEC anode with controlled porosity would also allow the electrochemical oxidation of CO (generated as a result of in-situ reverse Boudouard reaction or auto-thermal gasification). A number of MIEC anode materials such as doped titanates have been investigated as anode materials for gaseous and liquid hydrocarbon fuelled SOFC applications. However, these lack the required electrocatalytic properties as compared to Ni-YSZ anodes and act as predominantly electronic conductors in the anode environments [6]. In the present work, the performance of lanthanum strontium cobalt ferrite, La0.6Sr0.4Co0.2Fe0.8O3d (LSCF) has been evaluated as a first generation mixed conducting DCFC anode material. LSCF offers excellent electronic and ionic conductivities in air, and is a well known cathode material for SOFCs. All materials in the La1-xSrxCo1-yFeyO3d family exhibit high electronic conduction, however, due to the presence of large concentration of oxygen vacancies also provide fast bulk oxygen ion transport. Furthermore, the higher electronic conduction found in these materials leads to better surface catalytic properties [7]. LSCF type materials indeed have been used as oxidation catalysts for carbon and also in gaseous hydrocarbon fuel cells which have shown to have good performance [8e11]. Since the DCFC anode environments are not as reducing as those in SOFCs operating on hydrogen, such materials are expected to show reasonable stability and conductivity. Furthermore, there are benefits of using the same MIEC anode and cathode material in terms of the ease of fabrication of the fuel cell device. In this study, the structural stability of LSCF in direct carbon fuel cell operating environments has been investigated using Synchrotron X-ray powder diffraction and scanning electron microscopy. The electrical conductivities of LSCF have been determined in air, N2, CO2, CO and H2 atmospheres to establish the base line. Finally, the electrochemical performance of the LSCF as the anode material in DCFC button as well as tubular cells with solid carbon as the fuel has been reported.
2.
Experimental procedures
Commercial LSCF powder with the composition La0.6Sr0.4 Co0.2Fe0.8O3d was used in this investigation (LSCF-HP, Fuel Cell Materials Inc., OH, USA).
(1)
(Cathode: O2 þ 4e / 2O2, Anode: C þ 2O2 / CO2 þ 4e) (2) Overall (indirect): C þ CO2 / 2CO 2CO þ O2 / 2CO2
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(3)
(Cathode: O2 þ 4e / 2O2, Anode: 2CO þ 2O2 / 2CO2 þ 4e) In the present work a novel concept of using mixed ionic and electronic conductors (MIEC) as anode materials for DCFC
Extended Reaction Zone on Surface of Anode
MIEC Anode Electrolyte
Fig. 1 e A schematic representation of the surface reaction zone in MIEC materials as the DCFC anode.
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2.1.
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Characterisation
The microstructures of sintered (1300 C, air) and heat treated (for 50 h in CO or hydrogen at 800 C) samples were analysed using Hitachi S-900 FESEM. Structural analysis was carried out with Bruker advanced D8 X-ray diffractometer using Cu Ka radiation (5e130 angular range with 0.02 step size and a 1.6 s resonance time) and Synchrotron X-ray powder diffraction of as-received powder, sintered (1300 and 1500 C) as well as 1500 C sintered and heat treated (800 C, 50 h) specimens in CO or hydrogen (Table 1). Synchrotron X-ray powder diffraction data were collected at ambient temperature over the angular A at the range 5 e85 using X-rays of wavelength 0.82552 powder diffraction beamline (10BM-1) at the Australian Synchrotron facility [12]. NIST SRM 660b (LaB6) was used to accurately determine the wavelength. The samples were packed in 0.3 mm borosilicate capillaries and were rotated during measurements to improve the powder averaging. The diffraction pattern for each sample was investigated for phase purity and, where necessary, the structures were refined using TOPAS version 4.2 software. Peak shapes were modelled using a fundamental parameters approach and unit cell size and strain parameters were determined using Le Bail refinement. ICP chemical analysis was carried out using a Varian 730-ES. Simultaneous ICP-OES with the axially-viewed plasma was used for the analysis. A CCD detector was used to monitor the emission lines with the plasma being generated via a 40 MHz free running RF generator. The following emission wave lengths were used for the analysis: La - 333.749 nm, Sr 215.283 nm, Co - 237.863 nm, Fe - 259.837 nm and Zr 343.823 nm. The sample was prepared in duplicate by acid digestion using a Milestone Start D Microwave digestion system with Aqua Regia (3:1 HCI/HNO3) for 30 min at 200 C. After the reaction had ceased and the sample cooled down, the contents of the digestion vessels were made up to volume with Milli-Q (High purity) water and analysed by ICP-OES. Certified multi-element solutions were used to check the accuracy of the calibration and the method.
2.2.
Electrical conductivity
For electrical conductivity measurements the powder was isostatically pressed into rectangular bars and sintered at 1300 C (15 h) or 1500 C (2 h) in air. The bars were ground on all sides and a standard four-probe DC technique was used for conductivity measurements of each sintered sample with platinum wire potential probes wrapped around the bar
approximately 10 mm apart and with platinum paste (Englehard 6082) current collectors at both ends of the bar. A high precision electrometer (Keithley 617) was used for voltage measurements and a Keithley 224 current source was used for supplying current. For both 1300 and 1500 C sintered specimens, the conductivity was measured at 800 C as a function of time in five different atmospheres namely air, N2, CO2, H2 and CO at 800 C for 50 h in each gas. All gases used were industrial grade (BOC, Australia) except for CO2 which was food grade (BOC, Australia). The sequence followed consisted of heating the specimen to 800 C in air and making conductivity measurements in this atmosphere for a period of 50 h, changing the gas atmosphere in sequence to N2, CO2, H2, N2 and air. For the 1300 C sintered specimen, the sequence followed was initial measurements in air at 800 C for 50 h, followed by a rapid flushing of the furnace chamber and gas lines with N2 and then measurements in CO for 50 h.
2.3.
Fuel cell testing
For evaluation of LSCF in DCFCs, the electrolyte supported symmetrical button cells were fabricated using 8 mol% Y2O3eZrO2 (Tosoh, 8YSZ) discs. These were prepared by dry pressing and sintering at 1500 C for 2 h resulting in discs with nominal thickness of 270 mm and density over 99% of the theoretical. The LSCF anode and cathode were both screen printed on the electrolyte disc followed by heat treatment at 800 C in air. A carbon black Vulcan XC-72 (Cabot Corporation, USA) paste was prepared using butyl acetate. This paste was coated on the LSCF anode and dried at 100 C. The detailed DCFC testing procedure has been reported elsewhere [13]. In summary, an in-house made test fixture with spring loaded Pt mesh current collectors was used for cell testing. A commercially available sealant (Ceramabond 552 Aremco, USA) was used as a seal to separate fuel and air chambers, and ultra high purity helium gas (BOC, Australia) was used as fuel chamber purge gas. The evaluation of fuel cells was carried out at 800e820 C. The electrochemical measurements were carried out on fuel cells using Versastat 4 (PAR, USA) for voltage - current characterization and IM6e Impedance analyser (Zahner, Germany) for impedance spectroscopy (EIS). The EIS spectra were obtained at OCV and at 20 mA/cm2 current density. In order to deconvolute the anodic and cathodic polarisation losses and electrolyte resistance in the fuel cell, three-electrode cells with a Pt ring reference electrode were prepared and
Table 1 e Comparison of lattice parameter of the major phase La0.6Sr0.4Co0.2Fe0.8O3Ld in samples examined in this study with previously reported data [15]. The last digit in the lattice parameters with uncertainty is shown in brackets. Rwp X-ray source Sample Lattice parameter ( A) Volume ( A3)
As-received 1300 C, sintered in air 1500 C, sintered in air CO e heat treated Kostogloudis and Ftikos
a
c
5.5070(7) 5.517(0) 5.5221(1) 5.5421(6) 5.4919
13.3929(5) 13.410(5) 13.4058(2) 13.623(2) 13.364
351.76(2) 353.49(6) 354.04(2) 362.35(2) 349.1
4.80 3.94 4.59 4.54
Synchrotron Bruker Synchrotron Synchrotron [15]
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subjected to galvanostatic current interruption (GCI) under multiple currents at 804 C. The details of GCI measurement procedures and equipment used are presented in [14]. Briefly the GCI technique enables separation of electrolyte resistance and polarisation losses if a fast current interrupt switch is used and the voltage decay after current interruption is recorded using a fast data acquisition system. In the present work, it involved passing a constant current between working and counter electrodes and logging voltage e time data at 500 kHz sampling rate and multiple currents by sequencing current values from 0 to a preset value, back to zero and then to the next current value, automatically through interfacing software. For the determination of electrolyte resistance, two electrode arrangement was used i.e. working and counter electrodes were used for passing current as well as recording voltage. For the measurement of individual electrode polarisation losses, a three electrode arrangement with reference electrode was used, with current passing through the working and counter electrodes, and voltage being recorded between the reference and the working electrode. The electrolyte and cathode material (LSCF), and their fabrication methods for GCI experiments were identical to those used for DCFC button cells as described at the beginning of this section. The scaled-up version of DCFC with MIEC electrodes was demonstrated with tubular cells. Closed-end electrolyte tubes were fabricated using 3 mol% Y2O3 e ZrO2 (Tosoh, TZ3Y) by dry bag isostatic pressing and sintering at 1500 C for 2 h. The sintered tube dimensions were 6.42 mm OD 5.57 mm ID (wall thickness 0.425 mm) with density of over 99% of the theoretical value. The length of the tube employed for these investigations was 145 mm. Both LSCF anode and cathode were prepared by brush coating the LSCF ink both on the outside and inside the tube on a 40 mm long portion of the tube (5 mm before the closed-end of the tube). This was followed by heat treatment at 800 C in air. Pt wires with Pt paste were used as current collectors on both sides. This was again followed by heat treatment at 800 C in air. In the final step, the Vulcan XC-72 was applied on the anode as a fuel in the form of a paste made with butyl acetate. The evaluation of the tubular cell was carried out at 820 C in a stainless steel reactor. Fig. 2 shows the schematic of the reactor design. It consists of a 35 mm ID S.S. crucible with a flange. A brass flange (internally water cooled) with a number of ports for gases and electrical connections serves as a cover of the reactor as shown in the Fig. 2. After the assembly of the cell in the S.S. crucible it is inserted into a heated vertical chamber of a furnace up to its flange. The complete set-up resides inside a test station that allows the operation of the cell up to 1000 C, and has provision to supply gases such as H2, CO, N2, He and air to the inner or outer chamber of the tube. The station has multiple levels of safety redundancy for the operator and the cell, and functions in a failsafe mode. The cell is heated to the operating temperature with helium gas flowing in the outer chamber and air in the inner chamber of the tube. Once the temperature as measured by the thermocouple (placed inside the tube close to the active area) stabilises at the set temperature, voltageecurrent characteristics are obtained in different gaseous atmospheres as required.
Air in Air out Anode chamber gas in Cooling Jacket head Tubular cell
Anode chamber gas out Cathode current collector Anode current collector Carbon paste S.S. crucible
Thermocouple Fig. 2 e A schematic drawing of the test fixture used for electrochemical testing of tubular DCFCs.
3.
Results and discussion
3.1. Structure and stability of LSCF in DCFC operating environments X-ray diffraction studies were carried out on as-received powder and samples which underwent a number of different heat treatments, using both a Bruker (laboratory diffractometer) and Synchrotron X-ray sources. The Synchrotron X-ray diffraction patterns for the LSCF specimens sintered in air at 1500 C, and those 1500 C sintered and then heat treated (50 h at 800 C) either in CO or H2 are shown in Figs. 3e5 respectively with a summary of the lattice parameters provided in Table 1. The XRD pattern of the asreceived powder was similar to that of the 1500 C sintered specimen. Chemical analysis was also performed using ICP on the as-received powder. The chemical composition is provided in Table 2. Only elements detected above 0.1% were considered significant for this work with Cu, Cr, Mn and Na detected during analysis at levels below 0.1%. All samples, with the exception of the hydrogen treated material, which decomposed, were found to contain one dominant rhombohedral (R-3c) phase. Structurally this is generally in good agreement with previous studies of LSCF materials [15,16]. The chemical composition of the material was found to be deficient in lanthanum (Table 2). The deficiency in lanthanum appears to be compensated by the precipitation of a minor secondary cobalt oxide (CoO) phase which could be detected in the Synchrotron X-ray diffraction patterns of as-received and 1500 C sintered specimens (Fig. 3 (b)). It is worth noting here that this phase was not detected by Bruker X-ray diffractometer due to the use of Cu Ka radiation. The fluorescence of iron and cobalt in Cu Ka radiation make lab-based x-ray diffraction insensitive to iron or cobalt impurities unless they are present in large volumes (10%þ). A comparison of the lattice parameters observed in this study to previously reported work on the La0.6Sr0.4Fe1-xCoxO3-d series suggests that
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a
60,000
200,000
30,000
40,000
20,000
Counts
Counts
150,000 0
100,000
21
21.2
20,000
10,000
50,000
* *
*
*
0
0
10 10
15
20
25
30
35
40
45
50
15
20
25
30
35
40
45
50
2θ, Degrees
2θ, Degrees
b
* - CoO o - unknown phase
Ln (counts)
10.5 10.0 9.5 9.0 8.5
o
*
*
8.0 19
20
21
22
Fig. 4 e Synchrotron X-ray powder diffraction data for the LSCF sample heat treated in CO for 50 h at 800 C and refined with a Le Bail model (Space Group R-3c, a [ 5.5421(6), c [ 13.623(2), Rwp [ 4.54). The red line shows the Le Bail model data with the grey line as the difference curve shown directly below the main pattern. Note the shoulder on the low angle side of the 104 and 110 doublet indicate an unknown third phase that possibly may be LaSrFeO4. (*) indicate peaks associated with Co metal. (For interpretation of the references to colour in this figure legend, the reader is referred to the web version of this article.)
23
2θ, Degrees Fig. 3 e (a) Synchrotron X-ray powder diffraction data for the LSCF sample sintered at 1500 C in air and refined with a Le Bail model (Space Group R-3c, a [ 5.5221(1), c [ 13.4058(2), Rwp [ 4.59). The red line shows the Le Bail model data with the grey line as the difference curve shown directly below the main pattern. (Note: The poor fit of the 202 and 006 peaks in the inset indicate possible composition variations or additional strain contributions). (b) Expanded view of the XRD pattern from Fig. 3(a) showing CoO and an unknown phase peaks. For clarity figure shows only experimental data, unmarked peaks correspond to rhombohedral (R-3c) phase. (For interpretation of the references to colour in this figure legend, the reader is referred to the web version of this article.)
the major phase examined in this work contains significantly less Co at the B-site than indicated by the composition of asreceived powder [15e17]. Peaks for a third unknown phase (only trace levels) were also observed within Synchrotron X-ray diffraction patterns for the 1500 C air sintered sample at w19 (as obvious in Fig. 3(b)) and also at w24.6 . Due to the very low intensity and diffuse nature of these peaks it was not possible to identify this phase. Both CoO and the trace levels of unknown phases were estimated to account for substantially less than 5 vol% of the total composition. Comparison between diffraction pattern data for the as-received sample and those sintered in air at 1300 and 1500 C showed little difference structurally. The as-received powder had a slightly smaller unit cell than the 1300 and 1500 C sintered samples (Table 1). This is most
likely related to the thermal history of the sample which has been shown to have an effect on the lattice parameters and properties of these and related materials [17]. Impurity levels of Zr were detected with ICP analysis but no phase associated with Zr was observed in either the X-ray diffraction or SEM measurements carried out within this study. Attempts to carry out a full refinement of the structure including the determination of site occupancies and cation ratios were unsuccessful due to the high absorption of X-rays by the material leading to variation in peak intensity. It is theoretically possible to model the absorption effects and isolate them from the structural model but the increased error
Fig. 5 e Synchrotron X-ray powder diffraction pattern of the LSCF sample heat treated for 50 h in H2 at 800 C. Note the complex phase assemblage due to the complete breakdown of the LSCF structure: La2O3 (,); Fe (-); Fe0.75Co0.25 (:); SrO ( 3 ); LaSrFeO4 (A); La(OH)3 (y); Co (B).
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Table 2 e Chemical composition of material determined via ICP. Atomic fractions at the A- and B- sublattices exclude Zr impurity. Element
La Sr Co Fe Zr
Wt %
Relative standard deviation (%)
Actual atomic fractions at the A- and B- sublattices
Target atomic fractions at the A- and B- sublattices
35.4 16.5 5.5 21.0 0.4
1.4 2.1 2.1 1.9 6.3
0.53 0.39 0.20 0.80
0.60 0.40 0.20 0.80
involved with this would reduce the value of the results. Furthermore, closer inspection of the peak shapes indicated poor fits in some cases as highlighted in the inset from Fig. 3 (a) where the 202 and 006 peaks show additional broadening that may be attributed to a number of factors including subtle composition variations or additional strain contributions. This would again increase the complexity and error associated with full structural modelling. Peak shape modelling indicates strain within the unit cell (w3.4%) for the 1500 C air sintered sample is relatively low. Since lattice parameters of the LSCF family of materials are affected by La/Sr ratio at the A-site and Co/Fe ratio at the B-site as well as by the heat treatment at higher temperatures, it was not possible to determine the exact composition. However, based on lattice parameters from XRD data, La deficiency at the A-site from chemical composition determination and precipitation of CoO phase, it was estimated that the Co atomic percentage at the B-site is around 12e13%. Fig. 4 shows the Synchrotron X-ray powder diffraction pattern for LSCF heat treated at 800 C in CO for 50 h. The primary phase remains rhombohedral (R-3c) with lattice parameters: a ¼ 5.5421(6), c ¼ 13.623(2), Rwp ¼ 4.54. The expansion in the lattice (see Table 1) is due to the reduction of the transition metal cations which leads to an increase in their ionic radius. The primary phase has an overall relatively lower scattered intensity and significant peak broadening compared to the 1500 C sintered sample. Peak shape modelling for the primary phase shows that most of the broadening can be attributed to an increase in lattice strain, which increased to w53% in the CO heat treated sample from w3.4% for the 1500 C sintered sample. This strain is likely to be the result of lattice defects. This damage to the structure would be consistent with the loss of significant volumes of oxygen (forming oxygen lattice vacancies) caused by the reduction of the transition metal ions in the CO atmosphere. The secondary phases observed in this sample include a metallic Co phase (<5%) and an unknown trace phase that forms a lowangle shoulder on the 104 and 110 doublet (Fig. 4 inset). This unknown phase (<5%) may be tentatively assigned as LaSrFeO4 which has been shown to be stable in reducing environments [18,19]. The majority of the free Co metal is likely to be due to the reduction of the free cobalt oxide that was observed in the sample sintered at 1500 C in air (Fig. 3 (b)) and also in the as-received powder. From the Co metal peak intensities it is difficult to articulate whether any additional Co ions are removed from the lattice with exposure to CO. This is due to the relatively poor scattering from the primary phase in comparison to the 1500 C sintered sample.
On treatment in hydrogen the major perovskite phase completely decomposes (Fig. 5). The decomposition products are La2O3, La(OH)3, SrO, Fe, Co and Fe0.75Co0.25. Thus, the LSCF phase although showing stability in CO is unstable at 800 C over an extended period of treatment in hydrogen. In addition to the structural characterisation, microstructural analysis was also carried out. Fig. 6 compares the SEM micrographs of LSCF samples treated in CO, H2 and air for 50 h. The microstructure of 1300 C sintered specimen observed in Fig. 6 (a) is fairly homogeneous and showed no sign of phase separation at this magnification. This suggests that any secondary phase is finely dispersed and that it cannot be observed on this scale. After heat treating in CO there were clearly a number of small rounded nodules (Fig. 6 (b)) which form on the surface of the oxide. This is consistent with the observation of a small amount of cobalt metal being formed on exposure to CO as generally metals tend to coarsen or agglomerate once formed on the surface of a dense oxide material. After exposure to hydrogen, the surface is markedly different (Fig. 6 (c, d)) to that of the as-sintered or CO heat treated samples with the entire surface being comprised of a number of discrete grains or particles with a range of sizes and morphologies. Again this is consistent with a material which has decomposed into a range of different phases. Previously we had reported that no structural changes occurred to LSCF when placed in direct contact with carbon black for 2 h at 800 C in a helium atmosphere [13]. In this work, longer heat treatment times in CO of 50 h have been used. The high resolution Synchrotron data have shown that there are other minor phases present in the material examined but that overall LSCF is stable in DCFC operating conditions and therefore can be considered a potential anode material for use in DCFCs.
3.2.
Electrical conductivity measurements
The sintered specimen for conductivity measurements had a density of 6.23 (>99% of theoretical) and 5.94 g/cm3 (w95.0% of theoretical) respectively for the sintering temperatures of 1300 and 1500 C. The theoretical density used did not account for the small amount of impurity phase present, however, it is unlikely that a small amount of the impurity phase would have a measureable effect on the density measurement. Note the density of the specimen sintered at a higher temperature but for a much shorter period of time, was somewhat lower. Fig. 7 shows Arrhenius plots for the conductivity of both 1300 and 1500 C sintered specimens in air. The conductivity increased with increasing temperature with a broad maxima
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Fig. 6 e Scanning electron micrographs of (a) La0.6Sr0.4Co0.2Fe0.8O3Lx after sintering at 1300 C (15 h), HR&CR: 300 C/hr, (b) after conductivity measurements in CO at 800 C, (c, d) after conductivity measurements at 800 C in air, N2, CO2 and H2.
at around 550e600 C, and then decreased with further increase in the temperature. The conductivity of the 1500 C sintered specimen was somewhat lower as was its density. The general trend and the observed behaviour are similar to those reported in the literature [15,18,20]. In these materials, the conductivity is controlled by a small polaron hopping model with charge transfer being thermally activated. However, since there is some loss of oxygen with increasing temperature (which could vary with the specimen density and heating rate) and consequent reduction in the concentration of p-type charge carriers, the actual conductivity values and their behaviour as a function of temperature is somewhat complex. In fact Zeng et al. [17] have reported that the conductivity e temperature behaviour is significantly affected
2.60 AIR
Log (σ, S cm-1)
o
1300 C
2.50
2.40 1500oC
2.30
2.20 0.9
1.0
1.1
1.2
1.3
1.4
1.5
1000/T, K-1 Fig. 7 e Arrhenius plots of LSCF conductivity in air for 1300 and 1500 C sintered specimens.
by the sintering temperature of the nominal composition used in this study. Fig. 8 (a) and 8 (b) show the conductivity of 1300 and 1500 C sintered specimens respectively in various gas atmospheres as a function of time at a constant measurement temperature of 800 C. There was very little change in the conductivity of both specimens with time in air. On changing the gas atmosphere to nitrogen, the conductivity decreased rapidly over the first 500 min and then more slowly over the next 1500 min and then settled to a constant value. Before changing to CO2 atmosphere there was a short purge with air. The behaviour in CO2 was similar to that in N2, except that some fluctuations were clearly obvious in the CO2 atmosphere. On changing the gas atmosphere from CO2 to hydrogen (1300 and 1500 C sintered specimens) or from nitrogen to CO (1300 C sintered specimen), the conductivity decreased from 70 e 80 S-cm1 to around 1 S-cm1 in about 20 min and then over the next 4880 min, there was very little change with time. As obvious in the inset of Fig. 8 (a), the conductivity of LSCF phase assemblage in CO was about twice that in hydrogen (0.96 vs 0.46 S-cm1). Since there is complete decomposition of the LSCF phase following treatment in hydrogen at 800 C, the conductivity data for the hydrogen atmosphere has little significance. Teraoka et al. measured the ionic conductivity of various LSCF compositions in oxygen, air and He [21]. They report that the ionic conductivity is most affected by the A-site composition (substitution of La by Sr) and to a lesser extent by the B-site composition. For the composition used in this investigation, these authors report that the ionic conductivity increases with decreasing oxygen partial pressure and in He it was greater than 0.1Scm1 at 800 C At this temperature, it is
a
300
Conductivity, S cm-1
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250
800oC Air
1.5 CO
1.0
200
0.5
150
Hydrogen
0.0 0
1000
2000
3000
4000
5000
100
Carbon Dioxide
50
Nitrogen CO
0
Hydrogen
0
1000
2000
3000
4000
5000
b
300
Conductivity, S cm-1
Time, min 250
o
800 C Air
200 150 100
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discussion that there is a complex relationship between conductivity, oxygen activity, phase assemblage and the fuel cell operating temperature. However, it appears that, despite a substantial decrease in the conductivity, the LSCF phase is reasonably stable as discussed above where only minute quantities of other phases were detected in the anode environments of a DCFC operating around 800 C or below and is likely to remain stable over extended period of operation. The conductivity is a very sensitive parameter to composition of the phase assemblage and as mentioned above, after the first few minutes of changing gas atmosphere to CO, there was very little change in the conductivity over w5000 min of measurements. XRD work was also performed on specimens which had been heat treated for 50 h in CO. Thus it is reasonable to consider that observed phase assemblage would not change further with time. Based on the discussion above, the substantial decrease in the total conductivity of LSCF at 800 C in CO over that in air can be attributed to the electronic conductivity component due to a decrease in the concentration of p-type charge carriers.
Carbon Dioxide
50 0
0
1000
2000
3000
Nitrogen
3.3.
Hydrogen
Fig. 9 shows values of data obtained from GCI at 804 C under various currents. The electrolyte resistance was calculated using the two electrode arrangement while single electrode polarisation resistance was obtained using three electrode configurations. While the voltage drop across the electrolyte increases linearly with current, the cathode losses decrease slightly under higher currents as expected. The conductivity of the electrolyte was found to be 0.048 S-cm1 (after correction for the resistance of Pt lead wires at 800 C), and agrees well with literature values for YSZ electrolyte [22]. The cathode polarisation resistance was found to be 0.07 U-cm2 which is consistent with the literature values for LSCF electrodes at 800 C [23]. The values of cathode polarisation and the electrolyte resistance were used to calculate the anodic polarisation from the total cell losses at any particular current as discussed later in this section.
4000
5000
Time, min Fig. 8 e Conductivity of LSCF specimens sintered at 1300 C (a) and 1500 C (b) respectively in various gas atmospheres as a function of time at the constant measurement temperature of 800 C. The inset in Fig. 8(a) is expanded view of H2 and CO data.
twice that of 8 mol% Y2O3 e ZrO2 (8YSZ), a composition with the highest ionic conductivity in the Y2O3 e ZrO2 system [22]. Considering that the oxygen partial pressure is much lower in CO compared to He and that the main LSCF phase is stable, it is not unreasonable to assume that the ionic conductivity of the LSCF anode would be significantly higher than that of 8YSZ. Tai et al. [18] report that the LSCF composition used in this investigation, exhibits p-type conductivity at pO2 > 1013 atm at 1000 C. The conductivity decreases with decreasing oxygen partial pressure due to a loss in the concentration of p-type charge carriers. Below this oxygen concentration, there is a transition from p- to n-type conductivity. However, around this oxygen concentration other phases such as Co and (LaSr)(CoFe)O4 as discussed above, start appearing. The oxygen concentration, at which p- to n-type transition occurs and where the other phases start appearing, is temperature dependent and is considerably lower as the measurement temperature decreases. According to Tai et al. [18], the oxygen activity, down to which LSCF as the major phase, is observed, is 1010 at 1200 C and 1016 atm at 1000 C. It is expected that the LSCF phase will be observed up to considerably lower oxygen partial pressures at 800 C, the typical operating temperature of DCFC and at which most of the measurements reported here were carried out. It is obvious from the above
DCFC test results and discussion
T = 804oC
Electrolyte
Cathode
Fig. 9 e The cathodic polarisation and the electrolyte resistance as determined by GCI technique as a function of current at 804 C. Electrolyte thickness: 0.27 mm, electrode area: 0.5 cm2.
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Fig. 10 shows the OCV of the DCFC button cell as a function of temperature. The OCV showed a steep increase when cell was heated from 500 C to 600 C and after that it increased gradually to reach 0.99 V which is close to the theoretical value for electrochemical oxidation of carbon (Reaction 1). However, the OCV data alone is not sufficient to shed light on the dominant reaction mechanism and more investigation would be required for better understanding of the reaction mechanism. Fig. 11 shows the currentevoltageepower (IeVeP) curve for DCFC button cell at 804 C. The peak power densities over 50 mW/cm2 were achieved after correction for IR loss due to platinum lead wires. This is despite the fact that the kinetics of oxidation of solid carbon is slower than gaseous fuels such as H2. Since only ultra-high purity helium was used as the anode chamber purge gas, the generation of CO due to direct reaction between carbon and oxygen in the purge gas (maximum 2 ppm specified by BOC) would not be sufficient to drive the reaction. Thus, the power generation can be attributed to direct electrochemical oxidation of carbon in contact with the anode surface. The tail of IeV curve indicates the onset of limiting current behaviour possibly due to reduction of the number of reactive sites on the anode caused by the loss of contact between the carbon particles and the anode. The carbon can electrochemically oxidise to CO (2 electron transfer) or to CO2 (4 electron transfer) at the anode/carbon interface. Formation of a small amount of CO via the reverse Boudouard reaction between carbon and CO2 generated in-situ cannot be ruled out. The CO generated by the electrochemical oxidation of carbon or the reverse Boudouard reaction can then either oxidise at the anode/carbon interface or diffuses through the anode to oxidise at the anode/electrolyte interface. The unique ability of MIEC anodes such as LSCF is that they can facilitate the direct electrochemical oxidation of carbon on the surface as well as utilise CO which may be generated as a secondary product of electrochemical oxidation. It is worth mentioning here that although the electronic conductivity of the LSCF anode in DCFC operating conditions is considerably lower than that of the traditional SOFC anode (Ni-YSZ), the electrochemical performance is impressive and substantially higher than that reported for Ni/YSZ electrode when it was used as an anode in direct carbon fuel cells with
Fig. 11 e Voltageecurrent density and voltage-power density characteristics of the DCFC button cell tested at 804 C with Vulcan XC-72 carbon fuel. LSCF is used as anode as well as cathode. The inset is a pie chart showing the distribution of anodic (ha), cathodic (hc) and ohmic (IR) voltage losses at the peak power density.
solid carbon fuel [24,25]. Porous Ni/YSZ composites are commonly used as the anode in solid oxide fuel cells and give excellent performance with gaseous fuels, however, due to the difficulty of solid fuel diffusing through the porous structure, it does not appear to be as good an anode for solid carbon fuel. This may result from the lower ionic conductivity of YSZ as discussed above. The use of a mixed ionic/electronic electrode such as LSCF eliminates the need for solid carbon fuel to diffuse to electrode/electrolyte interface and the reasonable power output from button cells with LSCF anodes clearly demonstrates the feasibility of using such materials as an anode for DCFCs. Gorte et al. have indicated that the lower conductivity of the anode could be compensated by proper distribution of a current collector over the functional anode layer [26]. Alternatively, the electronic conductivity of the anode can be enhanced by the incorporation of a conducting metal or ceramic phase in the anode. The incorporation of the conducting phase within the anode has been reported to give better performance than the use of a pure metallic conducting current collector (such as silver) [27]. Inset of Fig. 11 shows a pie chart of ohmic (mainly electrolyte), anode and cathode voltage losses at the peak power density. As can be seen from
Carbon/Air He purge gas in fuel chamber -3
Z'', Ω-cm2
T = 804oC
-2
-1
0 0
1
2
3
4
5
6
7
Z', Ω-cm2 Fig. 10 e Open circuit voltage (OCV) of the DCFC button cell with Vulcan XC-72 carbon fuel as a function of temperature. LSCF is used as anode as well as cathode.
Fig. 12 e EIS spectrum of the DCFC button cell at 804 C at open circuit voltage (OCV) with Vulcan XC-72 carbon fuel. LSCF is used as anode as well as cathode.
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the pie chart, both anodic losses (70%) and electrolyte losses (28%) dominate overall cell voltage losses in the button cell DCFC at 804 C. The voltage losses associated with the electrolyte can easily be reduced with the use of state-of-the-art thin electrolyte films. The high anode polarisation can primarily be attributed to the slower rate of electrochemical oxidation of carbon and the low lateral conductivity of the anode which as mentioned above, can be improved with the incorporation of conducting metal (Cu, Ag or Ni, etc.) or an inert conducting ceramic phase in the anode. In the DCFC configuration used in this work with carbon layer pressed on the anode it is likely that the electrically conductive carbon black offers additional conduction paths for electrons in addition to the direct contact between the anode and the current collector. Fig. 12 shows the EIS plot under OCV condition and under a load of 20 mA/cm2. The values of IR losses including that from Pt lead wire of 1.82 U agree well with the GCI data. For a current density of 20 mA/ cm2, the total polarisation resistance shown by the lower frequency arc from the EIS spectrum (5.45 U-cm2) agrees well with the polarisation loss calculated from IeV curve (5.5 U-cm2) by subtracting electrolyte resistance losses and output voltage of the cell from OCV. As the button cells with an LSCF anode demonstrated promising performance, the concept was scaled-up using the tubular cell as discussed in the experimental section. Fig. 13 shows voltageecurrent characteristics of the tubular cell at 820 C, with carbon as a fuel and air as an oxidant, and helium gas flowing in the anode chamber of the cell. The cell produced an OCV of 1.11 V, and a peak power density of 25 mW/cm2 at a current density of 43 mA/cm2. The power densities obtained here are low compared to those achieved for button cells. Possible reasons for this could be: (a) the composition of the electrolyte material is 3 mol% yttria - zirconia with its conductivity one third that of 8 mol% yttria stabilised zirconia used for button cells; (b) the electrolyte thickness is higher in the case of the tubular cell (0.425 mm versus 0.27 mm for button cells) and (c) the LSCF ink has been brush coated in the case of the tubular cell compared to screen printing used for button cells. The total power produced by the tubular cell was w175 mW. The power output can be significantly
Fig. 13 e Voltageecurrent density and voltage-power density characteristics of the DCFC tubular cell tested at 820 C with Vulcan XC-72 carbon. LSCF was used as the anode as well as the cathode.
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improved with the use of a thinner 8 mol% yttria stabilised zirconia electrolyte, incorporation of a more conducting phase to enhance electronic conductivity in the anode and the use of a more optimised electrode coating techniques such as spray coating or dip coating onto a cathode supported tubular structure [28].
4.
Conclusion
LSCF, a commonly used perovskite air electrode, has been evaluated as the fuel electrode in a direct carbon fuel cell. The material used was deficient at the A-site sublattice. A significant observation was the detection, by high resolution Synchrotron XRD, of the CoO phase (B-site sublattice material), both in the as-received powder and sintered (in air) LSCF bars, which otherwise was not detected using a laboratory X-ray source using Cu Ka radiation. The total conductivity of LSCF in CO is around 1 S-cm1, however, the ionic conductivity appears to be twice than that of YSZ used in traditional Ni/YSZ anodes. The rhombohedral LSCF phase remained the most dominant phase even after 50 h of heat treatment in CO at the typical DCFC operating temperature of 800 C. Noteworthy was the finding that the total conductivity of LSCF did not change much after about 20 min of heat treatment in CO at 800 C up to further 4880 min of testing. The conductivity data clearly demonstrated the structural stability of LSCF in DCFC operating environments consistent with phase analysis by high resolution XRD. Although the electrical conductivity of LSCF in reducing environments is significantly lower than that of traditionally used Ni based SOFC anodes, power densities exceeding 50 mW/cm2 were achieved in a electrolyte supported button cell DCFC with LSCF as both the anode and the cathode. This power density is substantially higher than that reported for Ni/YSZ anode in a DCFC. This is most probably due to significantly higher ionic conductivity of LSCF compared with YSZ leading to lower voltage losses when carbon electrochemical oxidation reaction takes place at the two phase carbon/anode interface. The polarisation losses, when deconvoluted using GCI and EIS techniques, showed that the anodic polarisation was still quite high and along with resistive losses across the electrolyte were mainly responsible for the lower power densities in direct carbon contact DCFCs in this study. This was attributed to the low electronic conductivity of LSCF in reducing environments. Furthermore, a scaled up version of DCFC with LSCF anode, operating on solid carbon fuel, was demonstrated with tubular design electrolyte supported cell which delivered a total power up to 175 mW and a peak power density of 25 mW/cm2. Overall, mixed ionic electronic conducting LSCF material showed promising performance as the DCFC anode with solid carbon fuel, despite its low electronic conductivity which can be improved further with the addition of a metal or conducting inert ceramic phase to enhance electronic conductivity and optimisation of the anode microstructure. Thus the concept of using mixed ionic electronic conducting anode for a solid fuel to extend the reaction zone for carbon oxidation from anode/ electrolyte interface to anode/solid fuel interface for continuous operation of the direct carbon fuel cell has been clearly demonstrated. This type of fuel cell and operation mode
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would offer the highest efficiency for direct conversion of carbon from coal or biomass to electricity provided the anode material has both high ionic and electronic conductivities.
Acknowledgements Authors would like to thank Dr. David Hay for reviewing this manuscript. The work was carried out in CSIRO’s Advanced Coal Technology Portfolio.
references
[1] Giddey S, Badwal SPS, Kulkarni A, Munnings C. A comprehensive review of direct carbon fuel cell technology. Prog Ener Combust Sci 2012;38:360e99. [2] Cao D, Sun Y, Wang G. Direct carbon fuel cell: fundamentals and recent developments. J Power Sources 2007;167:250e7. [3] Li S, Lee A, Mitchell R, Gu¨r T. Conversion of solid carbonaceous fuels in a fluidized bed fuel cell. Electrochem Solid-state Lett 2008;11:B20e3. [4] Lee A, Mitchell R, Gu¨r T. Thermodynamic analysis of gasification-driven direct carbon fuel cells. J Power Sources 2009;194:774e85. [5] Li C, Shi Y, Cai N. Performance improvement of direct carbon fuel cell by introducing catalytic gasification process. J Power Sources 2010;195:4660e6. [6] Atkinson A, Barnett S, Gorte RJ, Irvine JTS, McEvoy AJ, Mogensen M, et al. Advanced anodes for high-temperature fuel cells. Nat Mat 2004;3:17e27. [7] Adler S. Factors governing oxygen reduction in solid oxide fuel cell cathodes. Chem Rev 2004;104:4791e843. [8] Fisher II JC, Chuang SC. Investigating the CH4 reaction pathway on a novel LSCF anode catalyst in the SOFC. Catal Comm 2009;10(6):772e6. [9] Hartley A, Sahibzada M, Weston M, Metcalfe IS, Mantzavinos D. La0.6Sr0.4Co0.2Fe0.8O3 as the anode and cathode for intermediate temperature solid oxide fuel cells. Catal Today 2000;55:197e204. [10] Huang T, Chou C, Chen W, Huang M. Coal syngas reactivity over Ni-added LSCFeGDC anode of solid oxide fuel cells. Electrochem Comm 2009;11:294e7. [11] Lai B, Kerman K, Ramanathan S. Nanostructured La0.6Sr0.4Co0.2Fe0.8O3/Y0.08Zr0.92O0.196/La0.6Sr0.4Co0.2Fe0.8O3 (LSCF/YSZ/LSCF) symmetric thin film solid oxide fuel cells. J Power Sources 2011;196(4):1826e32. [12] Wallwork KS, Kennedy BJ, Wang D. The high resolution powder diffraction beamline for the Australian synchrotron. AIP Conf Proc 2007;879:879e82.
[13] Kulkarni A, Giddey S, Badwal SPS. Electrochemical performance of ceria-gadolinia electrolyte based direct carbon fuel cells. Solid State Ion 2011;194(1):46e52. [14] Clarke R, Kulkarni A, Badwal SPS, Giddey S. Evaluation of solid electrolyte cells with a versatile electrochemical technique. Ionics 2012. http://dx.doi.org/10.1007/s11581-012-0749-1. [15] Kostogloudis GC, Ftikos Ch. Properties of A-site-deficient La0.6Sr0.4Co0.2Fe0.8O3d - based perovskite oxides. Solid State Ion 1999;126:143e51. [16] Hashimoto S, Fukuda Y, Kuhn M, Sato K, Yashiro K, Mizusaki J. Thermal and chemical lattice expansibility of La0.6Sr0.4Co1yFeyO3d (y ¼ 0.2, 0.4, 0.6 and 0.8). J Solid State Ion 2011;186:37e43. [17] Zeng P, Ran R, Chen Z, Gu H, Shao Z, Diniz da Costa JC, et al. Significant effects of sintering temperature on the performance of La0.6Sr0.4Co0.2Fe0.8O3d oxygen selective membranes. J Membr Sci 2007;302:171e9. [18] Tai LW, Nasrallah MM, Anderson HU. Thermochemical stability, electrical conductivity, and Seebeck coefficient of Sr-doped LaCo0.2Fe0.8O3d. J Solid State Chem 1995;118: 117e24. [19] Jennings AJ, Skinner SJ, Helgason O. Structural properties of LaxSr2xFeO4þ/ d at high temperature and under reducing conditions. J Solid State Chem 2003;175:207e17. [20] Tai LW, Nasrallah MM, Anderson HU, Sparlin DM, Sehlin SR. Oxygen stoichiometries in La1xSrxCo1yFeyO3d perovskites at reduced oxygen partial pressures. Solid State Ion 1995;76: 273e83. [21] Teraoka Y, Zhang HM, Okamoto K, Yamazoe N. Mixed ionic and electronic conductivity of LalxSrxColyFeyO3d perovskite type oxides. Mat Res Bull 1988;23:51e8. [22] Badwal SPS, Ciacchi FT. Oxygen-ion conducting electrolyte materials for solid oxide fuel cells. Ionics 2000;6:1e21. [23] Esquirol A, Brandon NP, Kilner JA, Mogensen M. Electrochemical characterization of La0.6Sr0.4Co0.2Fe0.8O3 cathodes for intermediate-temperature SOFCs. J Electrochem Soc 2004;151:A1847e55. [24] Li S, Lee AC, Mitchell RE, Gu¨r TM. Direct carbon conversion in a helium fluidized bed fuel cell. Solid State Ion 2008;179: 1549e52. [25] Li C, Shi Y, Cai N. Effect of contact type between anode and carbonaceous fuels on direct carbon fuel cell reaction characteristics. J Power Sources 2011;196:4588e93. [26] Gross MD, Vohs JM, Gorte RJ. A strategy for achieving high performance with SOFC ceramic anodes. Electrochem Solidstate Lett 2007;10:B65e9. [27] Tang Y, Liu J. Effect of anode and Boudouard reaction catalysts on the performance of direct carbon solid oxide fuel cells. Int J Hydrogen Energy 2010;35:11188e93. [28] Zhou J, Ye XF, Shao L, Zhang XP, Qian JQ, Wang SR. A promising direct carbon fuel cell based on the cathodesupported tubular solid oxide fuel cell technology. Electrochim Acta 2012;74:267e70.
Available online at www.sciencedirect.com
journal homepage: www.elsevier.com/locate/he
Mixed ionic electronic conducting perovskite anode for direct carbon fuel cells A. Kulkarni a, F.T. Ciacchi a, S. Giddey a, C. Munnings a, S.P.S. Badwal a,*, J.A. Kimpton b, D. Fini a a b
CSIRO Energy Technology, Private Bag 33, Clayton South, Victoria 3169, Australia Australian Synchrotron, 800 Blackburn Road, Clayton 3168, Victoria, Australia
article info
abstract
Article history:
The conversion of carbonaceous materials to electricity in a Direct Carbon Fuel Cell (DCFC)
Received 9 August 2012
offers the most efficient process with theoretical electric efficiency close to 100%. One of
Received in revised form
the key issues for fuel cells is the continuous availability of the fuel at the triple phase
21 September 2012
boundaries between fuel, electrode and electrolyte. While this can be easily achieved with
Accepted 23 September 2012
the use of a porous fuel electrode (anode) in the case of gaseous fuels, there are serious
Available online 18 October 2012
challenges for the delivery of solid fuels to the triple junctions. In this paper, a novel concept of using mixed ionic electronic conductors (MIEC) as anode materials for DCFCs
Keywords:
has been discussed. The lanthanum strontium cobalt ferrite, La0.6Sr0.4Co0.2Fe0.8O3d (LSCF)
Fuel cell
was chosen as the first generation anode material due to its well known high mixed ionic
DCFC
and electronic conductivities in air. This material has been investigated in detail with
Carbon
respect to its conductivity, phase and microstructural stability in DCFC operating envi-
MIEC anode
ronments. When used both as the anode and cathode in a DCFC, power densities in excess
LSCF
of 50 mW/cm2 were obtained at 804 C in electrolyte supported small button cells with solid
Synchrotron XRD
carbon as the fuel. The concept of using the same anode and cathode material has also been evaluated in electrolyte supported thick wall tubular cells where power densities around 25 mW/cm2 were obtained with carbon fuel at 820 C in the presence of helium as the purging gas. The concept of using a mixed ionic electronic conducting anode for a solid fuel, to extend the reaction zone for carbon oxidation from anode/electrolyte interface to anode/solid fuel interface, has been demonstrated. Copyright ª 2012, Hydrogen Energy Publications, LLC. Published by Elsevier Ltd. All rights reserved.
1.
Introduction
Fuel cell technology offers an efficient way to generate electricity via electrochemical oxidation of a fuel. Hydrogen is widely considered as a primary fuel for various fuel cells. However, the lack of hydrogen infrastructure and energy intensive hydrogen production processes are a barrier to wide scale commercialisation of hydrogen fuel cells. The direct
carbon fuel cell (DCFC), although in its early stages of development, is a unique concept which utilises abundantly available solid carbonaceous fuels such as coal for electricity generation. With near 100% thermodynamic efficiency of the carbon oxidation reaction, an electrical efficiency of up to 70% has been projected for practical DCFC systems [1]. This will have the beneficial effect of substantially reducing the greenhouse gas emissions and also extending the life of coal
* Corresponding author. E-mail address: [email protected] (S.P.S. Badwal). 0360-3199/$ e see front matter Copyright ª 2012, Hydrogen Energy Publications, LLC. Published by Elsevier Ltd. All rights reserved. http://dx.doi.org/10.1016/j.ijhydene.2012.09.141
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reserves. Moreover, the DCFC exhaust could potentially contain an almost pure stream of carbon dioxide ready for capture and subsequent sequestration leading to substantial cost and energy savings. Coal is the most widely used fuel for electricity generation accounting for nearly 40% of electricity generation worldwide and is well poised to retain this position for many more decades to come. Furthermore, carbonaceous fuels generated from biomass sources can be used with near net zero carbon foot print. Various DCFC designs have been proposed by research groups and companies across the globe. The detailed description of various DCFC designs, fuels used and technology status have been presented in various reviews [1,2]. Briefly, three types of DCFC systems are under consideration, differentiated by the type of electrolyte used: solid oxide, molten carbonate or molten hydroxide. Depending on the fuel feed method to the anode, the solid oxide electrolyte based designs are further classified into those that either use molten metal or molten carbonate as the media, and those that use a fluidised bed to deliver carbon to the anode. The designs based on molten media systems can lead to corrosion of various cell components such as electrode, electrolyte, current collectors, and containment materials in addition to other technical issues such as a low degree of wetting of electrolyte and fuel by the molten media. The designs based on conventional SOFC are comparatively robust as no corrosive media is involved. However, the major technical issue is the difficulty in getting the solid fuel to the anode reaction zone or triple phase boundaries (TPBs) through the porous anodes. In the case of the gaseous fuel, fuel can be easily delivered through the porous structure of the anode to reach the fuel/anode/electrolyte triple junction reaction sites for its oxidation with migrating oxygen-ions and electron transfer to the anode. However, for a solid fuel, continuous fuel delivery to reaction sites is a serious challenge. There are literature reports which indicate power densities in the range of 40e300 mW/cm2 are possible by using a solid carbon pellet as the consumable anode or the traditional Ni-YSZ anode with solid carbon fuels [1,3e5]. However, power densities in the upper range are attributed to the generation of CO in-situ in the anode chamber via the reverse Boudouard reaction and subsequent oxidation of CO at TPBs. While the in-situ gasification driven DCFCs are under investigation, the direct electrochemical oxidation of carbon is always sought, as it is the most efficient way of converting carbon into electricity, given that the thermodynamic efficiency for direct electrochemical oxidation is near 100% as compared to 70% in the case of CO oxidation. The overall chemical and electrochemical reactions in direct and indirect carbon oxidation in DCFCs are given by: Overall (direct): C þ O2 / CO2
has been proposed. The use of a MIEC material with sufficient ionic and electronic conducting properties would facilitate the direct electrochemical reaction on the surface of the anode itself as shown in Fig. 1 without the requirement of diffusion of carbon through the porous electrode to TPBs at the electrode/electrolyte interface. While facilitating the direct electrochemical oxidation, the MIEC anode with controlled porosity would also allow the electrochemical oxidation of CO (generated as a result of in-situ reverse Boudouard reaction or auto-thermal gasification). A number of MIEC anode materials such as doped titanates have been investigated as anode materials for gaseous and liquid hydrocarbon fuelled SOFC applications. However, these lack the required electrocatalytic properties as compared to Ni-YSZ anodes and act as predominantly electronic conductors in the anode environments [6]. In the present work, the performance of lanthanum strontium cobalt ferrite, La0.6Sr0.4Co0.2Fe0.8O3d (LSCF) has been evaluated as a first generation mixed conducting DCFC anode material. LSCF offers excellent electronic and ionic conductivities in air, and is a well known cathode material for SOFCs. All materials in the La1-xSrxCo1-yFeyO3d family exhibit high electronic conduction, however, due to the presence of large concentration of oxygen vacancies also provide fast bulk oxygen ion transport. Furthermore, the higher electronic conduction found in these materials leads to better surface catalytic properties [7]. LSCF type materials indeed have been used as oxidation catalysts for carbon and also in gaseous hydrocarbon fuel cells which have shown to have good performance [8e11]. Since the DCFC anode environments are not as reducing as those in SOFCs operating on hydrogen, such materials are expected to show reasonable stability and conductivity. Furthermore, there are benefits of using the same MIEC anode and cathode material in terms of the ease of fabrication of the fuel cell device. In this study, the structural stability of LSCF in direct carbon fuel cell operating environments has been investigated using Synchrotron X-ray powder diffraction and scanning electron microscopy. The electrical conductivities of LSCF have been determined in air, N2, CO2, CO and H2 atmospheres to establish the base line. Finally, the electrochemical performance of the LSCF as the anode material in DCFC button as well as tubular cells with solid carbon as the fuel has been reported.
2.
Experimental procedures
Commercial LSCF powder with the composition La0.6Sr0.4 Co0.2Fe0.8O3d was used in this investigation (LSCF-HP, Fuel Cell Materials Inc., OH, USA).
(1)
(Cathode: O2 þ 4e / 2O2, Anode: C þ 2O2 / CO2 þ 4e) (2) Overall (indirect): C þ CO2 / 2CO 2CO þ O2 / 2CO2
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(3)
(Cathode: O2 þ 4e / 2O2, Anode: 2CO þ 2O2 / 2CO2 þ 4e) In the present work a novel concept of using mixed ionic and electronic conductors (MIEC) as anode materials for DCFC
Extended Reaction Zone on Surface of Anode
MIEC Anode Electrolyte
Fig. 1 e A schematic representation of the surface reaction zone in MIEC materials as the DCFC anode.
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2.1.
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Characterisation
The microstructures of sintered (1300 C, air) and heat treated (for 50 h in CO or hydrogen at 800 C) samples were analysed using Hitachi S-900 FESEM. Structural analysis was carried out with Bruker advanced D8 X-ray diffractometer using Cu Ka radiation (5e130 angular range with 0.02 step size and a 1.6 s resonance time) and Synchrotron X-ray powder diffraction of as-received powder, sintered (1300 and 1500 C) as well as 1500 C sintered and heat treated (800 C, 50 h) specimens in CO or hydrogen (Table 1). Synchrotron X-ray powder diffraction data were collected at ambient temperature over the angular A at the range 5 e85 using X-rays of wavelength 0.82552 powder diffraction beamline (10BM-1) at the Australian Synchrotron facility [12]. NIST SRM 660b (LaB6) was used to accurately determine the wavelength. The samples were packed in 0.3 mm borosilicate capillaries and were rotated during measurements to improve the powder averaging. The diffraction pattern for each sample was investigated for phase purity and, where necessary, the structures were refined using TOPAS version 4.2 software. Peak shapes were modelled using a fundamental parameters approach and unit cell size and strain parameters were determined using Le Bail refinement. ICP chemical analysis was carried out using a Varian 730-ES. Simultaneous ICP-OES with the axially-viewed plasma was used for the analysis. A CCD detector was used to monitor the emission lines with the plasma being generated via a 40 MHz free running RF generator. The following emission wave lengths were used for the analysis: La - 333.749 nm, Sr 215.283 nm, Co - 237.863 nm, Fe - 259.837 nm and Zr 343.823 nm. The sample was prepared in duplicate by acid digestion using a Milestone Start D Microwave digestion system with Aqua Regia (3:1 HCI/HNO3) for 30 min at 200 C. After the reaction had ceased and the sample cooled down, the contents of the digestion vessels were made up to volume with Milli-Q (High purity) water and analysed by ICP-OES. Certified multi-element solutions were used to check the accuracy of the calibration and the method.
2.2.
Electrical conductivity
For electrical conductivity measurements the powder was isostatically pressed into rectangular bars and sintered at 1300 C (15 h) or 1500 C (2 h) in air. The bars were ground on all sides and a standard four-probe DC technique was used for conductivity measurements of each sintered sample with platinum wire potential probes wrapped around the bar
approximately 10 mm apart and with platinum paste (Englehard 6082) current collectors at both ends of the bar. A high precision electrometer (Keithley 617) was used for voltage measurements and a Keithley 224 current source was used for supplying current. For both 1300 and 1500 C sintered specimens, the conductivity was measured at 800 C as a function of time in five different atmospheres namely air, N2, CO2, H2 and CO at 800 C for 50 h in each gas. All gases used were industrial grade (BOC, Australia) except for CO2 which was food grade (BOC, Australia). The sequence followed consisted of heating the specimen to 800 C in air and making conductivity measurements in this atmosphere for a period of 50 h, changing the gas atmosphere in sequence to N2, CO2, H2, N2 and air. For the 1300 C sintered specimen, the sequence followed was initial measurements in air at 800 C for 50 h, followed by a rapid flushing of the furnace chamber and gas lines with N2 and then measurements in CO for 50 h.
2.3.
Fuel cell testing
For evaluation of LSCF in DCFCs, the electrolyte supported symmetrical button cells were fabricated using 8 mol% Y2O3eZrO2 (Tosoh, 8YSZ) discs. These were prepared by dry pressing and sintering at 1500 C for 2 h resulting in discs with nominal thickness of 270 mm and density over 99% of the theoretical. The LSCF anode and cathode were both screen printed on the electrolyte disc followed by heat treatment at 800 C in air. A carbon black Vulcan XC-72 (Cabot Corporation, USA) paste was prepared using butyl acetate. This paste was coated on the LSCF anode and dried at 100 C. The detailed DCFC testing procedure has been reported elsewhere [13]. In summary, an in-house made test fixture with spring loaded Pt mesh current collectors was used for cell testing. A commercially available sealant (Ceramabond 552 Aremco, USA) was used as a seal to separate fuel and air chambers, and ultra high purity helium gas (BOC, Australia) was used as fuel chamber purge gas. The evaluation of fuel cells was carried out at 800e820 C. The electrochemical measurements were carried out on fuel cells using Versastat 4 (PAR, USA) for voltage - current characterization and IM6e Impedance analyser (Zahner, Germany) for impedance spectroscopy (EIS). The EIS spectra were obtained at OCV and at 20 mA/cm2 current density. In order to deconvolute the anodic and cathodic polarisation losses and electrolyte resistance in the fuel cell, three-electrode cells with a Pt ring reference electrode were prepared and
Table 1 e Comparison of lattice parameter of the major phase La0.6Sr0.4Co0.2Fe0.8O3Ld in samples examined in this study with previously reported data [15]. The last digit in the lattice parameters with uncertainty is shown in brackets. Rwp X-ray source Sample Lattice parameter ( A) Volume ( A3)
As-received 1300 C, sintered in air 1500 C, sintered in air CO e heat treated Kostogloudis and Ftikos
a
c
5.5070(7) 5.517(0) 5.5221(1) 5.5421(6) 5.4919
13.3929(5) 13.410(5) 13.4058(2) 13.623(2) 13.364
351.76(2) 353.49(6) 354.04(2) 362.35(2) 349.1
4.80 3.94 4.59 4.54
Synchrotron Bruker Synchrotron Synchrotron [15]
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subjected to galvanostatic current interruption (GCI) under multiple currents at 804 C. The details of GCI measurement procedures and equipment used are presented in [14]. Briefly the GCI technique enables separation of electrolyte resistance and polarisation losses if a fast current interrupt switch is used and the voltage decay after current interruption is recorded using a fast data acquisition system. In the present work, it involved passing a constant current between working and counter electrodes and logging voltage e time data at 500 kHz sampling rate and multiple currents by sequencing current values from 0 to a preset value, back to zero and then to the next current value, automatically through interfacing software. For the determination of electrolyte resistance, two electrode arrangement was used i.e. working and counter electrodes were used for passing current as well as recording voltage. For the measurement of individual electrode polarisation losses, a three electrode arrangement with reference electrode was used, with current passing through the working and counter electrodes, and voltage being recorded between the reference and the working electrode. The electrolyte and cathode material (LSCF), and their fabrication methods for GCI experiments were identical to those used for DCFC button cells as described at the beginning of this section. The scaled-up version of DCFC with MIEC electrodes was demonstrated with tubular cells. Closed-end electrolyte tubes were fabricated using 3 mol% Y2O3 e ZrO2 (Tosoh, TZ3Y) by dry bag isostatic pressing and sintering at 1500 C for 2 h. The sintered tube dimensions were 6.42 mm OD 5.57 mm ID (wall thickness 0.425 mm) with density of over 99% of the theoretical value. The length of the tube employed for these investigations was 145 mm. Both LSCF anode and cathode were prepared by brush coating the LSCF ink both on the outside and inside the tube on a 40 mm long portion of the tube (5 mm before the closed-end of the tube). This was followed by heat treatment at 800 C in air. Pt wires with Pt paste were used as current collectors on both sides. This was again followed by heat treatment at 800 C in air. In the final step, the Vulcan XC-72 was applied on the anode as a fuel in the form of a paste made with butyl acetate. The evaluation of the tubular cell was carried out at 820 C in a stainless steel reactor. Fig. 2 shows the schematic of the reactor design. It consists of a 35 mm ID S.S. crucible with a flange. A brass flange (internally water cooled) with a number of ports for gases and electrical connections serves as a cover of the reactor as shown in the Fig. 2. After the assembly of the cell in the S.S. crucible it is inserted into a heated vertical chamber of a furnace up to its flange. The complete set-up resides inside a test station that allows the operation of the cell up to 1000 C, and has provision to supply gases such as H2, CO, N2, He and air to the inner or outer chamber of the tube. The station has multiple levels of safety redundancy for the operator and the cell, and functions in a failsafe mode. The cell is heated to the operating temperature with helium gas flowing in the outer chamber and air in the inner chamber of the tube. Once the temperature as measured by the thermocouple (placed inside the tube close to the active area) stabilises at the set temperature, voltageecurrent characteristics are obtained in different gaseous atmospheres as required.
Air in Air out Anode chamber gas in Cooling Jacket head Tubular cell
Anode chamber gas out Cathode current collector Anode current collector Carbon paste S.S. crucible
Thermocouple Fig. 2 e A schematic drawing of the test fixture used for electrochemical testing of tubular DCFCs.
3.
Results and discussion
3.1. Structure and stability of LSCF in DCFC operating environments X-ray diffraction studies were carried out on as-received powder and samples which underwent a number of different heat treatments, using both a Bruker (laboratory diffractometer) and Synchrotron X-ray sources. The Synchrotron X-ray diffraction patterns for the LSCF specimens sintered in air at 1500 C, and those 1500 C sintered and then heat treated (50 h at 800 C) either in CO or H2 are shown in Figs. 3e5 respectively with a summary of the lattice parameters provided in Table 1. The XRD pattern of the asreceived powder was similar to that of the 1500 C sintered specimen. Chemical analysis was also performed using ICP on the as-received powder. The chemical composition is provided in Table 2. Only elements detected above 0.1% were considered significant for this work with Cu, Cr, Mn and Na detected during analysis at levels below 0.1%. All samples, with the exception of the hydrogen treated material, which decomposed, were found to contain one dominant rhombohedral (R-3c) phase. Structurally this is generally in good agreement with previous studies of LSCF materials [15,16]. The chemical composition of the material was found to be deficient in lanthanum (Table 2). The deficiency in lanthanum appears to be compensated by the precipitation of a minor secondary cobalt oxide (CoO) phase which could be detected in the Synchrotron X-ray diffraction patterns of as-received and 1500 C sintered specimens (Fig. 3 (b)). It is worth noting here that this phase was not detected by Bruker X-ray diffractometer due to the use of Cu Ka radiation. The fluorescence of iron and cobalt in Cu Ka radiation make lab-based x-ray diffraction insensitive to iron or cobalt impurities unless they are present in large volumes (10%þ). A comparison of the lattice parameters observed in this study to previously reported work on the La0.6Sr0.4Fe1-xCoxO3-d series suggests that
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a
60,000
200,000
30,000
40,000
20,000
Counts
Counts
150,000 0
100,000
21
21.2
20,000
10,000
50,000
* *
*
*
0
0
10 10
15
20
25
30
35
40
45
50
15
20
25
30
35
40
45
50
2θ, Degrees
2θ, Degrees
b
* - CoO o - unknown phase
Ln (counts)
10.5 10.0 9.5 9.0 8.5
o
*
*
8.0 19
20
21
22
Fig. 4 e Synchrotron X-ray powder diffraction data for the LSCF sample heat treated in CO for 50 h at 800 C and refined with a Le Bail model (Space Group R-3c, a [ 5.5421(6), c [ 13.623(2), Rwp [ 4.54). The red line shows the Le Bail model data with the grey line as the difference curve shown directly below the main pattern. Note the shoulder on the low angle side of the 104 and 110 doublet indicate an unknown third phase that possibly may be LaSrFeO4. (*) indicate peaks associated with Co metal. (For interpretation of the references to colour in this figure legend, the reader is referred to the web version of this article.)
23
2θ, Degrees Fig. 3 e (a) Synchrotron X-ray powder diffraction data for the LSCF sample sintered at 1500 C in air and refined with a Le Bail model (Space Group R-3c, a [ 5.5221(1), c [ 13.4058(2), Rwp [ 4.59). The red line shows the Le Bail model data with the grey line as the difference curve shown directly below the main pattern. (Note: The poor fit of the 202 and 006 peaks in the inset indicate possible composition variations or additional strain contributions). (b) Expanded view of the XRD pattern from Fig. 3(a) showing CoO and an unknown phase peaks. For clarity figure shows only experimental data, unmarked peaks correspond to rhombohedral (R-3c) phase. (For interpretation of the references to colour in this figure legend, the reader is referred to the web version of this article.)
the major phase examined in this work contains significantly less Co at the B-site than indicated by the composition of asreceived powder [15e17]. Peaks for a third unknown phase (only trace levels) were also observed within Synchrotron X-ray diffraction patterns for the 1500 C air sintered sample at w19 (as obvious in Fig. 3(b)) and also at w24.6 . Due to the very low intensity and diffuse nature of these peaks it was not possible to identify this phase. Both CoO and the trace levels of unknown phases were estimated to account for substantially less than 5 vol% of the total composition. Comparison between diffraction pattern data for the as-received sample and those sintered in air at 1300 and 1500 C showed little difference structurally. The as-received powder had a slightly smaller unit cell than the 1300 and 1500 C sintered samples (Table 1). This is most
likely related to the thermal history of the sample which has been shown to have an effect on the lattice parameters and properties of these and related materials [17]. Impurity levels of Zr were detected with ICP analysis but no phase associated with Zr was observed in either the X-ray diffraction or SEM measurements carried out within this study. Attempts to carry out a full refinement of the structure including the determination of site occupancies and cation ratios were unsuccessful due to the high absorption of X-rays by the material leading to variation in peak intensity. It is theoretically possible to model the absorption effects and isolate them from the structural model but the increased error
Fig. 5 e Synchrotron X-ray powder diffraction pattern of the LSCF sample heat treated for 50 h in H2 at 800 C. Note the complex phase assemblage due to the complete breakdown of the LSCF structure: La2O3 (,); Fe (-); Fe0.75Co0.25 (:); SrO ( 3 ); LaSrFeO4 (A); La(OH)3 (y); Co (B).
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Table 2 e Chemical composition of material determined via ICP. Atomic fractions at the A- and B- sublattices exclude Zr impurity. Element
La Sr Co Fe Zr
Wt %
Relative standard deviation (%)
Actual atomic fractions at the A- and B- sublattices
Target atomic fractions at the A- and B- sublattices
35.4 16.5 5.5 21.0 0.4
1.4 2.1 2.1 1.9 6.3
0.53 0.39 0.20 0.80
0.60 0.40 0.20 0.80
involved with this would reduce the value of the results. Furthermore, closer inspection of the peak shapes indicated poor fits in some cases as highlighted in the inset from Fig. 3 (a) where the 202 and 006 peaks show additional broadening that may be attributed to a number of factors including subtle composition variations or additional strain contributions. This would again increase the complexity and error associated with full structural modelling. Peak shape modelling indicates strain within the unit cell (w3.4%) for the 1500 C air sintered sample is relatively low. Since lattice parameters of the LSCF family of materials are affected by La/Sr ratio at the A-site and Co/Fe ratio at the B-site as well as by the heat treatment at higher temperatures, it was not possible to determine the exact composition. However, based on lattice parameters from XRD data, La deficiency at the A-site from chemical composition determination and precipitation of CoO phase, it was estimated that the Co atomic percentage at the B-site is around 12e13%. Fig. 4 shows the Synchrotron X-ray powder diffraction pattern for LSCF heat treated at 800 C in CO for 50 h. The primary phase remains rhombohedral (R-3c) with lattice parameters: a ¼ 5.5421(6), c ¼ 13.623(2), Rwp ¼ 4.54. The expansion in the lattice (see Table 1) is due to the reduction of the transition metal cations which leads to an increase in their ionic radius. The primary phase has an overall relatively lower scattered intensity and significant peak broadening compared to the 1500 C sintered sample. Peak shape modelling for the primary phase shows that most of the broadening can be attributed to an increase in lattice strain, which increased to w53% in the CO heat treated sample from w3.4% for the 1500 C sintered sample. This strain is likely to be the result of lattice defects. This damage to the structure would be consistent with the loss of significant volumes of oxygen (forming oxygen lattice vacancies) caused by the reduction of the transition metal ions in the CO atmosphere. The secondary phases observed in this sample include a metallic Co phase (<5%) and an unknown trace phase that forms a lowangle shoulder on the 104 and 110 doublet (Fig. 4 inset). This unknown phase (<5%) may be tentatively assigned as LaSrFeO4 which has been shown to be stable in reducing environments [18,19]. The majority of the free Co metal is likely to be due to the reduction of the free cobalt oxide that was observed in the sample sintered at 1500 C in air (Fig. 3 (b)) and also in the as-received powder. From the Co metal peak intensities it is difficult to articulate whether any additional Co ions are removed from the lattice with exposure to CO. This is due to the relatively poor scattering from the primary phase in comparison to the 1500 C sintered sample.
On treatment in hydrogen the major perovskite phase completely decomposes (Fig. 5). The decomposition products are La2O3, La(OH)3, SrO, Fe, Co and Fe0.75Co0.25. Thus, the LSCF phase although showing stability in CO is unstable at 800 C over an extended period of treatment in hydrogen. In addition to the structural characterisation, microstructural analysis was also carried out. Fig. 6 compares the SEM micrographs of LSCF samples treated in CO, H2 and air for 50 h. The microstructure of 1300 C sintered specimen observed in Fig. 6 (a) is fairly homogeneous and showed no sign of phase separation at this magnification. This suggests that any secondary phase is finely dispersed and that it cannot be observed on this scale. After heat treating in CO there were clearly a number of small rounded nodules (Fig. 6 (b)) which form on the surface of the oxide. This is consistent with the observation of a small amount of cobalt metal being formed on exposure to CO as generally metals tend to coarsen or agglomerate once formed on the surface of a dense oxide material. After exposure to hydrogen, the surface is markedly different (Fig. 6 (c, d)) to that of the as-sintered or CO heat treated samples with the entire surface being comprised of a number of discrete grains or particles with a range of sizes and morphologies. Again this is consistent with a material which has decomposed into a range of different phases. Previously we had reported that no structural changes occurred to LSCF when placed in direct contact with carbon black for 2 h at 800 C in a helium atmosphere [13]. In this work, longer heat treatment times in CO of 50 h have been used. The high resolution Synchrotron data have shown that there are other minor phases present in the material examined but that overall LSCF is stable in DCFC operating conditions and therefore can be considered a potential anode material for use in DCFCs.
3.2.
Electrical conductivity measurements
The sintered specimen for conductivity measurements had a density of 6.23 (>99% of theoretical) and 5.94 g/cm3 (w95.0% of theoretical) respectively for the sintering temperatures of 1300 and 1500 C. The theoretical density used did not account for the small amount of impurity phase present, however, it is unlikely that a small amount of the impurity phase would have a measureable effect on the density measurement. Note the density of the specimen sintered at a higher temperature but for a much shorter period of time, was somewhat lower. Fig. 7 shows Arrhenius plots for the conductivity of both 1300 and 1500 C sintered specimens in air. The conductivity increased with increasing temperature with a broad maxima
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Fig. 6 e Scanning electron micrographs of (a) La0.6Sr0.4Co0.2Fe0.8O3Lx after sintering at 1300 C (15 h), HR&CR: 300 C/hr, (b) after conductivity measurements in CO at 800 C, (c, d) after conductivity measurements at 800 C in air, N2, CO2 and H2.
at around 550e600 C, and then decreased with further increase in the temperature. The conductivity of the 1500 C sintered specimen was somewhat lower as was its density. The general trend and the observed behaviour are similar to those reported in the literature [15,18,20]. In these materials, the conductivity is controlled by a small polaron hopping model with charge transfer being thermally activated. However, since there is some loss of oxygen with increasing temperature (which could vary with the specimen density and heating rate) and consequent reduction in the concentration of p-type charge carriers, the actual conductivity values and their behaviour as a function of temperature is somewhat complex. In fact Zeng et al. [17] have reported that the conductivity e temperature behaviour is significantly affected
2.60 AIR
Log (σ, S cm-1)
o
1300 C
2.50
2.40 1500oC
2.30
2.20 0.9
1.0
1.1
1.2
1.3
1.4
1.5
1000/T, K-1 Fig. 7 e Arrhenius plots of LSCF conductivity in air for 1300 and 1500 C sintered specimens.
by the sintering temperature of the nominal composition used in this study. Fig. 8 (a) and 8 (b) show the conductivity of 1300 and 1500 C sintered specimens respectively in various gas atmospheres as a function of time at a constant measurement temperature of 800 C. There was very little change in the conductivity of both specimens with time in air. On changing the gas atmosphere to nitrogen, the conductivity decreased rapidly over the first 500 min and then more slowly over the next 1500 min and then settled to a constant value. Before changing to CO2 atmosphere there was a short purge with air. The behaviour in CO2 was similar to that in N2, except that some fluctuations were clearly obvious in the CO2 atmosphere. On changing the gas atmosphere from CO2 to hydrogen (1300 and 1500 C sintered specimens) or from nitrogen to CO (1300 C sintered specimen), the conductivity decreased from 70 e 80 S-cm1 to around 1 S-cm1 in about 20 min and then over the next 4880 min, there was very little change with time. As obvious in the inset of Fig. 8 (a), the conductivity of LSCF phase assemblage in CO was about twice that in hydrogen (0.96 vs 0.46 S-cm1). Since there is complete decomposition of the LSCF phase following treatment in hydrogen at 800 C, the conductivity data for the hydrogen atmosphere has little significance. Teraoka et al. measured the ionic conductivity of various LSCF compositions in oxygen, air and He [21]. They report that the ionic conductivity is most affected by the A-site composition (substitution of La by Sr) and to a lesser extent by the B-site composition. For the composition used in this investigation, these authors report that the ionic conductivity increases with decreasing oxygen partial pressure and in He it was greater than 0.1Scm1 at 800 C At this temperature, it is
a
300
Conductivity, S cm-1
i n t e r n a t i o n a l j o u r n a l o f h y d r o g e n e n e r g y 3 7 ( 2 0 1 2 ) 1 9 0 9 2 e1 9 1 0 2
250
800oC Air
1.5 CO
1.0
200
0.5
150
Hydrogen
0.0 0
1000
2000
3000
4000
5000
100
Carbon Dioxide
50
Nitrogen CO
0
Hydrogen
0
1000
2000
3000
4000
5000
b
300
Conductivity, S cm-1
Time, min 250
o
800 C Air
200 150 100
19099
discussion that there is a complex relationship between conductivity, oxygen activity, phase assemblage and the fuel cell operating temperature. However, it appears that, despite a substantial decrease in the conductivity, the LSCF phase is reasonably stable as discussed above where only minute quantities of other phases were detected in the anode environments of a DCFC operating around 800 C or below and is likely to remain stable over extended period of operation. The conductivity is a very sensitive parameter to composition of the phase assemblage and as mentioned above, after the first few minutes of changing gas atmosphere to CO, there was very little change in the conductivity over w5000 min of measurements. XRD work was also performed on specimens which had been heat treated for 50 h in CO. Thus it is reasonable to consider that observed phase assemblage would not change further with time. Based on the discussion above, the substantial decrease in the total conductivity of LSCF at 800 C in CO over that in air can be attributed to the electronic conductivity component due to a decrease in the concentration of p-type charge carriers.
Carbon Dioxide
50 0
0
1000
2000
3000
Nitrogen
3.3.
Hydrogen
Fig. 9 shows values of data obtained from GCI at 804 C under various currents. The electrolyte resistance was calculated using the two electrode arrangement while single electrode polarisation resistance was obtained using three electrode configurations. While the voltage drop across the electrolyte increases linearly with current, the cathode losses decrease slightly under higher currents as expected. The conductivity of the electrolyte was found to be 0.048 S-cm1 (after correction for the resistance of Pt lead wires at 800 C), and agrees well with literature values for YSZ electrolyte [22]. The cathode polarisation resistance was found to be 0.07 U-cm2 which is consistent with the literature values for LSCF electrodes at 800 C [23]. The values of cathode polarisation and the electrolyte resistance were used to calculate the anodic polarisation from the total cell losses at any particular current as discussed later in this section.
4000
5000
Time, min Fig. 8 e Conductivity of LSCF specimens sintered at 1300 C (a) and 1500 C (b) respectively in various gas atmospheres as a function of time at the constant measurement temperature of 800 C. The inset in Fig. 8(a) is expanded view of H2 and CO data.
twice that of 8 mol% Y2O3 e ZrO2 (8YSZ), a composition with the highest ionic conductivity in the Y2O3 e ZrO2 system [22]. Considering that the oxygen partial pressure is much lower in CO compared to He and that the main LSCF phase is stable, it is not unreasonable to assume that the ionic conductivity of the LSCF anode would be significantly higher than that of 8YSZ. Tai et al. [18] report that the LSCF composition used in this investigation, exhibits p-type conductivity at pO2 > 1013 atm at 1000 C. The conductivity decreases with decreasing oxygen partial pressure due to a loss in the concentration of p-type charge carriers. Below this oxygen concentration, there is a transition from p- to n-type conductivity. However, around this oxygen concentration other phases such as Co and (LaSr)(CoFe)O4 as discussed above, start appearing. The oxygen concentration, at which p- to n-type transition occurs and where the other phases start appearing, is temperature dependent and is considerably lower as the measurement temperature decreases. According to Tai et al. [18], the oxygen activity, down to which LSCF as the major phase, is observed, is 1010 at 1200 C and 1016 atm at 1000 C. It is expected that the LSCF phase will be observed up to considerably lower oxygen partial pressures at 800 C, the typical operating temperature of DCFC and at which most of the measurements reported here were carried out. It is obvious from the above
DCFC test results and discussion
T = 804oC
Electrolyte
Cathode
Fig. 9 e The cathodic polarisation and the electrolyte resistance as determined by GCI technique as a function of current at 804 C. Electrolyte thickness: 0.27 mm, electrode area: 0.5 cm2.
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Fig. 10 shows the OCV of the DCFC button cell as a function of temperature. The OCV showed a steep increase when cell was heated from 500 C to 600 C and after that it increased gradually to reach 0.99 V which is close to the theoretical value for electrochemical oxidation of carbon (Reaction 1). However, the OCV data alone is not sufficient to shed light on the dominant reaction mechanism and more investigation would be required for better understanding of the reaction mechanism. Fig. 11 shows the currentevoltageepower (IeVeP) curve for DCFC button cell at 804 C. The peak power densities over 50 mW/cm2 were achieved after correction for IR loss due to platinum lead wires. This is despite the fact that the kinetics of oxidation of solid carbon is slower than gaseous fuels such as H2. Since only ultra-high purity helium was used as the anode chamber purge gas, the generation of CO due to direct reaction between carbon and oxygen in the purge gas (maximum 2 ppm specified by BOC) would not be sufficient to drive the reaction. Thus, the power generation can be attributed to direct electrochemical oxidation of carbon in contact with the anode surface. The tail of IeV curve indicates the onset of limiting current behaviour possibly due to reduction of the number of reactive sites on the anode caused by the loss of contact between the carbon particles and the anode. The carbon can electrochemically oxidise to CO (2 electron transfer) or to CO2 (4 electron transfer) at the anode/carbon interface. Formation of a small amount of CO via the reverse Boudouard reaction between carbon and CO2 generated in-situ cannot be ruled out. The CO generated by the electrochemical oxidation of carbon or the reverse Boudouard reaction can then either oxidise at the anode/carbon interface or diffuses through the anode to oxidise at the anode/electrolyte interface. The unique ability of MIEC anodes such as LSCF is that they can facilitate the direct electrochemical oxidation of carbon on the surface as well as utilise CO which may be generated as a secondary product of electrochemical oxidation. It is worth mentioning here that although the electronic conductivity of the LSCF anode in DCFC operating conditions is considerably lower than that of the traditional SOFC anode (Ni-YSZ), the electrochemical performance is impressive and substantially higher than that reported for Ni/YSZ electrode when it was used as an anode in direct carbon fuel cells with
Fig. 11 e Voltageecurrent density and voltage-power density characteristics of the DCFC button cell tested at 804 C with Vulcan XC-72 carbon fuel. LSCF is used as anode as well as cathode. The inset is a pie chart showing the distribution of anodic (ha), cathodic (hc) and ohmic (IR) voltage losses at the peak power density.
solid carbon fuel [24,25]. Porous Ni/YSZ composites are commonly used as the anode in solid oxide fuel cells and give excellent performance with gaseous fuels, however, due to the difficulty of solid fuel diffusing through the porous structure, it does not appear to be as good an anode for solid carbon fuel. This may result from the lower ionic conductivity of YSZ as discussed above. The use of a mixed ionic/electronic electrode such as LSCF eliminates the need for solid carbon fuel to diffuse to electrode/electrolyte interface and the reasonable power output from button cells with LSCF anodes clearly demonstrates the feasibility of using such materials as an anode for DCFCs. Gorte et al. have indicated that the lower conductivity of the anode could be compensated by proper distribution of a current collector over the functional anode layer [26]. Alternatively, the electronic conductivity of the anode can be enhanced by the incorporation of a conducting metal or ceramic phase in the anode. The incorporation of the conducting phase within the anode has been reported to give better performance than the use of a pure metallic conducting current collector (such as silver) [27]. Inset of Fig. 11 shows a pie chart of ohmic (mainly electrolyte), anode and cathode voltage losses at the peak power density. As can be seen from
Carbon/Air He purge gas in fuel chamber -3
Z'', Ω-cm2
T = 804oC
-2
-1
0 0
1
2
3
4
5
6
7
Z', Ω-cm2 Fig. 10 e Open circuit voltage (OCV) of the DCFC button cell with Vulcan XC-72 carbon fuel as a function of temperature. LSCF is used as anode as well as cathode.
Fig. 12 e EIS spectrum of the DCFC button cell at 804 C at open circuit voltage (OCV) with Vulcan XC-72 carbon fuel. LSCF is used as anode as well as cathode.
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the pie chart, both anodic losses (70%) and electrolyte losses (28%) dominate overall cell voltage losses in the button cell DCFC at 804 C. The voltage losses associated with the electrolyte can easily be reduced with the use of state-of-the-art thin electrolyte films. The high anode polarisation can primarily be attributed to the slower rate of electrochemical oxidation of carbon and the low lateral conductivity of the anode which as mentioned above, can be improved with the incorporation of conducting metal (Cu, Ag or Ni, etc.) or an inert conducting ceramic phase in the anode. In the DCFC configuration used in this work with carbon layer pressed on the anode it is likely that the electrically conductive carbon black offers additional conduction paths for electrons in addition to the direct contact between the anode and the current collector. Fig. 12 shows the EIS plot under OCV condition and under a load of 20 mA/cm2. The values of IR losses including that from Pt lead wire of 1.82 U agree well with the GCI data. For a current density of 20 mA/ cm2, the total polarisation resistance shown by the lower frequency arc from the EIS spectrum (5.45 U-cm2) agrees well with the polarisation loss calculated from IeV curve (5.5 U-cm2) by subtracting electrolyte resistance losses and output voltage of the cell from OCV. As the button cells with an LSCF anode demonstrated promising performance, the concept was scaled-up using the tubular cell as discussed in the experimental section. Fig. 13 shows voltageecurrent characteristics of the tubular cell at 820 C, with carbon as a fuel and air as an oxidant, and helium gas flowing in the anode chamber of the cell. The cell produced an OCV of 1.11 V, and a peak power density of 25 mW/cm2 at a current density of 43 mA/cm2. The power densities obtained here are low compared to those achieved for button cells. Possible reasons for this could be: (a) the composition of the electrolyte material is 3 mol% yttria - zirconia with its conductivity one third that of 8 mol% yttria stabilised zirconia used for button cells; (b) the electrolyte thickness is higher in the case of the tubular cell (0.425 mm versus 0.27 mm for button cells) and (c) the LSCF ink has been brush coated in the case of the tubular cell compared to screen printing used for button cells. The total power produced by the tubular cell was w175 mW. The power output can be significantly
Fig. 13 e Voltageecurrent density and voltage-power density characteristics of the DCFC tubular cell tested at 820 C with Vulcan XC-72 carbon. LSCF was used as the anode as well as the cathode.
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improved with the use of a thinner 8 mol% yttria stabilised zirconia electrolyte, incorporation of a more conducting phase to enhance electronic conductivity in the anode and the use of a more optimised electrode coating techniques such as spray coating or dip coating onto a cathode supported tubular structure [28].
4.
Conclusion
LSCF, a commonly used perovskite air electrode, has been evaluated as the fuel electrode in a direct carbon fuel cell. The material used was deficient at the A-site sublattice. A significant observation was the detection, by high resolution Synchrotron XRD, of the CoO phase (B-site sublattice material), both in the as-received powder and sintered (in air) LSCF bars, which otherwise was not detected using a laboratory X-ray source using Cu Ka radiation. The total conductivity of LSCF in CO is around 1 S-cm1, however, the ionic conductivity appears to be twice than that of YSZ used in traditional Ni/YSZ anodes. The rhombohedral LSCF phase remained the most dominant phase even after 50 h of heat treatment in CO at the typical DCFC operating temperature of 800 C. Noteworthy was the finding that the total conductivity of LSCF did not change much after about 20 min of heat treatment in CO at 800 C up to further 4880 min of testing. The conductivity data clearly demonstrated the structural stability of LSCF in DCFC operating environments consistent with phase analysis by high resolution XRD. Although the electrical conductivity of LSCF in reducing environments is significantly lower than that of traditionally used Ni based SOFC anodes, power densities exceeding 50 mW/cm2 were achieved in a electrolyte supported button cell DCFC with LSCF as both the anode and the cathode. This power density is substantially higher than that reported for Ni/YSZ anode in a DCFC. This is most probably due to significantly higher ionic conductivity of LSCF compared with YSZ leading to lower voltage losses when carbon electrochemical oxidation reaction takes place at the two phase carbon/anode interface. The polarisation losses, when deconvoluted using GCI and EIS techniques, showed that the anodic polarisation was still quite high and along with resistive losses across the electrolyte were mainly responsible for the lower power densities in direct carbon contact DCFCs in this study. This was attributed to the low electronic conductivity of LSCF in reducing environments. Furthermore, a scaled up version of DCFC with LSCF anode, operating on solid carbon fuel, was demonstrated with tubular design electrolyte supported cell which delivered a total power up to 175 mW and a peak power density of 25 mW/cm2. Overall, mixed ionic electronic conducting LSCF material showed promising performance as the DCFC anode with solid carbon fuel, despite its low electronic conductivity which can be improved further with the addition of a metal or conducting inert ceramic phase to enhance electronic conductivity and optimisation of the anode microstructure. Thus the concept of using mixed ionic electronic conducting anode for a solid fuel to extend the reaction zone for carbon oxidation from anode/ electrolyte interface to anode/solid fuel interface for continuous operation of the direct carbon fuel cell has been clearly demonstrated. This type of fuel cell and operation mode
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would offer the highest efficiency for direct conversion of carbon from coal or biomass to electricity provided the anode material has both high ionic and electronic conductivities.
Acknowledgements Authors would like to thank Dr. David Hay for reviewing this manuscript. The work was carried out in CSIRO’s Advanced Coal Technology Portfolio.
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