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Cathodes for fuel cells using proton-conducting Li2SO4–Al2O3 electrolyte
Journal of
Electroanalytical Chemistry Journal of Electroanalytical Chemistry 575 (2005) 183–193 www.elsevier.com/locate/jelechem
Cathodes for fuel cells using proton-conducting Li2SO4–Al2O3 electrolyte Guo-Lin Wei, Juri Melnik, Jing-Li Luo *, Alan R. Sanger, Karl T. Chuang Department of Chemical and Materials Engineering, University of Alberta, Room 536 CME, Edmonton, Alta., Canada T6G 2G6 Received 24 June 2004; received in revised form 18 September 2004; accepted 21 September 2004 Available online 11 November 2004
Abstract The performances of three widely different cathode materials (Pt, strontium-doped lanthanum manganite (LSM), and NiO) have been compared for use with proton conducting Li2SO4–Al2O3 composite electrolyte, using H2S–air and H2–air fuel cells operating at 600 °C. Surface analysis and electrochemical techniques were used to characterize fresh and used electrode materials. Pt or LSM cathodes each became covered with Li2SO4 and Al2O3 and, as a consequence, the fuel cells showed poor performance. In contrast, the NiO cathode catalyst did not become covered with Li2SO4 and good fuel cell performance was achieved. Exceptionally good current densities of over 100 mA/cm2 and power densities of over 30 mW/cm2 were obtained for H2S–air fuel cells having Mo– Ni–S anode catalysts. Slight agglomeration of NiO particles during fuel cell operation had only a minor effect on performance. Ó 2004 Elsevier B.V. All rights reserved. Keywords: Cathode polarization; Fuel cell; Proton conductor; NiO; LSM; Pt
1. Introduction Fuel cells provide a clean and efficient option for local electricity generation [1]. The performance of a fuel cell depends not only on the chemical and physical nature of the electrode materials and electrolyte, but also on their mutual chemical and physical compatibilities at the interfaces, under the conditions during assembly and for operation of the fuel cell. Significant efforts have been directed to optimization of electrode catalyst utilization, based on the activity of the catalyst, support materials, dispersion and loading techniques. Further efforts have been directed to development of the electrolyte materials, and toward more effective and economical techniques to prepare electrolyte layers from currently available materials. The selection of electrolyte materials is dependent on many factors, especially their ionic conductivity and *
Corresponding author. Tel.: +1 780 4922232; fax: +1 780 4922881. E-mail address: [email protected] (J.-L. Luo).
0022-0728/$ - see front matter Ó 2004 Elsevier B.V. All rights reserved. doi:10.1016/j.jelechem.2004.09.010
physical/chemical stability at operating temperatures. All electrolytes to date operate within a limited set of parameters, and many have limited lifetimes, with consequent economic drawbacks. Conventional solid oxide fuel cells (SOFCs) use oxygen ion conducting yttriadoped zirconia (YSZ), and usually operate at temperatures higher than 800 °C, and problems can arise from thermal instability and incompatibility of the materials. Consequently, new materials have been sought that offer superior electrochemical performance. To this end, recent efforts have also been directed to the development of intermediate temperature proton conducting fuel cells (<700 °C) using proton conducting perovskite-type oxide ceramics [2–4], as well as inorganic salts and their composites [1,5–7]. Ceramic oxides have been widely studied as proton conducting electrolytes, and the best candidates to date, such as BaCeO3, show conductivity values of the order of 102 S/cm at 600 °C [2]. Recently, Y-doped BaZrO3-based oxides have shown promise for use as electrolytes for intermediate-temperature SOFCs [4]. These proton conductors have conductivities higher
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than that of the best oxide ion conductors at temperatures below 700 °C. Inorganic salts like alkali metal sulfates/bisulfates, phosphates/biphosphates and nitrates, and their composites, have also been investigated as electrolytes in fuel cells [5–7]. Proton conduction has been demonstrated for neutral Li2SO4 [8–10]. Voltage responses were measured with hydrogen concentration cells, and the proton transport number was shown to be higher than 0.9. No similar voltage response was found using oxygen concentration cells, showing that this material was not an oxygen ion conductor. Protons were taken up by the electrolyte during the wet-chemical preparation of the salt. Infrared spectroscopy showed that the protons became bonded to the oxygen atoms of the sulfate group and formed co-ordination bonds, H–SO4 [11]. A ‘‘paddle-wheel’’ mechanism was proposed for proton conduction in which there was a dynamic process to transfer protons between sulphate groups [12]. Progress has been made in the development of such proton conducting inorganic salt electrolytes for fuel cells operated at intermediate (300–700 °C) [5,6] and more moderate temperatures (100–200 °C) [7]. Protons are considered to be the only widely mobile species in these electrolytes during fuel cell operation. All other cations or anions are less mobile, and are considered to be blocked or restricted from travelling significant distances through the conducting medium [5,8,11,12]. To date no such electrolyte material has been developed that combines the conductivity and long-term stability required for economic commercial operation. Issues that still need to be addressed include development of techniques to fabricate micrometer-thin, instead of millimeter-thick, electrolyte layers, and long-term stability of the electrolyte under cell operating conditions [1]. For example, it was found that CsHSO4 electrolyte decomposed to Cs2SO4 and H2S under an atmosphere of H2 in the presence of Pt/C catalyst, even though that electrolyte was stable under identical conditions in the absence of the catalyst [13]. The performance of H2–air fuel cells using proton conducting inorganic salts is generally far below that for state-of-the-art PEM systems due, in part, to various slow but as yet unidentified electrochemical processes [1,7]. Challenges remaining for the enhancement of electrode performance include the development of active and stable electrode materials, and optimization of their microstructures. H2S–air fuel cells are attracting attention as options for recovery of the chemical energy from H2S oxidation processes [14–23]. H2S is one of the most noxious, poisonous and abundant air pollutants and is produced in millions of tons annually. In conventional thermal catalytic systems, H2S is converted to elemental sulfur using the Claus process. Fuel cell technology provides an economically and environmentally desirable alternative to generate high-grade electric power from the large amount of chemical energy associated with the oxida-
tion of H2S, which would otherwise be either vented or partially recovered as low-grade steam. To date, most H2S–air fuel cells use oxide ion-conducting yttria-stabilized zirconia (YSZ) as the electrolyte, and are operated around 750–900 °C [14–19]. In these systems, either LSM [14,15] or Pt [16–23] is used as the cathode catalyst. Metal sulfides are the preferred anode catalysts. High-performance composite anode catalysts developed recently in our laboratory, and in particular a Mo–Ni–S composite prepared by heat-treating a 1:1 mixture by weight of MoS2 and NiS, showed excellent performance and durability [23]. When the fuel is H2, the only product of fuel cells is water, either at the cathode (proton-conducting) or at the anode (oxide-ion-conducting). In contrast, a mixture of products can potentially be formed in the anode chamber of H2S–air fuel cells using oxide-conducting electrolytes, including elemental sulfur, H2O and SO2 [15]. Consequently, the operating conditions must be controlled to prevent formation of SO2, or further processing of the effluent stream is required to capture SO2. However, when a proton-conducting electrolyte is used, high-purity sulfur is the only product obtained in the anode chamber, and water is the only product formed in the cathode chamber, so no subsequent SO2 separation process is required. Product sulfur exits the anode chamber as vapour since the fuel cell operating temperature is usually higher than the boiling point of sulfur (444.6 °C). Several proton-conducting electrolyte materials have been investigated for use in H2S–air fuel cells [20,21,24]. Li2SO4 has been used for H2–air fuel cells, but was found to be unstable during prolonged operation in pure H2 [25]. Thermodynamic analysis and XRD experiments [25] confirmed that the reaction products were Li2SO4 Æ H2O, Li2S and LiOH. Li2SO4 reacts with H2 according to the following reactions: Li2 SO4 þ 4H2 ¼ 4H2 O þ Li2 S
ð1Þ
Li2 SO4 þ 4H2 ¼ 2H2 O þ 2LiOH þ H2 S
ð2Þ
XRD analysis [25] also confirmed the formation of H2S in the anode chamber by showing that CuS was formed by reaction of CuSO4 solution, which was used to detect H2S, with anode effluent from fuel cells operated at 750 °C: CuSO4 þ H2 S ¼ CuS # þH2 SO4
ð3Þ
Li2SO4 was also among the first proton conducting materials tested for use as an electrolyte in H2S–air fuel cells, and a maximum short-circuit current density of 12 mA/cm2 was achieved at 725 °C [21]. Although Li2SO4 is unstable in a reducing H2 atmosphere, this electrolyte is stable when used in a H2S atmosphere [21,24]. However, membranes made of Li2SO4 alone have low mechanical strength. Membranes constructed from Li2SO4–Al2O3 composites have enhanced mechanical strength and elec-
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trical conductivity compared with those made from pure Li2SO4. Recently, we developed an improved method for preparation of strong and gas-impermeable Li2SO4–Al2O3 composite membranes [10]. We have now found that the performance of H2S– air fuel cells having Li2SO4–Al2O3 composite membranes depends upon the nature of the cathode catalyst. We will show that this effect is due to electrolyte transport onto some cathode surfaces, and that this is a phenomenon that also affects the performance of conventional H2–air fuel cells using inorganic salts as electrolytes. We have compared the performance using three widely different cathode materials (Pt; strontiumdoped lanthanum manganite (LSM); and NiO). We will describe how transport of Li2SO4 from the electrolyte onto the cathode surface compromises the activity of Pt or LSM, but not NiO, in fuel cells operated at 600 °C.
2. Experimental 2.1. Electrolyte membrane To make electrolyte membranes, 0.8–1.1 g wet mixtures having 60:40:10 weight ratios of Li2SO4 (Alfa Aesar), Al2O3 (99.95% submicron powder, Alfa Aesar) and deionized H2O were pressed to form green disks. The disks were dried in air at 105 °C for 2 h and sintered at 850 °C for 8 h [10]. The resulting disks were 0.7–1.0 mm thick and 24–25 mm in diameter. 2.2. Electrodes The three different cathode materials (Pt, LSM and NiO) were applied by screen-printing the corresponding paste onto one side of the electrolyte membrane. Pt was applied as Heraeus CL11-5100 Pt paste. Both LSM (NextTech Materials, admixed with 5 wt% Ag (Alfa Aesar) to improve electrical conductivity) and NiO (Alfa Aesar) catalysts were applied as dispersions in a-terpineol. The electrolyte–cathode assembly was cured in air in a furnace before applying the anode. The temperature was first raised at 2 °C/min to 230 °C, kept at that temperature for 1 h, then raised at 2 °C/min to 750 °C and kept at that temperature for a further 1 h. The anode catalysts were a Mo–Ni–S-based composite [19] for H2S–air fuel cells and Pt for H2–air fuel cells. Catalysts were applied to the membrane as a paste in aterpineol or a Pt paste (Heraeus CL11-5100). The metal sulfide-based composite anode contained 95 wt% Mo– Ni–S, prepared by heat treatment of a 1:1 weight ratio mixture of MoS2 and NiS (both from Alfa Aesar), which was then admixed with 5 wt% Ag to enhance electrical conductivity [23]. A paste was prepared by dispersing these well-mixed anode materials in a-terpineol. The
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paste was converted into the catalytic anode by sequentially applying the paste to the side of the electrolyte membrane opposite that with the cured cathode, then drying the paste on the membrane in air, and finally curing the whole assembly in a nitrogen atmosphere in a furnace. The temperature was first raised at 2 °C/min to 230 °C and kept at that temperature for 1 h. It was then raised at 2 °C/min to 750 °C and kept at that temperature for a further 1 h. The cured assembly was then slowly cooled to room temperature. Pt mesh was used as the current collector at each electrode. The reference electrode was a spiral Pt wire in contact with a Pt spot 2 mm in diameter formed by applying and curing Pt paste on the cathode side of the electrolyte–electrode assembly. 2.3. Fuel cell set-up and electrochemical measurements The fuel cell test station was similar to that described in [18], with minor modifications that did not alter the principles of operation. To install the electrolyte–electrode assembly into a test fuel cell, first the cathode side of the assembly was attached to a supporting annular alumina disk, 3.2 cm in diameter and 0.3 cm in thickness, using ceramic sealant (Aremco 503). An opening 1.1 cm in diameter in the center of the supporting ceramic disk allowed air to access the cathode. The combination so made was then sealed with ceramic sealant between two alumina tubular chambers (outer dimension 2.54 cm, length 40 cm). The surface of the Pt mesh used as the anode current collector was refreshed in the flame of a gas lamp prior to installation for each experiment. The cell was then heated in a tubular furnace (Thermolyne F79300), with nitrogen passing through the anode chamber and air through the cathode chamber, each at 25 ml/min. To cure the sealant (Aremco 503), the furnace temperature was increased to 230 °C at 1 °C/min, and held at that temperature for 1 h. The temperature then was increased to a selected testing temperature, typically 600 °C, and held at that temperature for the duration of a series of tests. Cell performances were first determined using H2S as the anode feed. The cell open circuit voltage (ocv) and electrode potentials were monitored as functions of time on stream. Electrode potentials were recorded versus a Pt reference electrode using a National Instruments V-Bench system and other data were recorded with a Gamry PC4-750 measurement system. Since the reference electrode was located at the cathode side of the electrolyte membrane, the recorded change in anode potential included both the anodic electrochemical overpotential and the ohmic drop arising from the electrolyte resistance. After a steady ocv was achieved, potentiodynamic measurements were conducted at a scan rate of 5 mV/s to determine the current–voltage performance of the H2S–air fuel cell in the IR compensation mode using
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Fig. 1. Change in cell voltage and electrode potentials with anode feed (anode: Mo–Ni–S + Ag; cathode: Pt T: 600 °C, H2S–air feed rate: 25/25 ml/ min. Switching N2–H2S at time 0. Starting voltage scanning at time 2475 s).
1.4
35 I-V, Pt I-V, LSM I-V, NiO I-P, Pt I-P, LSM I-P, NiO
Voltage (V)
1 0.8
30 25 20
0.6
15
0.4
10
0.2
5
0 0
20
40
60
80
100
Power (mW/cm2)
1.2
0 120
Current (mA/cm-2)
Fig. 2. Performances for H2S–air fuel cell using different cathodes (T: 600 °C. H2S–air feed rate: 25/25 ml/min. Anode: Mo–Ni–S + Ag).
Anode potential (V)
-1.5
-1.3
-1.1
-0.9
-0.7 0
20
40
60
80
100
120
140
-2
Current (mA/cm )
Fig. 3. Current-potential curves for Mo–Ni–S composite anode in H2S stream (T: 600 °C; H2S/air feed rate: 25/25 ml/min; Cathode: NiO).
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the Gamry system. The performance of H2–air fuel cells was also then determined to investigate whether the nature of the anode feed affected the performance of these cathodes. 2.4. Surface analysis The electrode morphologies and compositions were studied with scanning electron microscopy (SEM) and energy dispersive X-ray (EDX) techniques using a Hitachi S-2700 scanning electron microscope and PGT Imix system with a PRISM IG or a JEOL 6301F field emission scanning microscope and an attached PGT Analyzer. In some cases, a layer of gold was sputtered onto the SEM samples to increase the electrical conduction. X-ray diffraction (XRD) analyses were conducted using a Rigaku Rotaflex X-ray diffractometer. XRD spectra were acquired and compared with samples of mixtures of various cathodes and electrolyte (Li2SO4 + Al2O3) with and without heat-treatment in air at 750 °C for 6 h.
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was scanned from 1.411 to 0 V, the anode potential apparently changed from 1.447 to 1.332 V, while the cathode potential changed from 0.036 to 1.332 V (Fig. 1). Thus, the cathode became more highly polarized (1.296 V) than the anode (less than 0.115 V), indicating that the cathodic kinetics were much slower than those at the anode. The recorded anode potential change included both the anodic polarization and the ohmic drop from the electrolyte resistance. This result differed from that for H2S–air fuel cells using oxide-ion-conducting YSZ as the electrolyte and
3. Results and discussion 3.1. Gas impermeability of the electrolyte membrane The membrane had to be gas impermeable so as to ensure that the data obtained during this research were free from effects arising from cross-over of gases. To serve this purpose, the cell voltage and both the cathode and anode potentials were monitored when the anode feed was changed from N2 to H2S and, during cell performance, were measured. Fig. 1 shows the typical experimental results obtained with the Li2SO4–Al2O3 electrolyte membrane used in our tests. It is obvious that the cathode potential was unaffected by the anode feed, and remained constant with time on stream, until voltage scanning was performed. In contrast, when a less dense Li2SO4–Al2O3 membrane was used, there was cross-over of H2S into the cathode chamber, which resulted in fluctuations of the order of hundreds of mV in the cell voltage and deterioration of the reference electrode. Thus, the data in Fig. 1 show that there was no cross-over in the present system using a dense Li2SO4– Al2O3 membrane. 3.2. Polarization of the Pt cathode and Mo–Ni–S based anode in the H2S–air fuel cell To discern which electrode limited cell performance, polarizations of both cathode and anode were monitored during current–voltage behaviour measurements of H2S–air fuel cells. When a Pt cathode and Mo– Ni–S based anode were used, and the cell voltage
Fig. 4. SEM of Mo–Ni–S anode (a) before and (b) after H2S–air fuel cell tests.
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Pt as the cathode, for which the cathode polarization was 200 mV and the anode polarization 400 mV at a current density of 130 mA/cm2 [26]. Thus, there was a dramatic effect on the performance of the Pt cathode attributable to the switch from oxide-ion-conducting YSZ to proton-conducting Li2SO4–Al2O3 as the electrolyte. Consequently, two additional, widely different cathode materials were also examined to determine whether the effect was due to the cathode, the electrolyte or the combination of these materials. 3.3. H2S–air fuel cell performance using different cathodes Fig. 2 shows in more detail the dependence of H2S– air fuel cell performance on the different cathode material used. The results were typical and reproducible for each of the specific cathodes. In each case the electrolyte was Li2SO4–Al2O3 and the anode catalyst was Mo–Ni–S intimately admixed with 5 wt% Ag. The maximum current densities and power densities achieved using the three widely different cathode materials were: Pt, 29 mA/cm2 and 10 mW/cm2; LSM, 60 mA/cm2 and 19 mW/cm2; and NiO, 96 mA/cm2 and 33 mW/cm2. Thus, the performance of the fuel cell depended strongly on the cathode used.
3.4. Performance of Mo–Ni–S composite anode in an H2S–air fuel cell Fig. 3 shows the current–electrode potential curve with a Mo–Ni–S composite anode as the working electrode on Li2SO4–Al2O3 as electrolyte with NiO as the counter electrode when using H2S as the fuel and air as the oxidant. The potential at the anode was controlled relative to that at the Pt reference electrode, which contacted the electrolyte from the cathode side. It showed a potential change (polarization) of 460 mV at 100 mA/cm2. Comparison of Fig. 3 with Fig. 2 shows that, at the same current density, the change in anode potential was much less than the change in potential for each cathode material. At the maximum current density, the anodic polarization was 57 mV with a Pt cathode (at 29 mA/cm2), 168 mV with LSM (at 60 mA/cm2) and 440 mV with NiO (at 98 mA/cm2), respectively. Meanwhile, the corresponding cathode polarizations were 1.353 V for Pt, 1.242 V for LSM, and 0.97 V for NiO. Therefore, in each case, the majority of the change in cell voltage occurred as a result of cathodic polarization. The activity of the anode under the present conditions was sufficiently high that the anodic reaction kinetics were not the limiting step.
Fig. 5. EDX of Mo–Ni–S anode (a) before and (b) after H2S–air fuel cell tests.
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3.5. Morphologies and composition of anode surfaces The surfaces of Mo–Ni–S anodes were examined before and after testing in a H2S–air fuel cell for several hours, using SEM (Fig. 4) and EDX (Fig. 5), to determine whether there were any morphological or compositional changes resulting from operation of the material in a fuel cell. The anode retained its highly porous microstructure. There were no significant changes in
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morphology (Fig. 4) or chemical composition (Fig. 5) arising from the fuel cell tests. 3.6. Changes at the cathodes during testing of H2S–air fuel cells The morphologies and compositions of different cathode materials before and after fuel cell testing were determined to provide insight into the causes of the high cathodic polarizations observed with Pt, LSM and NiO cathodes. Results are shown in Figs. 6–10 and Table 1. 3.6.1. Pt electrodes 3.6.1.1. Pt as cathode. The Pt cathodes were applied to the membrane as pastes and then heat-treated in air. Fresh Pt cathodes were very porous (Fig. 6(a)) and highly conductive, showing a dc ohmic resistance of less than 1 X, which arose mainly from contact resistances of the measuring probes. The performances of cells having Pt cathodes were very poor and deteriorated rapidly during use in H2S–air fuel cells. The used cathode was found to have developed a less porous structure (Fig. 6(b)) and showed very poor conduction, with the dc ohmic resistance being higher than 100 MX. EDX showed that the used Pt cathode comprised only 37a/o Pt, and had 20a/o S and 43a/o O, in contrast to the fresh Pt cathode (100a/o Pt) (Table 1). Pt was exposed only at
Fig. 6. Morphology of Pt cathodes (a) before and (b)–(c) after fuel cell tests: (b) overall; (c) detail. Compositions at Points 1 and 2 in (c) are in Table 1.
Fig. 7. LSM + Ag cathodes (a) before and (b) after fuel cell tests.
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Fig. 8. EDX results of LSM + Ag cathode (a) before and (b) after fuel cell tests.
isolated sites represented by site 1 in Fig. 6(c), and the cathode was under an extensive coverage by electrolyte materials represented by site 2. When these sites were compared, EDX showed that site 1 maintained a high Pt content of 81a/o, and site 2 had only 7a/o Pt but 35a/o S and 58a/o O. Thus, the performance of the used cathodes was affected detrimentally as the electrolyte blocked access of oxygen molecules to Pt catalytic sites and inhibited transport of electrons between different Pt particles, resulting in high polarization of the Pt cathode. In contrast, contacts between Pt cathodes on YSZ membranes were found to remain intact during prolonged operation of the fuel cells. Consequently, the effect is directly related to the nature of the Li2SO4–Al2 O3 electrolyte, as well as the Pt catalyst, as shown below. 3.6.1.2. Pt as anode. The performance and stability of Pt as the anode catalyst was also investigated for use with Li2SO4–Al2O3 electrolyte, to determine whether the above surface-covering effect arising during fuel cell testing was related only to the nature of the materials. An H2–air fuel cell using Li2SO4–Al2O3 electrolyte and operated at 600 °C was used instead of an H2S–air fuel cell, since it was known that reaction of Pt with H2S could cause delamination of the anode [16,17]. The EDX spectrum for a Pt anode after fuel cell tests showed only a single Pt peak. No signals were observed for Al,
O or S, which showed that the used Pt anode had not become covered by electrolyte, in contrast to the extensive coverage of the Pt cathode by electrolyte. Thus, transport of electrolyte onto the Pt electrode was not a function of the nature of either Pt alone or the combination of Pt and Li2SO4–Al2O3 electrolyte. Transport of electrolyte to cause covering of the Pt surface occurred only when Pt was used as the cathode during fuel cell testing. In contrast, no compositional change was observed in XRD spectra of Pt and electrolyte mixtures subjected to heat-treatment in air at 750 °C for 6 h under non-fuel cell operating conditions. Also, no covering by electrolyte was detected on Pt reference electrodes that were prepared from the same paste and located in the same cathode chamber during the fuel cell tests when the Pt cathodes became coated. Thus transport and covering of electrolyte were related directly to the electrochemical process(es) during fuel cell operation. It is unclear whether the local transport of electrolyte and its mixing with cathode materials arose through mobility of ions other than protons, as there was unlikely to be significant mobility over significant distances of ions other than protons [5,8,11,12]. 3.6.2. LSM cathodes LSM has been used extensively as the cathode catalyst in SOFCs using a variety of fuels. However, we have
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now found that similar electrolyte transport effects were observed with LSM cathodes when used with Li2SO4– Al2O3 membranes. LSM admixed with 5 wt% Ag was applied as a paste to one face of Li2SO4–Al2O3 membranes. The Ag enhanced the electrical conductivity of the LSM cathode, and was unaffected by O2 at the operating temperature of the fuel cell. The fresh catalysts were heat-treated in air, after which they were very porous (Fig. 7(a)). The elemental composition of the fresh cathode surface was entirely LSM and Ag (Au signals arose from Au sputtered as the electrical conductor for SEM and as the reference for EDX determinations) (Fig. 8(a)). In contrast, the surface of used LSM cathodes comprised mostly particles significantly larger than those for fresh catalysts (Fig. 7(b)). The large surface particles on the used LSM cathode formed a much denser surface that was less accessible by O2. EDX spectra showed only traces of La, Mn, Sr and Ag in this layer (Fig. 8(b)). The strong presence of sulfur peaks indicated that the cathode surface had been covered almost entirely by the migrated sulfate. The cathode agglomeration effect appeared to be directly related to the transport of electrolyte, possibly due to adherence of particles caused by the coating of electrolyte. Thus the
Fig. 9. NiO cathode (a) before and (b) after fuel cell tests.
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performance of LSM cathodes was affected detrimentally both by coverage of catalyst sites by the electrolyte and by reduced access by O2 to the available catalyst sites. Mixtures of LSM and electrolyte behaved in a manner similar to mixtures of Pt and electrolyte. Mixtures that had undergone heat-treatment in air at 750 °C for 6 h showed no compositional change in XRD spectra. In contrast, there was even more severe covering of LSM by electrolyte after fuel cell tests than was found for Pt (Figs. 7, 8). 3.6.3. NiO cathodes In contrast to either Pt or LSM cathodes, comparison of fresh and used NiO cathodes showed only a small difference in surface morphology (Fig. 9). EDX analysis showed that the composition of the surface of used NiO catalyst was mainly similar to that of fresh catalyst (Fig. 10), which indicated very little transport of electrolyte materials onto the surface of NiO cathodes. Consequently, the performance of fuel cells having NiO cathodes was better and more stable than that for fuel cells with Pt or LSM cathodes (Fig. 2). Although it was found that cathode NiO particles agglomerated slightly during use, the cathode layer remained porous, and so the catalyst sites remained accessible by O2 (Fig. 9). Because there was no surface coating of electrolyte on the NiO particles, agglomeration of catalyst particles was not a consequence of electrolyte material transport, in contrast to the effect found for Pt and LSM (Fig. 7). Lithium doping has been shown to improve the electrical conductivity, chemical diffusivity, and oxygen exchange kinetics of NiO [27]. During operation of the present fuel cells, Li ions could have doped into NiO particles located adjacent to the electrolyte membrane, instead of covering its surface, and thereby could have enhanced both the ionic conductivity of the catalyst and oxygen reduction kinetics. However, no change was detected in XRD spectra of the NiO cathode after the fuel cell test, but this may have been a consequence of the inability of EDX to detect Li. A combination of NiO powder with electrolyte that was heat-treated in air also showed no change in the XRD spectrum. To summarize: the performance of proton-conducting fuel cells having Li2SO4–Al2O3 membranes depended strongly on the compatibility of the electrode materials with the Li2SO4–Al2O3 membrane material. Pt and LSM cathodes were unsuitable for use with Li2SO4–Al2O3 membranes during operation at temperatures at 600 °C, as the electrolyte material migrated onto the cathode catalyst surface. However, neither NiO cathode catalysts nor Pt or Mo–Ni–S anode catalysts became covered with this electrolyte. Thus, we have shown that the use of Li2SO4–Al2O3 as the electrolyte in proton-conducting fuel cells depends on the selection of cathode materials.
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Fig. 10. EDX of NiO cathode (a) before and (b) after fuel cell tests.
Mo–Ni–S anode catalysts and the electrolyte transport effect appears to be limited to the cathode.
Table 1 Surface atomic composition in at% of a used Pt cathode (EDX) Analysis area
O
S
Pt
Overall (Fig. 6(b)) Site 1 (Fig. 6(c)) Site 2 (Fig. 6(c))
43 9 58
20 10 35
37 81 7
4. Conclusions Transport of electrolyte materials onto the surface of Pt or LSM cathodes adversely affected the performance of proton-conducting fuel cells using Li2SO4–Al2O3 as the electrolyte and operating at temperatures at 600 °C. In contrast, no significant transport of electrolyte onto the cathode was observed when NiO was the cathode catalyst. Covering the cathode surface with electrolyte reduced the accessibility to catalyst sites by O2 and reduced the electrical conductivity of the cathode layer; consequently the covered cathodes became highly polarized. In contrast, NiO remained mainly uncovered after use. Consequently, H2S–air fuel cells having NiO cathodes showed the highest cell performance, with a current density of over 100 mA/cm2 and a power density of over 30 mW/cm2. Slight agglomeration of NiO particles did not appear to affect the performance of the cathode. The same electrolyte did not migrate onto either Pt or
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T. Norby, Nature 410 (2001) 877. T. Norby, Solid State Ionics 125 (1999) 1. G. Alberti, M. Casciola, Solid State Ionics 145 (2001) 3. K.D. Kreuer, Annu. Rev. Mater. Res. 33 (2003) 333. B. Zhu, Solid State Ionics 125 (1999) 397, and references therein. B. Zhu, J. Power Sources 114 (2003) 1, and references therein. S.M. Haile, D.A. Boysen, C.R.I. Chisholm, R.B. Merle, Nature 410 (2001) 910. B. Heed, B. Zhu, B.-E. Mellander, Solid State Ionics 46 (1991) 121. K.D. Kreuer, Chem. Mater. 8 (1996) 610. J. Melnik, W.Z. An, G.-L. Wei, J.-L. Luo, A.R. Sanger, K.T. Chuang, Solid State Ionics, submitted. B. Zhu, B.-E. Mellander, Ferroelectrics 167 (1995) 1. B. Zhu, B.-E. Mellander, Solid State Ionics 77 (1995) 244. B. Yang, A.M. Kannan, A. Manthiram, Mater. Res. Bull. 38 (2003) 691. N.U. Pujare, K.W. Semkow, A.F. Sammells, J. Electrochem. Soc. 134 (1987) 2639. C. Yates, J. Winnick, J. Electrochem. Soc. 146 (1999) 2841. M. Liu, P. He, J.-L. Luo, A.R. Sanger, K.T. Chuang, J. Power Sources 94 (2001) 20. P. He, M. Liu, J.-L. Luo, A.R. Sanger, K.T. Chuang, J. Electrochem. Soc. 149 (2002) A808.
G.-L. Wei et al. / Journal of Electroanalytical Chemistry 575 (2005) 183–193 [18] G.-L. Wei, M. Liu, J.-L. Luo, A.R. Sanger, K.T. Chuang, J. Electrochem. Soc. 150 (2003) A463. [19] K.T. Chuang, J.-L. Luo, G.-L. Wei, A.R. Sanger, US Patent Application 20030215697, 2003. [20] T.J. Kirk, J. Winnick, J. Electrochem. Soc. 140 (1993) 3494. [21] D. Peterson, J. Winnick, J. Electrochem. Soc. 143 (1996) L55. [22] D. Peterson, J. Winnick, J. Electrochem. Soc. 145 (1998) 1449.
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[23] G.-L. Wei, J.-L. Luo, A.R. Sanger, K.T. Chuang, J. Electrochem. Soc. 151 (2004) A232. [24] B. Zhu, S. Tao, Solid State Ionics 127 (2000) 83. [25] S.W. Tao, Z.L. Zhan, P. Wang, G.Y. Meng, Solid State Ionics 116 (1999) 29. [26] M. Liu, Ph.D. Thesis, University of Alberta, 2004. [27] E. Antolini, Mater. Chem. Phys. 82 (2003) 937, and references therein.
Electroanalytical Chemistry Journal of Electroanalytical Chemistry 575 (2005) 183–193 www.elsevier.com/locate/jelechem
Cathodes for fuel cells using proton-conducting Li2SO4–Al2O3 electrolyte Guo-Lin Wei, Juri Melnik, Jing-Li Luo *, Alan R. Sanger, Karl T. Chuang Department of Chemical and Materials Engineering, University of Alberta, Room 536 CME, Edmonton, Alta., Canada T6G 2G6 Received 24 June 2004; received in revised form 18 September 2004; accepted 21 September 2004 Available online 11 November 2004
Abstract The performances of three widely different cathode materials (Pt, strontium-doped lanthanum manganite (LSM), and NiO) have been compared for use with proton conducting Li2SO4–Al2O3 composite electrolyte, using H2S–air and H2–air fuel cells operating at 600 °C. Surface analysis and electrochemical techniques were used to characterize fresh and used electrode materials. Pt or LSM cathodes each became covered with Li2SO4 and Al2O3 and, as a consequence, the fuel cells showed poor performance. In contrast, the NiO cathode catalyst did not become covered with Li2SO4 and good fuel cell performance was achieved. Exceptionally good current densities of over 100 mA/cm2 and power densities of over 30 mW/cm2 were obtained for H2S–air fuel cells having Mo– Ni–S anode catalysts. Slight agglomeration of NiO particles during fuel cell operation had only a minor effect on performance. Ó 2004 Elsevier B.V. All rights reserved. Keywords: Cathode polarization; Fuel cell; Proton conductor; NiO; LSM; Pt
1. Introduction Fuel cells provide a clean and efficient option for local electricity generation [1]. The performance of a fuel cell depends not only on the chemical and physical nature of the electrode materials and electrolyte, but also on their mutual chemical and physical compatibilities at the interfaces, under the conditions during assembly and for operation of the fuel cell. Significant efforts have been directed to optimization of electrode catalyst utilization, based on the activity of the catalyst, support materials, dispersion and loading techniques. Further efforts have been directed to development of the electrolyte materials, and toward more effective and economical techniques to prepare electrolyte layers from currently available materials. The selection of electrolyte materials is dependent on many factors, especially their ionic conductivity and *
Corresponding author. Tel.: +1 780 4922232; fax: +1 780 4922881. E-mail address: [email protected] (J.-L. Luo).
0022-0728/$ - see front matter Ó 2004 Elsevier B.V. All rights reserved. doi:10.1016/j.jelechem.2004.09.010
physical/chemical stability at operating temperatures. All electrolytes to date operate within a limited set of parameters, and many have limited lifetimes, with consequent economic drawbacks. Conventional solid oxide fuel cells (SOFCs) use oxygen ion conducting yttriadoped zirconia (YSZ), and usually operate at temperatures higher than 800 °C, and problems can arise from thermal instability and incompatibility of the materials. Consequently, new materials have been sought that offer superior electrochemical performance. To this end, recent efforts have also been directed to the development of intermediate temperature proton conducting fuel cells (<700 °C) using proton conducting perovskite-type oxide ceramics [2–4], as well as inorganic salts and their composites [1,5–7]. Ceramic oxides have been widely studied as proton conducting electrolytes, and the best candidates to date, such as BaCeO3, show conductivity values of the order of 102 S/cm at 600 °C [2]. Recently, Y-doped BaZrO3-based oxides have shown promise for use as electrolytes for intermediate-temperature SOFCs [4]. These proton conductors have conductivities higher
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than that of the best oxide ion conductors at temperatures below 700 °C. Inorganic salts like alkali metal sulfates/bisulfates, phosphates/biphosphates and nitrates, and their composites, have also been investigated as electrolytes in fuel cells [5–7]. Proton conduction has been demonstrated for neutral Li2SO4 [8–10]. Voltage responses were measured with hydrogen concentration cells, and the proton transport number was shown to be higher than 0.9. No similar voltage response was found using oxygen concentration cells, showing that this material was not an oxygen ion conductor. Protons were taken up by the electrolyte during the wet-chemical preparation of the salt. Infrared spectroscopy showed that the protons became bonded to the oxygen atoms of the sulfate group and formed co-ordination bonds, H–SO4 [11]. A ‘‘paddle-wheel’’ mechanism was proposed for proton conduction in which there was a dynamic process to transfer protons between sulphate groups [12]. Progress has been made in the development of such proton conducting inorganic salt electrolytes for fuel cells operated at intermediate (300–700 °C) [5,6] and more moderate temperatures (100–200 °C) [7]. Protons are considered to be the only widely mobile species in these electrolytes during fuel cell operation. All other cations or anions are less mobile, and are considered to be blocked or restricted from travelling significant distances through the conducting medium [5,8,11,12]. To date no such electrolyte material has been developed that combines the conductivity and long-term stability required for economic commercial operation. Issues that still need to be addressed include development of techniques to fabricate micrometer-thin, instead of millimeter-thick, electrolyte layers, and long-term stability of the electrolyte under cell operating conditions [1]. For example, it was found that CsHSO4 electrolyte decomposed to Cs2SO4 and H2S under an atmosphere of H2 in the presence of Pt/C catalyst, even though that electrolyte was stable under identical conditions in the absence of the catalyst [13]. The performance of H2–air fuel cells using proton conducting inorganic salts is generally far below that for state-of-the-art PEM systems due, in part, to various slow but as yet unidentified electrochemical processes [1,7]. Challenges remaining for the enhancement of electrode performance include the development of active and stable electrode materials, and optimization of their microstructures. H2S–air fuel cells are attracting attention as options for recovery of the chemical energy from H2S oxidation processes [14–23]. H2S is one of the most noxious, poisonous and abundant air pollutants and is produced in millions of tons annually. In conventional thermal catalytic systems, H2S is converted to elemental sulfur using the Claus process. Fuel cell technology provides an economically and environmentally desirable alternative to generate high-grade electric power from the large amount of chemical energy associated with the oxida-
tion of H2S, which would otherwise be either vented or partially recovered as low-grade steam. To date, most H2S–air fuel cells use oxide ion-conducting yttria-stabilized zirconia (YSZ) as the electrolyte, and are operated around 750–900 °C [14–19]. In these systems, either LSM [14,15] or Pt [16–23] is used as the cathode catalyst. Metal sulfides are the preferred anode catalysts. High-performance composite anode catalysts developed recently in our laboratory, and in particular a Mo–Ni–S composite prepared by heat-treating a 1:1 mixture by weight of MoS2 and NiS, showed excellent performance and durability [23]. When the fuel is H2, the only product of fuel cells is water, either at the cathode (proton-conducting) or at the anode (oxide-ion-conducting). In contrast, a mixture of products can potentially be formed in the anode chamber of H2S–air fuel cells using oxide-conducting electrolytes, including elemental sulfur, H2O and SO2 [15]. Consequently, the operating conditions must be controlled to prevent formation of SO2, or further processing of the effluent stream is required to capture SO2. However, when a proton-conducting electrolyte is used, high-purity sulfur is the only product obtained in the anode chamber, and water is the only product formed in the cathode chamber, so no subsequent SO2 separation process is required. Product sulfur exits the anode chamber as vapour since the fuel cell operating temperature is usually higher than the boiling point of sulfur (444.6 °C). Several proton-conducting electrolyte materials have been investigated for use in H2S–air fuel cells [20,21,24]. Li2SO4 has been used for H2–air fuel cells, but was found to be unstable during prolonged operation in pure H2 [25]. Thermodynamic analysis and XRD experiments [25] confirmed that the reaction products were Li2SO4 Æ H2O, Li2S and LiOH. Li2SO4 reacts with H2 according to the following reactions: Li2 SO4 þ 4H2 ¼ 4H2 O þ Li2 S
ð1Þ
Li2 SO4 þ 4H2 ¼ 2H2 O þ 2LiOH þ H2 S
ð2Þ
XRD analysis [25] also confirmed the formation of H2S in the anode chamber by showing that CuS was formed by reaction of CuSO4 solution, which was used to detect H2S, with anode effluent from fuel cells operated at 750 °C: CuSO4 þ H2 S ¼ CuS # þH2 SO4
ð3Þ
Li2SO4 was also among the first proton conducting materials tested for use as an electrolyte in H2S–air fuel cells, and a maximum short-circuit current density of 12 mA/cm2 was achieved at 725 °C [21]. Although Li2SO4 is unstable in a reducing H2 atmosphere, this electrolyte is stable when used in a H2S atmosphere [21,24]. However, membranes made of Li2SO4 alone have low mechanical strength. Membranes constructed from Li2SO4–Al2O3 composites have enhanced mechanical strength and elec-
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trical conductivity compared with those made from pure Li2SO4. Recently, we developed an improved method for preparation of strong and gas-impermeable Li2SO4–Al2O3 composite membranes [10]. We have now found that the performance of H2S– air fuel cells having Li2SO4–Al2O3 composite membranes depends upon the nature of the cathode catalyst. We will show that this effect is due to electrolyte transport onto some cathode surfaces, and that this is a phenomenon that also affects the performance of conventional H2–air fuel cells using inorganic salts as electrolytes. We have compared the performance using three widely different cathode materials (Pt; strontiumdoped lanthanum manganite (LSM); and NiO). We will describe how transport of Li2SO4 from the electrolyte onto the cathode surface compromises the activity of Pt or LSM, but not NiO, in fuel cells operated at 600 °C.
2. Experimental 2.1. Electrolyte membrane To make electrolyte membranes, 0.8–1.1 g wet mixtures having 60:40:10 weight ratios of Li2SO4 (Alfa Aesar), Al2O3 (99.95% submicron powder, Alfa Aesar) and deionized H2O were pressed to form green disks. The disks were dried in air at 105 °C for 2 h and sintered at 850 °C for 8 h [10]. The resulting disks were 0.7–1.0 mm thick and 24–25 mm in diameter. 2.2. Electrodes The three different cathode materials (Pt, LSM and NiO) were applied by screen-printing the corresponding paste onto one side of the electrolyte membrane. Pt was applied as Heraeus CL11-5100 Pt paste. Both LSM (NextTech Materials, admixed with 5 wt% Ag (Alfa Aesar) to improve electrical conductivity) and NiO (Alfa Aesar) catalysts were applied as dispersions in a-terpineol. The electrolyte–cathode assembly was cured in air in a furnace before applying the anode. The temperature was first raised at 2 °C/min to 230 °C, kept at that temperature for 1 h, then raised at 2 °C/min to 750 °C and kept at that temperature for a further 1 h. The anode catalysts were a Mo–Ni–S-based composite [19] for H2S–air fuel cells and Pt for H2–air fuel cells. Catalysts were applied to the membrane as a paste in aterpineol or a Pt paste (Heraeus CL11-5100). The metal sulfide-based composite anode contained 95 wt% Mo– Ni–S, prepared by heat treatment of a 1:1 weight ratio mixture of MoS2 and NiS (both from Alfa Aesar), which was then admixed with 5 wt% Ag to enhance electrical conductivity [23]. A paste was prepared by dispersing these well-mixed anode materials in a-terpineol. The
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paste was converted into the catalytic anode by sequentially applying the paste to the side of the electrolyte membrane opposite that with the cured cathode, then drying the paste on the membrane in air, and finally curing the whole assembly in a nitrogen atmosphere in a furnace. The temperature was first raised at 2 °C/min to 230 °C and kept at that temperature for 1 h. It was then raised at 2 °C/min to 750 °C and kept at that temperature for a further 1 h. The cured assembly was then slowly cooled to room temperature. Pt mesh was used as the current collector at each electrode. The reference electrode was a spiral Pt wire in contact with a Pt spot 2 mm in diameter formed by applying and curing Pt paste on the cathode side of the electrolyte–electrode assembly. 2.3. Fuel cell set-up and electrochemical measurements The fuel cell test station was similar to that described in [18], with minor modifications that did not alter the principles of operation. To install the electrolyte–electrode assembly into a test fuel cell, first the cathode side of the assembly was attached to a supporting annular alumina disk, 3.2 cm in diameter and 0.3 cm in thickness, using ceramic sealant (Aremco 503). An opening 1.1 cm in diameter in the center of the supporting ceramic disk allowed air to access the cathode. The combination so made was then sealed with ceramic sealant between two alumina tubular chambers (outer dimension 2.54 cm, length 40 cm). The surface of the Pt mesh used as the anode current collector was refreshed in the flame of a gas lamp prior to installation for each experiment. The cell was then heated in a tubular furnace (Thermolyne F79300), with nitrogen passing through the anode chamber and air through the cathode chamber, each at 25 ml/min. To cure the sealant (Aremco 503), the furnace temperature was increased to 230 °C at 1 °C/min, and held at that temperature for 1 h. The temperature then was increased to a selected testing temperature, typically 600 °C, and held at that temperature for the duration of a series of tests. Cell performances were first determined using H2S as the anode feed. The cell open circuit voltage (ocv) and electrode potentials were monitored as functions of time on stream. Electrode potentials were recorded versus a Pt reference electrode using a National Instruments V-Bench system and other data were recorded with a Gamry PC4-750 measurement system. Since the reference electrode was located at the cathode side of the electrolyte membrane, the recorded change in anode potential included both the anodic electrochemical overpotential and the ohmic drop arising from the electrolyte resistance. After a steady ocv was achieved, potentiodynamic measurements were conducted at a scan rate of 5 mV/s to determine the current–voltage performance of the H2S–air fuel cell in the IR compensation mode using
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Fig. 1. Change in cell voltage and electrode potentials with anode feed (anode: Mo–Ni–S + Ag; cathode: Pt T: 600 °C, H2S–air feed rate: 25/25 ml/ min. Switching N2–H2S at time 0. Starting voltage scanning at time 2475 s).
1.4
35 I-V, Pt I-V, LSM I-V, NiO I-P, Pt I-P, LSM I-P, NiO
Voltage (V)
1 0.8
30 25 20
0.6
15
0.4
10
0.2
5
0 0
20
40
60
80
100
Power (mW/cm2)
1.2
0 120
Current (mA/cm-2)
Fig. 2. Performances for H2S–air fuel cell using different cathodes (T: 600 °C. H2S–air feed rate: 25/25 ml/min. Anode: Mo–Ni–S + Ag).
Anode potential (V)
-1.5
-1.3
-1.1
-0.9
-0.7 0
20
40
60
80
100
120
140
-2
Current (mA/cm )
Fig. 3. Current-potential curves for Mo–Ni–S composite anode in H2S stream (T: 600 °C; H2S/air feed rate: 25/25 ml/min; Cathode: NiO).
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the Gamry system. The performance of H2–air fuel cells was also then determined to investigate whether the nature of the anode feed affected the performance of these cathodes. 2.4. Surface analysis The electrode morphologies and compositions were studied with scanning electron microscopy (SEM) and energy dispersive X-ray (EDX) techniques using a Hitachi S-2700 scanning electron microscope and PGT Imix system with a PRISM IG or a JEOL 6301F field emission scanning microscope and an attached PGT Analyzer. In some cases, a layer of gold was sputtered onto the SEM samples to increase the electrical conduction. X-ray diffraction (XRD) analyses were conducted using a Rigaku Rotaflex X-ray diffractometer. XRD spectra were acquired and compared with samples of mixtures of various cathodes and electrolyte (Li2SO4 + Al2O3) with and without heat-treatment in air at 750 °C for 6 h.
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was scanned from 1.411 to 0 V, the anode potential apparently changed from 1.447 to 1.332 V, while the cathode potential changed from 0.036 to 1.332 V (Fig. 1). Thus, the cathode became more highly polarized (1.296 V) than the anode (less than 0.115 V), indicating that the cathodic kinetics were much slower than those at the anode. The recorded anode potential change included both the anodic polarization and the ohmic drop from the electrolyte resistance. This result differed from that for H2S–air fuel cells using oxide-ion-conducting YSZ as the electrolyte and
3. Results and discussion 3.1. Gas impermeability of the electrolyte membrane The membrane had to be gas impermeable so as to ensure that the data obtained during this research were free from effects arising from cross-over of gases. To serve this purpose, the cell voltage and both the cathode and anode potentials were monitored when the anode feed was changed from N2 to H2S and, during cell performance, were measured. Fig. 1 shows the typical experimental results obtained with the Li2SO4–Al2O3 electrolyte membrane used in our tests. It is obvious that the cathode potential was unaffected by the anode feed, and remained constant with time on stream, until voltage scanning was performed. In contrast, when a less dense Li2SO4–Al2O3 membrane was used, there was cross-over of H2S into the cathode chamber, which resulted in fluctuations of the order of hundreds of mV in the cell voltage and deterioration of the reference electrode. Thus, the data in Fig. 1 show that there was no cross-over in the present system using a dense Li2SO4– Al2O3 membrane. 3.2. Polarization of the Pt cathode and Mo–Ni–S based anode in the H2S–air fuel cell To discern which electrode limited cell performance, polarizations of both cathode and anode were monitored during current–voltage behaviour measurements of H2S–air fuel cells. When a Pt cathode and Mo– Ni–S based anode were used, and the cell voltage
Fig. 4. SEM of Mo–Ni–S anode (a) before and (b) after H2S–air fuel cell tests.
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Pt as the cathode, for which the cathode polarization was 200 mV and the anode polarization 400 mV at a current density of 130 mA/cm2 [26]. Thus, there was a dramatic effect on the performance of the Pt cathode attributable to the switch from oxide-ion-conducting YSZ to proton-conducting Li2SO4–Al2O3 as the electrolyte. Consequently, two additional, widely different cathode materials were also examined to determine whether the effect was due to the cathode, the electrolyte or the combination of these materials. 3.3. H2S–air fuel cell performance using different cathodes Fig. 2 shows in more detail the dependence of H2S– air fuel cell performance on the different cathode material used. The results were typical and reproducible for each of the specific cathodes. In each case the electrolyte was Li2SO4–Al2O3 and the anode catalyst was Mo–Ni–S intimately admixed with 5 wt% Ag. The maximum current densities and power densities achieved using the three widely different cathode materials were: Pt, 29 mA/cm2 and 10 mW/cm2; LSM, 60 mA/cm2 and 19 mW/cm2; and NiO, 96 mA/cm2 and 33 mW/cm2. Thus, the performance of the fuel cell depended strongly on the cathode used.
3.4. Performance of Mo–Ni–S composite anode in an H2S–air fuel cell Fig. 3 shows the current–electrode potential curve with a Mo–Ni–S composite anode as the working electrode on Li2SO4–Al2O3 as electrolyte with NiO as the counter electrode when using H2S as the fuel and air as the oxidant. The potential at the anode was controlled relative to that at the Pt reference electrode, which contacted the electrolyte from the cathode side. It showed a potential change (polarization) of 460 mV at 100 mA/cm2. Comparison of Fig. 3 with Fig. 2 shows that, at the same current density, the change in anode potential was much less than the change in potential for each cathode material. At the maximum current density, the anodic polarization was 57 mV with a Pt cathode (at 29 mA/cm2), 168 mV with LSM (at 60 mA/cm2) and 440 mV with NiO (at 98 mA/cm2), respectively. Meanwhile, the corresponding cathode polarizations were 1.353 V for Pt, 1.242 V for LSM, and 0.97 V for NiO. Therefore, in each case, the majority of the change in cell voltage occurred as a result of cathodic polarization. The activity of the anode under the present conditions was sufficiently high that the anodic reaction kinetics were not the limiting step.
Fig. 5. EDX of Mo–Ni–S anode (a) before and (b) after H2S–air fuel cell tests.
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3.5. Morphologies and composition of anode surfaces The surfaces of Mo–Ni–S anodes were examined before and after testing in a H2S–air fuel cell for several hours, using SEM (Fig. 4) and EDX (Fig. 5), to determine whether there were any morphological or compositional changes resulting from operation of the material in a fuel cell. The anode retained its highly porous microstructure. There were no significant changes in
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morphology (Fig. 4) or chemical composition (Fig. 5) arising from the fuel cell tests. 3.6. Changes at the cathodes during testing of H2S–air fuel cells The morphologies and compositions of different cathode materials before and after fuel cell testing were determined to provide insight into the causes of the high cathodic polarizations observed with Pt, LSM and NiO cathodes. Results are shown in Figs. 6–10 and Table 1. 3.6.1. Pt electrodes 3.6.1.1. Pt as cathode. The Pt cathodes were applied to the membrane as pastes and then heat-treated in air. Fresh Pt cathodes were very porous (Fig. 6(a)) and highly conductive, showing a dc ohmic resistance of less than 1 X, which arose mainly from contact resistances of the measuring probes. The performances of cells having Pt cathodes were very poor and deteriorated rapidly during use in H2S–air fuel cells. The used cathode was found to have developed a less porous structure (Fig. 6(b)) and showed very poor conduction, with the dc ohmic resistance being higher than 100 MX. EDX showed that the used Pt cathode comprised only 37a/o Pt, and had 20a/o S and 43a/o O, in contrast to the fresh Pt cathode (100a/o Pt) (Table 1). Pt was exposed only at
Fig. 6. Morphology of Pt cathodes (a) before and (b)–(c) after fuel cell tests: (b) overall; (c) detail. Compositions at Points 1 and 2 in (c) are in Table 1.
Fig. 7. LSM + Ag cathodes (a) before and (b) after fuel cell tests.
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Fig. 8. EDX results of LSM + Ag cathode (a) before and (b) after fuel cell tests.
isolated sites represented by site 1 in Fig. 6(c), and the cathode was under an extensive coverage by electrolyte materials represented by site 2. When these sites were compared, EDX showed that site 1 maintained a high Pt content of 81a/o, and site 2 had only 7a/o Pt but 35a/o S and 58a/o O. Thus, the performance of the used cathodes was affected detrimentally as the electrolyte blocked access of oxygen molecules to Pt catalytic sites and inhibited transport of electrons between different Pt particles, resulting in high polarization of the Pt cathode. In contrast, contacts between Pt cathodes on YSZ membranes were found to remain intact during prolonged operation of the fuel cells. Consequently, the effect is directly related to the nature of the Li2SO4–Al2 O3 electrolyte, as well as the Pt catalyst, as shown below. 3.6.1.2. Pt as anode. The performance and stability of Pt as the anode catalyst was also investigated for use with Li2SO4–Al2O3 electrolyte, to determine whether the above surface-covering effect arising during fuel cell testing was related only to the nature of the materials. An H2–air fuel cell using Li2SO4–Al2O3 electrolyte and operated at 600 °C was used instead of an H2S–air fuel cell, since it was known that reaction of Pt with H2S could cause delamination of the anode [16,17]. The EDX spectrum for a Pt anode after fuel cell tests showed only a single Pt peak. No signals were observed for Al,
O or S, which showed that the used Pt anode had not become covered by electrolyte, in contrast to the extensive coverage of the Pt cathode by electrolyte. Thus, transport of electrolyte onto the Pt electrode was not a function of the nature of either Pt alone or the combination of Pt and Li2SO4–Al2O3 electrolyte. Transport of electrolyte to cause covering of the Pt surface occurred only when Pt was used as the cathode during fuel cell testing. In contrast, no compositional change was observed in XRD spectra of Pt and electrolyte mixtures subjected to heat-treatment in air at 750 °C for 6 h under non-fuel cell operating conditions. Also, no covering by electrolyte was detected on Pt reference electrodes that were prepared from the same paste and located in the same cathode chamber during the fuel cell tests when the Pt cathodes became coated. Thus transport and covering of electrolyte were related directly to the electrochemical process(es) during fuel cell operation. It is unclear whether the local transport of electrolyte and its mixing with cathode materials arose through mobility of ions other than protons, as there was unlikely to be significant mobility over significant distances of ions other than protons [5,8,11,12]. 3.6.2. LSM cathodes LSM has been used extensively as the cathode catalyst in SOFCs using a variety of fuels. However, we have
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now found that similar electrolyte transport effects were observed with LSM cathodes when used with Li2SO4– Al2O3 membranes. LSM admixed with 5 wt% Ag was applied as a paste to one face of Li2SO4–Al2O3 membranes. The Ag enhanced the electrical conductivity of the LSM cathode, and was unaffected by O2 at the operating temperature of the fuel cell. The fresh catalysts were heat-treated in air, after which they were very porous (Fig. 7(a)). The elemental composition of the fresh cathode surface was entirely LSM and Ag (Au signals arose from Au sputtered as the electrical conductor for SEM and as the reference for EDX determinations) (Fig. 8(a)). In contrast, the surface of used LSM cathodes comprised mostly particles significantly larger than those for fresh catalysts (Fig. 7(b)). The large surface particles on the used LSM cathode formed a much denser surface that was less accessible by O2. EDX spectra showed only traces of La, Mn, Sr and Ag in this layer (Fig. 8(b)). The strong presence of sulfur peaks indicated that the cathode surface had been covered almost entirely by the migrated sulfate. The cathode agglomeration effect appeared to be directly related to the transport of electrolyte, possibly due to adherence of particles caused by the coating of electrolyte. Thus the
Fig. 9. NiO cathode (a) before and (b) after fuel cell tests.
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performance of LSM cathodes was affected detrimentally both by coverage of catalyst sites by the electrolyte and by reduced access by O2 to the available catalyst sites. Mixtures of LSM and electrolyte behaved in a manner similar to mixtures of Pt and electrolyte. Mixtures that had undergone heat-treatment in air at 750 °C for 6 h showed no compositional change in XRD spectra. In contrast, there was even more severe covering of LSM by electrolyte after fuel cell tests than was found for Pt (Figs. 7, 8). 3.6.3. NiO cathodes In contrast to either Pt or LSM cathodes, comparison of fresh and used NiO cathodes showed only a small difference in surface morphology (Fig. 9). EDX analysis showed that the composition of the surface of used NiO catalyst was mainly similar to that of fresh catalyst (Fig. 10), which indicated very little transport of electrolyte materials onto the surface of NiO cathodes. Consequently, the performance of fuel cells having NiO cathodes was better and more stable than that for fuel cells with Pt or LSM cathodes (Fig. 2). Although it was found that cathode NiO particles agglomerated slightly during use, the cathode layer remained porous, and so the catalyst sites remained accessible by O2 (Fig. 9). Because there was no surface coating of electrolyte on the NiO particles, agglomeration of catalyst particles was not a consequence of electrolyte material transport, in contrast to the effect found for Pt and LSM (Fig. 7). Lithium doping has been shown to improve the electrical conductivity, chemical diffusivity, and oxygen exchange kinetics of NiO [27]. During operation of the present fuel cells, Li ions could have doped into NiO particles located adjacent to the electrolyte membrane, instead of covering its surface, and thereby could have enhanced both the ionic conductivity of the catalyst and oxygen reduction kinetics. However, no change was detected in XRD spectra of the NiO cathode after the fuel cell test, but this may have been a consequence of the inability of EDX to detect Li. A combination of NiO powder with electrolyte that was heat-treated in air also showed no change in the XRD spectrum. To summarize: the performance of proton-conducting fuel cells having Li2SO4–Al2O3 membranes depended strongly on the compatibility of the electrode materials with the Li2SO4–Al2O3 membrane material. Pt and LSM cathodes were unsuitable for use with Li2SO4–Al2O3 membranes during operation at temperatures at 600 °C, as the electrolyte material migrated onto the cathode catalyst surface. However, neither NiO cathode catalysts nor Pt or Mo–Ni–S anode catalysts became covered with this electrolyte. Thus, we have shown that the use of Li2SO4–Al2O3 as the electrolyte in proton-conducting fuel cells depends on the selection of cathode materials.
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Fig. 10. EDX of NiO cathode (a) before and (b) after fuel cell tests.
Mo–Ni–S anode catalysts and the electrolyte transport effect appears to be limited to the cathode.
Table 1 Surface atomic composition in at% of a used Pt cathode (EDX) Analysis area
O
S
Pt
Overall (Fig. 6(b)) Site 1 (Fig. 6(c)) Site 2 (Fig. 6(c))
43 9 58
20 10 35
37 81 7
4. Conclusions Transport of electrolyte materials onto the surface of Pt or LSM cathodes adversely affected the performance of proton-conducting fuel cells using Li2SO4–Al2O3 as the electrolyte and operating at temperatures at 600 °C. In contrast, no significant transport of electrolyte onto the cathode was observed when NiO was the cathode catalyst. Covering the cathode surface with electrolyte reduced the accessibility to catalyst sites by O2 and reduced the electrical conductivity of the cathode layer; consequently the covered cathodes became highly polarized. In contrast, NiO remained mainly uncovered after use. Consequently, H2S–air fuel cells having NiO cathodes showed the highest cell performance, with a current density of over 100 mA/cm2 and a power density of over 30 mW/cm2. Slight agglomeration of NiO particles did not appear to affect the performance of the cathode. The same electrolyte did not migrate onto either Pt or
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