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A new solid oxide fuel cell design based on thin film electrolytes
$3.00 + 0.00 03~5442/90 Copyright @ 1990 Pergamon Press plc
Energy Vol. 15, No. 1, pp. 1-9, 1990 Printed in Great B&am. All rights reserved
A NEW SDLID OXIDE NEL
CELL DESIRR BASED OISTHIN FILM ELECTROLYTES
S.A. BARNETT Department of Materials Science and Engineering, Northwestern University, Evanston, IL 60208 (Received 11 July
1989)
Abstract - A novel solid oxide fuel cell (SOFC) design that can be fabricated entirely using low-temperature, thin-film processing is described. Potential advantages of the cell are reduced materials costs and improved fuel-cell characteristics. The crftical design feature is the use of thin (% 50 nm), catalytically-active oxide layers on a < 10 em thick yttria-stabilized tfrconia (YSZ) supported electrolyte to minimize reaction overpotentials and ohmic losses. Doped cerfa at the fuel electrode side and doped bismuth oxide at the oxygen electrode side are proposed for the surface layers. The surface reaction rates and overall electrolyte conductance in this design are high enough at < 75O'C to allow efficient SOFC operation. This operating temperature is low enough that low-resistance, thfnfilm metal electrodes, Ni at the fuel side and Ag at the oxygen side, can be used to provide low ohmic losses. The overpotential behavior of the proposed cell is estimated on the basis of literature data an leads to fuel efficiencies > 50% at a power density of I, 0.5 W/cmP when operated at 750°C. The thin film monolithic design will lead readily to the incorporation of the cells into stacks with high power-to-weight and power-to-volume ratios. Methods for cell fabrication are discussed. 1.
INTRWCTIOlY
Fuel cells possess significant advantages for power generation including electrical conversion efficiencies of 40 to 60%. 90% overall conversion efficiency in pofnt-of-use cogeneration of electrfffty and heat, independence of efficiency on power plant size, and Amongst the various fuel cell schemes, solid oxide fuel cells ;e&fble pollution. are believed to have the widest potential range of applications and hi hest system efficiencies because of their all-solid-state design and hfgh (~lOOO°C s operating temperaturesITc which lead to internal reformfng of fuel gases and efficient utilization of waste heat. Successful prototype SOFCs based of 3~ttrfa-stabilized zfrconfa (YSZ) electrolytes are currently under further development. The economic viability of SOFC power generation hinges on processing costs, materials costs, and cell performance. While the operating characteristics of exfstfng SOFCs are good, it is not yet clear that costs will be competitive. In particular, the use of high processing temperatures fncreases fabrication costs while materials costs are relatfvely high since thick (1 a$Bfilms of rare-earth-containingcompounds, e.g. LaMnO , are used as the oxygen electrode. It is thus of interest to consider SOFC-design opsions that may reduce costs or exhibit enhanced characteristics. An alternate approach that would reduce materials costs is an all-thin-film SOFC desfgn. Low temperature thin-film deposition methods such as magnetron sputtering are available that provide good control over composition, properties, and morphology. Since magnetron sputter deposition is commonly used as a large-scale production technique, one might expect processing costs to be relatively low. Detailed studies of cell processing will be required before reliable cost predictions can be made. The reduced amount of material in a thfn film cell would also cut cell mass leading to higher power-to-mass ratios. However, the feasibility of a thin-film SOFC has not been established. Two technical advances are required to make thin film SOFCs a viable approach. First, methods for fabricating supported thin film structures must be developed. Second, a thin film oxygen electrode is needed. In SOFCs operating at Tc4=510000C, various stability crfterfa dfctate the use of low conductivity oxide cathodes l that must be-1 tmn thick in order to reduce ohmic losses to tolerable levels. In order to make a thin film SOFC, it m ,I&&
1
S. A.BARNFIT
2
is thus necessary to reduce Tc sufficiently to allow the use of a high conductivity metal cathode, namely Ag, that eliminates ohmic losses at thicknesses of ~1 pm. Since the volatility of Ag becomes a problem at T, > 750°C, technical feasibility depends on being able to achieve acceptable reaction overpotentials and electrolyte ohmic losses at this temperature. In thfs article, the technical feasibility of a supported thin film SOFC operating at < 750°C is discussed. A detailed, thin-film SOFC design is developed to allow estimation of the expected low T operating characteristics. The design is based on a thin (< 101.1m), YSZ-supported electrolyte that provides very high oxygen conductance. The YSZ is coated with thin oxide layers in order to reduce reaction overpotentials. Likely candidates are doped cerfa on the fuel electrode side and stabilized bismuth oxide on the oxygen electrode side. Estimates based on reaction rates for these materials indicate that the overall cell will operate efffcfently at temperatures T, as low as 700°C. In addition to allowing the use of thin Ag films for the oxygen electrode, the low T, also leads to higher thennodynamfc efficiencies.and a wider range of allowable reducing fuel mixtures. The predicted operating characteristics of the proposed thin-film SOFC design are at least as good as existing cells. The thin-film approach would thus seem to merit further research. 2.
THIN FILM SOFC DESISN CRITERIA
The main limitations on state-ofithe-art SOFC performance are oxygen electrode ohmic loss and activation overpotential. Ohmic losses are important despite electrode thicknesses of 1 mm because of the relatively poor conductivity of the oxide cathode. While a metallic electrode with much higher conductivity would reduce ohmic losses even for thicknesses in the micron range, economically-viablemetals are not sufficiently stable at T = 1ooooc. The proposed design thus depends on being able to reduce T sufficiently in t6e thin-film cell in order that a metallfc cathode, specifically Ag, can 6e used. In this section, the main factors limiting T, values - reaction overpotentials and ohmic loss in the electrolyte - are discussed. It is concluded that T, can be reduced sufffclently to allow the use of Ag cathodes. A. -
Electrode reaction kinetics
Slow reaction kinetics at the fuel cell electrodes can lead to large interfacial resistances and hence significant losses at high current densities. Figure 1 is a comparison of the low voltage interfacial resistance Rf in ai6 8s a function of temperature for several different electrolyte/electrode combinations. In the case of a YSZ Temperature, T, ("C) 1200
1000
600
600
500
Fig. 1. Interfacial resistances in air as a function of inverse temperature
for several different electrode-electrolyte Data were taken from Refs. 6-8.
combinations.
The the dependence on electrode material and morphology is very strong. electrolyte, shaded reafon shows the ranoe of R; values obtained for Pt electrodes with different morphologfes. While there is-a wide Lange of values observed, the ~10~ is independent of
Results for the moroholoav. fndfcatfna that the rate-limitina mechanism is the same. mixed co&ctfng oxide La 6Sro 5Co03 are also shown and are in the same range as observed for Pt. Much lower resfs ?*ances'are observed for Ag on YSZ and the dependence on morphology is considerably reduced. In all cases, R. increases rapidly with decreasing temperature and the resulting overpotential has precludea attempts to reduce T, in YSZ-based cells.
New solid oxidefuelcell design
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Substarftial decreases in R. can be obtained by using a stabilized bismuth-oxide electrolyte. Results are shown i'nFig. 1 for Pt on 20% Er-stabilized bismuth oxide (ESB) and Ag on 20% Ba-stabilized bismuth oxide (BSB). For doped bismuth oxide, the oxygen reaction has been shown to take place directly on the electrolyte quiface, resulting in a For example, mixed very low Ri value and only a weak dependence on electrode material. ’ exhibited only slightly higher Ri values than Ag. conducting oxide electrodes on BSB Clearly, the use of doped bismuth oxide as the cathode rather than YSZ would reduce overpotentials. It should be noted thlt other catalytically-activeoxides such as U02 have been successfully used to decrease Ri. Catalytic properties for H20 reduction are be{der at ceria electrolytes than YSZ leading to lower overpotentials in electrolysis CellS. As in the case of bismuth oxide, the higher rate kinetics are apparently due to a reaction on the electrolyte surface. SOFC electrolyte applications for gadolinia-doped ceria (GOCj11a12compositionexhibiting high oxygen conductivity, have attracted considerable interest. ’ The applicability of ceria and bismuth oxide electrolytes is limited, however. Bismuth oxide is unstable in reducing envim ts and the use of ceria in fuel cells is On the other hz#, ceria-doped surface problematic due to low transference numbers. ’ layers on YSZ have been used successfully in reducing environments. ":hi: electronic conduction through the electrolyte. Doping with bismuthThoex:dSeZ1yr~~enat: oxygen electrode side of the electrolyte has also been successfully used to enhance reaction rates. These results suggest that a multilayer electrolyte, for example GDC/YSZ/BSB, would minimize activation overpotentials while still retaining chemical stability and 100% ionic conduction. Figure 2 shows the effective oxygen partial pressure as a function of position three-layer electrolyte with air on the cathode side and an 0 partial pressure P = 36 atm (a hydrogen to water partial pressure ratio of 500 at 700$), for example, on ? he fuel side. Conductivity data for the three electrolytes at 700°C was used in constructing this figure (see section 1I.B below). As long as the BSB layer was sufficiently thin ,GDC
BSB \
25
0
I
fuel electrode
Electrolyte Depth (urn)
I
air electrode
Fig. 2. Oxygen partial pressure as a function of distance in a thre -layer electrolyte at 700°C with air on one side and PO = 10-56 atm on the other. Electrolyte resistances were taken from Refs. 8, 16, and 17. relative to the YSZ, it would be exposed only to an oxidizing environment and hence remain stable. The GDC would remain chemically stable in the reducing fue$flectrode environment. but would exhibit a significant component of electronic conduction. The GDC layer would thus also have to be quite thin relative to the YSZ in order to prevent shorting of a significant portion of the cell voltage. Nonetheless, very thin GDC and BSB layers should be sufficient in order to enhance surface reaction rates. The use of very thin layers combined with the good thermal expansion match between i@ three oxides eliminates any thermal stability problem for the three-layer electrolyte. As discussed in section IV. the deposition of this three-layer electrolyte will not significantly complicate the cell
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Cell Temperature. 900
850
800
T, (‘C)
750
700
650
'Orn
Fig. 3. Air and fuel electrode interfacial voltages at a current density of 1 A/cm2 as a function of inverse temperature. Data were taken from Refs. 8 and 9. SOFC
11001000so0
Operating Temperature, T, (OC) 800
1000/l (K-l) Fig. 4. Electrolyte thickness at which the ohmic voltage dsop across the electrolyte is 0.05 V for a current density of 1 A/cm , plotted vs. inverse temperature for three electrolytes. Resistivity data were taken from Refs. 8, 16, and 17.
New solid oxide fuel cell design
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processing. Existing data indicate that overpotentials low enough for efficient SOFC operation can be achieved at T < 75O'C. For example, resistance losses in a SOFC with a Ni/Ybstabilized ZrO (ebSZ) anode and (In,Sn)oxide/U-dopedYbSZ cathode we% dominated by the Fie. 3 is a plot 0.4 mn thick Y& electrolyte conductivity, not Ri, at T down to 8OO'C. vs. T, of anode and cathode overpotentials for a currenE density9J of 1 A/cm . Data for the anode was taken for a Ni cermet on YbSZ in a 50% H -H20 mixture. (No data of this type is available for GDC.) Data for the Ag/BSB cathode are extrapolated from the resistance results shown in Fig. 1. Also shown are data where UO2 wab added to the YbSZ electrolyte Each of these combinations to enhance reaction rates at a In 0 -SnO air electrode. yields an acceptable overpotential aQ$, %ooc. B_
Electrolyte resistance
Electrolyte resistance is currently not a limitation on SOFC performance since the electrolyte layers are < 50 Am thick and T, = 1000°C. However, at T, ?r 700 to 750°C, electrolyte resistance becomes a limiting factor. Figure 4 is a $lot of the electrolyte thickness t at which the ohmic voltage dro&is 0.05V at J = 1 A/B versus 1000/T , based on conductivity data for YSZ (10% yttria), GDC (10% gadolinia), and BSB (20% baria). Since the GDC and BSB layers are quite thin in the proposed electrolyte and have conductivities similar to the YSZ, they will make negligible contributions to the overall resistance. The electrolyte must be < 10 Pm thick in order to keep ohmic losses < 0.05V at 7OO'C. Lower voltage drops can be achieved by using thicknesses < 10 #II and/or T, > 7OO'C. In addition, the electrolyte conductivity must be due mainly to the ionic contribution di, with only a minor conductivity~~. That is, the YSZ ionic transference number t be near unity in order that electrical conduction does not Fig. 5 shows the range of T, and PO values where ti > 0.99 for stabilized zirconia. Efficient fuel cell operation is possible Temperature 02) 16 6
-32
1 Ca-
stabilized ZrO,
3.0
5.0
7.0
9.0
11.0
13.0
1O"lT (K-l)
Fig. 5. A plot of the temperature and oxygen partial pressure range where the conductance of stabilized zirconia is primarily ionic, i.e. where the ionic transference number is > 0.99. Data were taken from Ref. 18. anywhere within this region. Lowering T, from 1000 to c 700°C has the advantage of expanding th~of~olI;!~~trode PD range that can be used without electronic conduction by a factor of 10 atm, as shown in Fig. 5. The low P8 value at the fuel electrode shown in Fig. 2 is thus acceptable for a cell operated at 700 C. C,
Oxygen Electrode Material
A SOFC based on a three-layer electrolyte < 10 Pm thick would be expected. based on the above discussion, to operate with acceptable overpotentials and electrolyte ohmic losses at T, < 75O'C. In this temperature range, Ag is a desireable alternative to conducting oxides for the air electrode since it 4ceines the good catalytic activity of the noble metals with high oxygen solubility. l Furthermore, Ag is an excellent conductor such that z 1 Pm thick Ag electrodes would be sufficient to eliminate ohmic losses, and hence materials cost would not be excessive. Ag has alwaylgbeen disqualified for use in high-T, cells because of its volatility. The Harwell group found that 2 Pm of Ag would be lost after 1 year in air at BOD'C. Extrapolations based on vapor-pressure data indicate that only 50 nm would be lost in 1 year at 700°C and 0.3~m at 750°C. Furthermore. in a thin-film approach it would be
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strafghforward to deposit a mixed-conducting oxide cap layer on the Ag electrc ,de in order to reduce Ag evaporation. This would oennft small 'increases in T_ above 7'50°C, and a consequent reduction in overpotentials. Finally, since 700°C is a siinfficant fraction of the Ag melting point (960°C), it will be necessary to co-deposit an oxide (i.e. BSB) with the Ag In order to minimize sfntering. 3.
THIW FILM SDFC DESIGW MD
PERFDRMWE
In this section, details of the design of a thin-film SOFC are consfdered and the performance of the resulting cell predicted. Fig. 6 shows schematically the SOFC design. Air
Fig. 6. A schematic drawing of the proposed thin-film fuel cell. Note that the drawing is not to scale.
In addition to the three-layer electrolyte and the Ag-cermet cathode, a Nl-cermet anode is proposed I-VSZ cermet is currently successfully used in SOFCs as the fuel electrod;.1'3*TI The chemical compatabilfty of Ag and Ni with the electrolytes is good. based upon their previous extensive use in solid electrolyte cells. The overall SOFC structure is supported on a Ni mesh frame which also serves for current collection. Techniques for fabricating this structure are discussed in section IV. A_
Cell dimensions
The maximum thickness t 1 of the electrolyte for T, = 700°C is 10~ m. The minimum thickness, on the other ban%, is given by the requirements that the electrolyte does not break down at operatfng voltages, that the film is stable against Interdiffusion, that it possess good mechanical stability, and that it fully covers imperfections on the substrate surface during deposition to form a pore-free fflm. Electrical breakdown can be discounted for fflms thicker than ~10 nm since breakdown fields are > lo6 V/cm and cell voltages are Sl v. Interdiffusion should not be a problem fo50tel > 500 nm and T, s700°C where oxide Finally, previous results have shown films have been found to be stable for long times. that 0.5 u 91is thick enough to obtain a pore-free electrolyte layer on a variety of substrates. Based on the above arguments, the overall electrolyte should be 0.5 to 10 urn thick. Mechanlcal stability and economic considerations will likely determine the optlmal thickness. Both the cell thickness and the free-standing film area (determtied by the Nf support frame mesh size) can be varied to achieve the required mechanical stability. There is considerable evidence suggesting that a mechanlcally stable SOFC can be based on thin ceramic films. which are sufficiently robust and resistant to fracwe and fatigue that a variety of coarnercfalelectromechanical devices are based on them. For example, freestandfng films such as cantilever beams up to 120 vrn long of 100 - 900 nm thick SfO2, en fabricated and reliably used as high frequency and 7059 glass ha$$,,#$ si3Nfg Nb2059 From an economics point of view, lower thicknesses oscf lators and pressure sensors. are desfreable to minimize processing time when techniques such as magnetron sputtering are used for film deposftlon (see section IV). Detailed studies of thin-film mechanical properties ~111 be required to determine an optimal electrolyte thickness. Electrode film thicknesses can be estimated using the criterion that < 10% of the to 1 cell volvge be dropped across the electrodes. The lateral voltage drop V 1 Is given by" V,l = JeL /2tf where J is the current density, P the electrode resistivfty, tf the
New solidoxide fuelcell
design
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film thickness, and L the distance between current carrying wires. P for Ni and Ag were taken at 700°C, giving p i = 23.3 n cm and PA = 5.2 ncm. These values are multiplied by a factor of 5 to accoun'jcfor the porous ce&et nature of the electrode films. Sputter deposited and annealed cermet films with metal volume fr#c#dons > 50% have been shown $0 J is taken to be 1 A/cm . have p values within a factor of 5 of the pure metals. * Taking, for example, tf = 100 nm gives L = 1 mn for Ni and 2 mm for Ag. A Ni foil mesh would be used as the mechanical support for the fuel cell and also as a network for current collection. Since the free-standing electrolyte area will likely be limited to < 1 mm by mechanical stability criteria, a 100 nm Ni fuel electrode would essentially eliminate ohmic losses at that electrode. On the other hand, a thicker ( % 2 Pm) Ag film would be useful in order to achieve L = 1 cm. The Ni-foil mesh size would range from 0.1 to 1 mm , with the upper limit due to electrolyte mechanical stability and the lower limit due to ease of fabrication. A foil thickness of 2511m would provide sufficient rigidity over a 1 cm width while minimizing materials cost. The percentage open area of the mesh, a compromise between maximizing the cell area and obtaining sufficient rigidity and conductivity, should be s 70%.
2B
Predicted SOFC performance
The proposed cell has negligible ohmic losses at the electrolyte (for T, = 75O'C) and Ni electrode (for t&la> 200 nml and a voltaoe droo of 0.05 V at the Aa electrode for J = 1 A/&The major "sburcesof *losses in ce'iloperation are the overpotentials associated with kinetic rate limitations at the electrodes. The overvoltages at both electrodes at 750°C range from 0.07 to 0.25 V (Fig. 3). A conservative estpte of the overpotential is thus 0.3 V, giving s 0.35 V total voltage drop at 1 A/cm . Based on these losses, a predic$ion of cell performance can be made. Using the 2/O reaction potential of 1.1 V at 750°C, this gives a cell voltage of% 0.75 V at 1 A/cmv ._ ii oting the fact that the actual cell open area is 70X, the power density is ?. 0.5 W/cmL with a fuel efficiency;ff %50X. This compares favorably w$th the Westinghouse cell operftfng on H2 at 0.5 A/cm with a power density of 0.33 W/cm and a fuel efficiency of 46%. 4. A -I
TRIM FILH SOFC FABRICATIOW.INTERCORNECTS, ANDCELL STACKS
Fabrication
Thick-film electrolytes in state-of-the-art SOFC designs are deposited directly onto porous s bstrates using the electrochemical vapor deposition (EVD) technique to fill the pores.letl In this case,there is a limitation that the minimum film thickness be on the order of the substrate pore size, i.e. %5O#m, in order to fully cover all pores. To work with 0.5 to 10um thick electrolytes, it will be necessary to deposit on fully dense substrates that can subsequently be removed or rendered porous. While many schemes are possible, the following seems to be a simple, economically-feasible approach. To fabricate the proposed cell, the entire cell structure, including Nf-cermet. GDC, YSZ. BSB, and Ag-cermet. will be sequentially deposited directly onto a continuous Ni mesh/polymer sheet. This substrate could be formed by lamination or by coating the Nf mash with a liquid polymer. In the latter case, a liquid solution that fills the pores in the mesh due to surface tension would be dried to provide a solfdz5dense substrate. Polymer solutions such as commonly used in semiconductor lithography would be suitable since they can withstand the temperatures ( <400°C) that would arise during deposition. The polymer would be removed, a er depositing the cell, either by wet chemical etching or ashing in an oxygen plasma>% A variety of thin-film deposition techniques are available for fabricating the structure shown in Fig. 6. The present discussion will center on the use of magnetron sputter deposition for preparing the cells. Magnetron sputtering is chosen because it can be used to deposit all of the cell materials at low temperature in a single production-line deposition chamber, it gives high deposition rates and high materials utilization efficiency, and provides excellej&,2fjontrolover film thickness, structure, intrinsic stress, adhesion, and composition. The oxide films would be deposited by reactive magnetron sputtering of Zr-Y, Ce-Gd and Bi-Ba metallic targets in Ar-O2 mixtures. Deposition rates of 150 nm/min have been reported for reactive magnetron sputtering of Zr02 such t&t, for example, only ?r6 mins would be required to deposit a lr m thick electrolyte. Nf-Ce-Gd and Ag-Bi-Ba would be cosputtered in Ar-O2 mixtures to form the cennet electr #es. Rates for Ag and Ni deposition are expected to be even higher than for the oxides. The deposition chamber would be a standard in-line sputtering system with 5 sputtering targets and a mechanism for moving substrates continuously beneath the targets. Substrate preparation and loading into the deposition chamber should be carried out in a clean room in order to eliminate dust particles which can lead to pinholes in the films. Depending upon the growth conditions, sputtering can be used to deposit either dense or porous films. s een widely used for depositing porous metal electrodes Sputteri?o for electrochemical cells. * 9!n,zB On the other hand, many groups have shown in the last
S. A.BARNER
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several years that low energy ion bombardment of the substrate and growing film during obtain high density oxide films such as Zr02 and CeO2 at B z-
Cell interconnectionand stacking
The present monolithic cel1Idfygn would be suitable for incorporation into stacks as proposed by the Argonne group. ’ The basic cell stacking configuration is shown schematically in Fig. 7. In this approach, the Ni/polymer substrate would be formed into lntercon"ecl z?zz hl0leell uxmmnmm
air fuel
fuel
Fig. 7. A schematic drawing of the stacked fuel cell and interconnect. the V-shaped configuratlon shown and mounted on a ceramic cell support structure. The films making up the SOFC would be deposited directly onto this substrate. After removal of the polymer, the interconnect sheet would be added and the cells stacked. The fnterconnect would be a Ni sheet coated on the side exposed to oxygen with a conducting oxfde with low oxygen d ffusfvity to prevent oxidation of the Ni. iaCrOS, used successfully In existing SOFCs '-' would be a suitable choice. A separate sputter deposftfon system would be used to deiosft the oxide layer on the Ni sheet to form the interconnect. 5.
DISCuSSIDll
The thin-film SOFC overpotentials estimated fn section 1II.B indicate that for operatlon near 75O'C. the cell power density and efficiency would be as good or better than existing cells. Furthermore, it may be possible to increase the useful cell Tc range slightly. and hence fmprove cell performance, by capping the Ag cathode wfth a mfxedconducting oxide in order to reduce evaporation (see section 1I.C). An all-thfn-film deslgn employing a Ag cathode is thus not only feasible from an electrochemical pofnt of view, but may provide Improved efficiency and power density over existing SOFCs. The thin-film approach offers the advantages of monolithic cell design, reduced materials costs, higher power-to-weight and power-to-volume ratios relative tf existing SOFCs. Estimates were made for cell stacks as proposed by the Argonne group. The main materials cost is for the Ni (taken to be $ZO/kg) structural supports and interconnect, with a fraction due to the Ag (taken to be $250/kg) films and connectfons. A very conservative estimate gives a materialsScost cjf< $2O/kW. A power-to-weight ratio of > 2 kW/kg and a power-to-volume rati? > 10 kW/m are estimated for this type of cell stack using a power density of 0.5 W/cm . The lower T, value of the present SOFC design offers a few potential advantages. First, less insulation would be required for cell stacks to be thermally self-sufflcfent. a 'ons such as synthesis reactions, lower temperatures will Second, in non fuel cell appl$$,$3 Third. the low T value and low thermal mass of the provide improved selectivity. thin-film SOFC suggests the possibility that a short-ctrcufted cell can heat itself rapidly from an "idling" temperature to T,. A simple calculation based on a cell current limited by the oxygen electrode resistance (as expected for T, < 700°C based on the data presented above) and assuming a 5 pm total cell thickness yields a 1 degree/set heating rate at 400°C. and 10 degree/set at 450'. In transportation applications Involving intenaittant use, a well-insulated cell stack could thus be maintained at 400°C by burning very small quantities of fuel and be able to reach operating temperature in -1 min. The above discussion suggests that a thin-fflm approach to SOFC design merits further investigation. Considerable research and development would be required to make thls a viable technology. First, the fabrication processes described above and the resulting films would have to be characterized in detail. Second, detailed fnvestlgatfons of the
New solid oxidefuelcell design
9
mechanical stability of the thin-film structure would be needed. One issue is that the structure have sufficient toughness to withstand any gas pressure gradients or impacts that This requirement will impose limits on the Nf mesh it would be subjected to in normal use. dimensions and on electrolyte film thickness. In addition, a partially-stabf#red Zr02 could be electrolyte, known to exhibit much better toughness than fully-stabilized ZrO e partiallyused to improve mechanical stability. The lower fonic conductivity of t8' stabilized ZrO2 would not significantly degrade cell characteristics since they are not limited by electrolyte resistance. Another important issue is the stability in thermal cycling applications. While the thermal expansion coefficients of the Nf components, ceramic cell supports (presumably made of stabilized ZrO ), and_.,e;ectrolytesare well yatched, Ag has a relatively large mismatch with the oxides 2i1.9x10 / C for Ag VS. 1.2~10 /OC for YSZ). While the extent of the problems posed by this mismatch are not known, the First, the fracture toughness of ceramic-based following points should be noted. multilayer structyges with different thermal expansion coefficients improves as the layers favoring the present thin-fflm design. Second, it may be possible to are made thinner, minimize thermal strains by minfmiring the temperature cycling range as described above. Thfrd, most of the thermal strains will taken up in the porous, low elastic modulus Agcermet layer. High temperature fatigue in the Ag would thus be the likely thermal cycling failure mechanism. Acknowlecf~nts - We thank the Donors of The Petroleum Research Fund, administered by the American Chemical Society, for support of this research. REFERENCES J.T. Brown, Energy 11, 209 (1986). A.O. Isenberg. Solid State Ionics 3/4. 431 (1981). J.T. Brown and S.E. Veyo, Mod. Power Syst. 6, 43 (1986). C.S. Tedmon, H.S. Spacil. and S.P. Mitoff, 7. Electrochem. Sot. 116, 1171 (1969). 0. Yamamoto, Y. Takeda. R. Kanno, and M. Noda, Solid State Ionfcn2, 241 (1987). J. Sasaki. J. Mfzuasaki. S. Yamauchf, and K. Fueki, Solid State IosiTcs 3/4, 531 (1981). 70 (1983). 7. M.J. Verkerk, M.W.J. Hamnfnk, and A.J. Burggraaf, J. Electrochem. Sot. m, 8. T. Suzuki, T. Yamazaki, K. Kaku, and M. Ikegami, Solid State Ionics I5,xl (1985). 9. R. Accorsi and E. Bermann, J. Electrochem. Sot. 127. 804 (1980). L.J. Olmer. J.C. Viguie. and E.J.L. Schouler, Sam State Ionfcs I, 23 (1982). ::: T. Kudo and H. Obayashi, J. Electrochem. Sot. 123, 415 (1976). H.L. Tuller and A.S. Nowfck, J. Electrochem. Sot. 122, 255 (1975). :23:A.J.A. Winnubst, A.H.A. Scharenborg. aud A.J. Burggraaf, Solid State Ionfcs 14, 319 (1984). Kflner, P.F. Dennis, A.E. McHale, M. Van Hemert, and A.J. 14. 8.C.H. Steele, J.A. Burggraaf, Solid State Ionfcs 18 19, 1038 (1986). 15. S. Majumdar, T. Claar, and 8. I+andermeyer, J. Am. Ceram. Sot. 69, 628 (1986). 8.C.H. Steele, Solid State Ionfcs I& 391 (1984). :;: T. Kudo and H. Obayashi, J. Electrochem. Sot. 122, 142 (1975). of Electronic Ceramics, p. 143, L.L. Hench and D.8. Dove eds., 18. J. Patterson, in b Marcel Dekker, New YorkT971) T.L. Markfn, Power Sources Q, i. 583, D.H. Collins ed., Oriel, New York (1973). Ch. 3, S.M. Sze ed., McGraw-Hill, New York (1983). ::: A.C. Adams, inITechnolosy, L.B. Welsh, and F.R. Szofran, J. Vat. Scf. Technol. I& 177 21. J.E. Greene, R.mlinger, (1977). K.E. Petersen, Proc. IEEE, 70, 420 (1982). f3: K.E. Petersen and C.R. GuarSireri,J. Appl. Phys. 50, 6761 (1979). 8. Abeles, P. Sheng, M.D. Coutts, and Y. Arie, Adv. Phys. 2. 407 (1975). Ht: S.K. Ghandhi, in VLSI Fabrication Princi les, Wiley, New York (1983). 26. See, for example,TE. Sundgren,-EX&- o ansson, A. Rockett, S.A. Barnett, and J.E. Greene, in Chemistry and Physics of Hard Coatings, p. 95, B. Sproul, J.E. Greene, and J.A. Thornton eds., AKNew York-m 27. J.A. Thornton, in Thin Film Processes, p. 76. J.L. Vossen and W. Kern,Academfc, New York (1978). F. Jones, J. Vat. Sci. Technol. A 5, 3088 (1988). z: T.M. Gur. I.D. Raistrick, and R.A. Huggfns, J. Electrochem. Sot. 127, 2620 (1980). 30. J.E. Greene, S.A. Bamett, J.-E. Sundgren, and A. Rockett, in _Ion Beam Assisted Film Growth, Ch. 5, T. Itoh ed., Elsevier, Tokyo (1988). Martin, J. Mater. Sci. 21, 1 (1986). and references contained therein. 31. mP.J. T.M. Gur and R.A. Huggins, SolidState Ionics 5, 563 (1981). R.D. Farr and C.6. :32:C.T. Sigal and C.G. Vayenas, Solid State Ionics 5, 567 (1981); Vayenas, J. Electrochem. Sot. 127. 1478 (1981). 34. T. Yoshida, T. Hoshfna, I. Mukxawa, and S. Sakurada, J. Electrochem. Sot. 136, 2604 (1989).
::
2 2
Energy Vol. 15, No. 1, pp. 1-9, 1990 Printed in Great B&am. All rights reserved
A NEW SDLID OXIDE NEL
CELL DESIRR BASED OISTHIN FILM ELECTROLYTES
S.A. BARNETT Department of Materials Science and Engineering, Northwestern University, Evanston, IL 60208 (Received 11 July
1989)
Abstract - A novel solid oxide fuel cell (SOFC) design that can be fabricated entirely using low-temperature, thin-film processing is described. Potential advantages of the cell are reduced materials costs and improved fuel-cell characteristics. The crftical design feature is the use of thin (% 50 nm), catalytically-active oxide layers on a < 10 em thick yttria-stabilized tfrconia (YSZ) supported electrolyte to minimize reaction overpotentials and ohmic losses. Doped cerfa at the fuel electrode side and doped bismuth oxide at the oxygen electrode side are proposed for the surface layers. The surface reaction rates and overall electrolyte conductance in this design are high enough at < 75O'C to allow efficient SOFC operation. This operating temperature is low enough that low-resistance, thfnfilm metal electrodes, Ni at the fuel side and Ag at the oxygen side, can be used to provide low ohmic losses. The overpotential behavior of the proposed cell is estimated on the basis of literature data an leads to fuel efficiencies > 50% at a power density of I, 0.5 W/cmP when operated at 750°C. The thin film monolithic design will lead readily to the incorporation of the cells into stacks with high power-to-weight and power-to-volume ratios. Methods for cell fabrication are discussed. 1.
INTRWCTIOlY
Fuel cells possess significant advantages for power generation including electrical conversion efficiencies of 40 to 60%. 90% overall conversion efficiency in pofnt-of-use cogeneration of electrfffty and heat, independence of efficiency on power plant size, and Amongst the various fuel cell schemes, solid oxide fuel cells ;e&fble pollution. are believed to have the widest potential range of applications and hi hest system efficiencies because of their all-solid-state design and hfgh (~lOOO°C s operating temperaturesITc which lead to internal reformfng of fuel gases and efficient utilization of waste heat. Successful prototype SOFCs based of 3~ttrfa-stabilized zfrconfa (YSZ) electrolytes are currently under further development. The economic viability of SOFC power generation hinges on processing costs, materials costs, and cell performance. While the operating characteristics of exfstfng SOFCs are good, it is not yet clear that costs will be competitive. In particular, the use of high processing temperatures fncreases fabrication costs while materials costs are relatfvely high since thick (1 a$Bfilms of rare-earth-containingcompounds, e.g. LaMnO , are used as the oxygen electrode. It is thus of interest to consider SOFC-design opsions that may reduce costs or exhibit enhanced characteristics. An alternate approach that would reduce materials costs is an all-thin-film SOFC desfgn. Low temperature thin-film deposition methods such as magnetron sputtering are available that provide good control over composition, properties, and morphology. Since magnetron sputter deposition is commonly used as a large-scale production technique, one might expect processing costs to be relatively low. Detailed studies of cell processing will be required before reliable cost predictions can be made. The reduced amount of material in a thfn film cell would also cut cell mass leading to higher power-to-mass ratios. However, the feasibility of a thin-film SOFC has not been established. Two technical advances are required to make thin film SOFCs a viable approach. First, methods for fabricating supported thin film structures must be developed. Second, a thin film oxygen electrode is needed. In SOFCs operating at Tc4=510000C, various stability crfterfa dfctate the use of low conductivity oxide cathodes l that must be-1 tmn thick in order to reduce ohmic losses to tolerable levels. In order to make a thin film SOFC, it m ,I&&
1
S. A.BARNFIT
2
is thus necessary to reduce Tc sufficiently to allow the use of a high conductivity metal cathode, namely Ag, that eliminates ohmic losses at thicknesses of ~1 pm. Since the volatility of Ag becomes a problem at T, > 750°C, technical feasibility depends on being able to achieve acceptable reaction overpotentials and electrolyte ohmic losses at this temperature. In thfs article, the technical feasibility of a supported thin film SOFC operating at < 750°C is discussed. A detailed, thin-film SOFC design is developed to allow estimation of the expected low T operating characteristics. The design is based on a thin (< 101.1m), YSZ-supported electrolyte that provides very high oxygen conductance. The YSZ is coated with thin oxide layers in order to reduce reaction overpotentials. Likely candidates are doped cerfa on the fuel electrode side and stabilized bismuth oxide on the oxygen electrode side. Estimates based on reaction rates for these materials indicate that the overall cell will operate efffcfently at temperatures T, as low as 700°C. In addition to allowing the use of thin Ag films for the oxygen electrode, the low T, also leads to higher thennodynamfc efficiencies.and a wider range of allowable reducing fuel mixtures. The predicted operating characteristics of the proposed thin-film SOFC design are at least as good as existing cells. The thin-film approach would thus seem to merit further research. 2.
THIN FILM SOFC DESISN CRITERIA
The main limitations on state-ofithe-art SOFC performance are oxygen electrode ohmic loss and activation overpotential. Ohmic losses are important despite electrode thicknesses of 1 mm because of the relatively poor conductivity of the oxide cathode. While a metallic electrode with much higher conductivity would reduce ohmic losses even for thicknesses in the micron range, economically-viablemetals are not sufficiently stable at T = 1ooooc. The proposed design thus depends on being able to reduce T sufficiently in t6e thin-film cell in order that a metallfc cathode, specifically Ag, can 6e used. In this section, the main factors limiting T, values - reaction overpotentials and ohmic loss in the electrolyte - are discussed. It is concluded that T, can be reduced sufffclently to allow the use of Ag cathodes. A. -
Electrode reaction kinetics
Slow reaction kinetics at the fuel cell electrodes can lead to large interfacial resistances and hence significant losses at high current densities. Figure 1 is a comparison of the low voltage interfacial resistance Rf in ai6 8s a function of temperature for several different electrolyte/electrode combinations. In the case of a YSZ Temperature, T, ("C) 1200
1000
600
600
500
Fig. 1. Interfacial resistances in air as a function of inverse temperature
for several different electrode-electrolyte Data were taken from Refs. 6-8.
combinations.
The the dependence on electrode material and morphology is very strong. electrolyte, shaded reafon shows the ranoe of R; values obtained for Pt electrodes with different morphologfes. While there is-a wide Lange of values observed, the ~10~ is independent of
Results for the moroholoav. fndfcatfna that the rate-limitina mechanism is the same. mixed co&ctfng oxide La 6Sro 5Co03 are also shown and are in the same range as observed for Pt. Much lower resfs ?*ances'are observed for Ag on YSZ and the dependence on morphology is considerably reduced. In all cases, R. increases rapidly with decreasing temperature and the resulting overpotential has precludea attempts to reduce T, in YSZ-based cells.
New solid oxidefuelcell design
3
Substarftial decreases in R. can be obtained by using a stabilized bismuth-oxide electrolyte. Results are shown i'nFig. 1 for Pt on 20% Er-stabilized bismuth oxide (ESB) and Ag on 20% Ba-stabilized bismuth oxide (BSB). For doped bismuth oxide, the oxygen reaction has been shown to take place directly on the electrolyte quiface, resulting in a For example, mixed very low Ri value and only a weak dependence on electrode material. ’ exhibited only slightly higher Ri values than Ag. conducting oxide electrodes on BSB Clearly, the use of doped bismuth oxide as the cathode rather than YSZ would reduce overpotentials. It should be noted thlt other catalytically-activeoxides such as U02 have been successfully used to decrease Ri. Catalytic properties for H20 reduction are be{der at ceria electrolytes than YSZ leading to lower overpotentials in electrolysis CellS. As in the case of bismuth oxide, the higher rate kinetics are apparently due to a reaction on the electrolyte surface. SOFC electrolyte applications for gadolinia-doped ceria (GOCj11a12compositionexhibiting high oxygen conductivity, have attracted considerable interest. ’ The applicability of ceria and bismuth oxide electrolytes is limited, however. Bismuth oxide is unstable in reducing envim ts and the use of ceria in fuel cells is On the other hz#, ceria-doped surface problematic due to low transference numbers. ’ layers on YSZ have been used successfully in reducing environments. ":hi: electronic conduction through the electrolyte. Doping with bismuthThoex:dSeZ1yr~~enat: oxygen electrode side of the electrolyte has also been successfully used to enhance reaction rates. These results suggest that a multilayer electrolyte, for example GDC/YSZ/BSB, would minimize activation overpotentials while still retaining chemical stability and 100% ionic conduction. Figure 2 shows the effective oxygen partial pressure as a function of position three-layer electrolyte with air on the cathode side and an 0 partial pressure P = 36 atm (a hydrogen to water partial pressure ratio of 500 at 700$), for example, on ? he fuel side. Conductivity data for the three electrolytes at 700°C was used in constructing this figure (see section 1I.B below). As long as the BSB layer was sufficiently thin ,GDC
BSB \
25
0
I
fuel electrode
Electrolyte Depth (urn)
I
air electrode
Fig. 2. Oxygen partial pressure as a function of distance in a thre -layer electrolyte at 700°C with air on one side and PO = 10-56 atm on the other. Electrolyte resistances were taken from Refs. 8, 16, and 17. relative to the YSZ, it would be exposed only to an oxidizing environment and hence remain stable. The GDC would remain chemically stable in the reducing fue$flectrode environment. but would exhibit a significant component of electronic conduction. The GDC layer would thus also have to be quite thin relative to the YSZ in order to prevent shorting of a significant portion of the cell voltage. Nonetheless, very thin GDC and BSB layers should be sufficient in order to enhance surface reaction rates. The use of very thin layers combined with the good thermal expansion match between i@ three oxides eliminates any thermal stability problem for the three-layer electrolyte. As discussed in section IV. the deposition of this three-layer electrolyte will not significantly complicate the cell
S. A.BAIw~
Cell Temperature. 900
850
800
T, (‘C)
750
700
650
'Orn
Fig. 3. Air and fuel electrode interfacial voltages at a current density of 1 A/cm2 as a function of inverse temperature. Data were taken from Refs. 8 and 9. SOFC
11001000so0
Operating Temperature, T, (OC) 800
1000/l (K-l) Fig. 4. Electrolyte thickness at which the ohmic voltage dsop across the electrolyte is 0.05 V for a current density of 1 A/cm , plotted vs. inverse temperature for three electrolytes. Resistivity data were taken from Refs. 8, 16, and 17.
New solid oxide fuel cell design
5
processing. Existing data indicate that overpotentials low enough for efficient SOFC operation can be achieved at T < 75O'C. For example, resistance losses in a SOFC with a Ni/Ybstabilized ZrO (ebSZ) anode and (In,Sn)oxide/U-dopedYbSZ cathode we% dominated by the Fie. 3 is a plot 0.4 mn thick Y& electrolyte conductivity, not Ri, at T down to 8OO'C. vs. T, of anode and cathode overpotentials for a currenE density9J of 1 A/cm . Data for the anode was taken for a Ni cermet on YbSZ in a 50% H -H20 mixture. (No data of this type is available for GDC.) Data for the Ag/BSB cathode are extrapolated from the resistance results shown in Fig. 1. Also shown are data where UO2 wab added to the YbSZ electrolyte Each of these combinations to enhance reaction rates at a In 0 -SnO air electrode. yields an acceptable overpotential aQ$, %ooc. B_
Electrolyte resistance
Electrolyte resistance is currently not a limitation on SOFC performance since the electrolyte layers are < 50 Am thick and T, = 1000°C. However, at T, ?r 700 to 750°C, electrolyte resistance becomes a limiting factor. Figure 4 is a $lot of the electrolyte thickness t at which the ohmic voltage dro&is 0.05V at J = 1 A/B versus 1000/T , based on conductivity data for YSZ (10% yttria), GDC (10% gadolinia), and BSB (20% baria). Since the GDC and BSB layers are quite thin in the proposed electrolyte and have conductivities similar to the YSZ, they will make negligible contributions to the overall resistance. The electrolyte must be < 10 Pm thick in order to keep ohmic losses < 0.05V at 7OO'C. Lower voltage drops can be achieved by using thicknesses < 10 #II and/or T, > 7OO'C. In addition, the electrolyte conductivity must be due mainly to the ionic contribution di, with only a minor conductivity~~. That is, the YSZ ionic transference number t be near unity in order that electrical conduction does not Fig. 5 shows the range of T, and PO values where ti > 0.99 for stabilized zirconia. Efficient fuel cell operation is possible Temperature 02) 16 6
-32
1 Ca-
stabilized ZrO,
3.0
5.0
7.0
9.0
11.0
13.0
1O"lT (K-l)
Fig. 5. A plot of the temperature and oxygen partial pressure range where the conductance of stabilized zirconia is primarily ionic, i.e. where the ionic transference number is > 0.99. Data were taken from Ref. 18. anywhere within this region. Lowering T, from 1000 to c 700°C has the advantage of expanding th~of~olI;!~~trode PD range that can be used without electronic conduction by a factor of 10 atm, as shown in Fig. 5. The low P8 value at the fuel electrode shown in Fig. 2 is thus acceptable for a cell operated at 700 C. C,
Oxygen Electrode Material
A SOFC based on a three-layer electrolyte < 10 Pm thick would be expected. based on the above discussion, to operate with acceptable overpotentials and electrolyte ohmic losses at T, < 75O'C. In this temperature range, Ag is a desireable alternative to conducting oxides for the air electrode since it 4ceines the good catalytic activity of the noble metals with high oxygen solubility. l Furthermore, Ag is an excellent conductor such that z 1 Pm thick Ag electrodes would be sufficient to eliminate ohmic losses, and hence materials cost would not be excessive. Ag has alwaylgbeen disqualified for use in high-T, cells because of its volatility. The Harwell group found that 2 Pm of Ag would be lost after 1 year in air at BOD'C. Extrapolations based on vapor-pressure data indicate that only 50 nm would be lost in 1 year at 700°C and 0.3~m at 750°C. Furthermore. in a thin-film approach it would be
S. A.BARNE~~
6
strafghforward to deposit a mixed-conducting oxide cap layer on the Ag electrc ,de in order to reduce Ag evaporation. This would oennft small 'increases in T_ above 7'50°C, and a consequent reduction in overpotentials. Finally, since 700°C is a siinfficant fraction of the Ag melting point (960°C), it will be necessary to co-deposit an oxide (i.e. BSB) with the Ag In order to minimize sfntering. 3.
THIW FILM SDFC DESIGW MD
PERFDRMWE
In this section, details of the design of a thin-film SOFC are consfdered and the performance of the resulting cell predicted. Fig. 6 shows schematically the SOFC design. Air
Fig. 6. A schematic drawing of the proposed thin-film fuel cell. Note that the drawing is not to scale.
In addition to the three-layer electrolyte and the Ag-cermet cathode, a Nl-cermet anode is proposed I-VSZ cermet is currently successfully used in SOFCs as the fuel electrod;.1'3*TI The chemical compatabilfty of Ag and Ni with the electrolytes is good. based upon their previous extensive use in solid electrolyte cells. The overall SOFC structure is supported on a Ni mesh frame which also serves for current collection. Techniques for fabricating this structure are discussed in section IV. A_
Cell dimensions
The maximum thickness t 1 of the electrolyte for T, = 700°C is 10~ m. The minimum thickness, on the other ban%, is given by the requirements that the electrolyte does not break down at operatfng voltages, that the film is stable against Interdiffusion, that it possess good mechanical stability, and that it fully covers imperfections on the substrate surface during deposition to form a pore-free fflm. Electrical breakdown can be discounted for fflms thicker than ~10 nm since breakdown fields are > lo6 V/cm and cell voltages are Sl v. Interdiffusion should not be a problem fo50tel > 500 nm and T, s700°C where oxide Finally, previous results have shown films have been found to be stable for long times. that 0.5 u 91is thick enough to obtain a pore-free electrolyte layer on a variety of substrates. Based on the above arguments, the overall electrolyte should be 0.5 to 10 urn thick. Mechanlcal stability and economic considerations will likely determine the optlmal thickness. Both the cell thickness and the free-standing film area (determtied by the Nf support frame mesh size) can be varied to achieve the required mechanical stability. There is considerable evidence suggesting that a mechanlcally stable SOFC can be based on thin ceramic films. which are sufficiently robust and resistant to fracwe and fatigue that a variety of coarnercfalelectromechanical devices are based on them. For example, freestandfng films such as cantilever beams up to 120 vrn long of 100 - 900 nm thick SfO2, en fabricated and reliably used as high frequency and 7059 glass ha$$,,#$ si3Nfg Nb2059 From an economics point of view, lower thicknesses oscf lators and pressure sensors. are desfreable to minimize processing time when techniques such as magnetron sputtering are used for film deposftlon (see section IV). Detailed studies of thin-film mechanical properties ~111 be required to determine an optimal electrolyte thickness. Electrode film thicknesses can be estimated using the criterion that < 10% of the to 1 cell volvge be dropped across the electrodes. The lateral voltage drop V 1 Is given by" V,l = JeL /2tf where J is the current density, P the electrode resistivfty, tf the
New solidoxide fuelcell
design
I
film thickness, and L the distance between current carrying wires. P for Ni and Ag were taken at 700°C, giving p i = 23.3 n cm and PA = 5.2 ncm. These values are multiplied by a factor of 5 to accoun'jcfor the porous ce&et nature of the electrode films. Sputter deposited and annealed cermet films with metal volume fr#c#dons > 50% have been shown $0 J is taken to be 1 A/cm . have p values within a factor of 5 of the pure metals. * Taking, for example, tf = 100 nm gives L = 1 mn for Ni and 2 mm for Ag. A Ni foil mesh would be used as the mechanical support for the fuel cell and also as a network for current collection. Since the free-standing electrolyte area will likely be limited to < 1 mm by mechanical stability criteria, a 100 nm Ni fuel electrode would essentially eliminate ohmic losses at that electrode. On the other hand, a thicker ( % 2 Pm) Ag film would be useful in order to achieve L = 1 cm. The Ni-foil mesh size would range from 0.1 to 1 mm , with the upper limit due to electrolyte mechanical stability and the lower limit due to ease of fabrication. A foil thickness of 2511m would provide sufficient rigidity over a 1 cm width while minimizing materials cost. The percentage open area of the mesh, a compromise between maximizing the cell area and obtaining sufficient rigidity and conductivity, should be s 70%.
2B
Predicted SOFC performance
The proposed cell has negligible ohmic losses at the electrolyte (for T, = 75O'C) and Ni electrode (for t&la> 200 nml and a voltaoe droo of 0.05 V at the Aa electrode for J = 1 A/&The major "sburcesof *losses in ce'iloperation are the overpotentials associated with kinetic rate limitations at the electrodes. The overvoltages at both electrodes at 750°C range from 0.07 to 0.25 V (Fig. 3). A conservative estpte of the overpotential is thus 0.3 V, giving s 0.35 V total voltage drop at 1 A/cm . Based on these losses, a predic$ion of cell performance can be made. Using the 2/O reaction potential of 1.1 V at 750°C, this gives a cell voltage of% 0.75 V at 1 A/cmv ._ ii oting the fact that the actual cell open area is 70X, the power density is ?. 0.5 W/cmL with a fuel efficiency;ff %50X. This compares favorably w$th the Westinghouse cell operftfng on H2 at 0.5 A/cm with a power density of 0.33 W/cm and a fuel efficiency of 46%. 4. A -I
TRIM FILH SOFC FABRICATIOW.INTERCORNECTS, ANDCELL STACKS
Fabrication
Thick-film electrolytes in state-of-the-art SOFC designs are deposited directly onto porous s bstrates using the electrochemical vapor deposition (EVD) technique to fill the pores.letl In this case,there is a limitation that the minimum film thickness be on the order of the substrate pore size, i.e. %5O#m, in order to fully cover all pores. To work with 0.5 to 10um thick electrolytes, it will be necessary to deposit on fully dense substrates that can subsequently be removed or rendered porous. While many schemes are possible, the following seems to be a simple, economically-feasible approach. To fabricate the proposed cell, the entire cell structure, including Nf-cermet. GDC, YSZ. BSB, and Ag-cermet. will be sequentially deposited directly onto a continuous Ni mesh/polymer sheet. This substrate could be formed by lamination or by coating the Nf mash with a liquid polymer. In the latter case, a liquid solution that fills the pores in the mesh due to surface tension would be dried to provide a solfdz5dense substrate. Polymer solutions such as commonly used in semiconductor lithography would be suitable since they can withstand the temperatures ( <400°C) that would arise during deposition. The polymer would be removed, a er depositing the cell, either by wet chemical etching or ashing in an oxygen plasma>% A variety of thin-film deposition techniques are available for fabricating the structure shown in Fig. 6. The present discussion will center on the use of magnetron sputter deposition for preparing the cells. Magnetron sputtering is chosen because it can be used to deposit all of the cell materials at low temperature in a single production-line deposition chamber, it gives high deposition rates and high materials utilization efficiency, and provides excellej&,2fjontrolover film thickness, structure, intrinsic stress, adhesion, and composition. The oxide films would be deposited by reactive magnetron sputtering of Zr-Y, Ce-Gd and Bi-Ba metallic targets in Ar-O2 mixtures. Deposition rates of 150 nm/min have been reported for reactive magnetron sputtering of Zr02 such t&t, for example, only ?r6 mins would be required to deposit a lr m thick electrolyte. Nf-Ce-Gd and Ag-Bi-Ba would be cosputtered in Ar-O2 mixtures to form the cennet electr #es. Rates for Ag and Ni deposition are expected to be even higher than for the oxides. The deposition chamber would be a standard in-line sputtering system with 5 sputtering targets and a mechanism for moving substrates continuously beneath the targets. Substrate preparation and loading into the deposition chamber should be carried out in a clean room in order to eliminate dust particles which can lead to pinholes in the films. Depending upon the growth conditions, sputtering can be used to deposit either dense or porous films. s een widely used for depositing porous metal electrodes Sputteri?o for electrochemical cells. * 9!n,zB On the other hand, many groups have shown in the last
S. A.BARNER
8
several years that low energy ion bombardment of the substrate and growing film during obtain high density oxide films such as Zr02 and CeO2 at B z-
Cell interconnectionand stacking
The present monolithic cel1Idfygn would be suitable for incorporation into stacks as proposed by the Argonne group. ’ The basic cell stacking configuration is shown schematically in Fig. 7. In this approach, the Ni/polymer substrate would be formed into lntercon"ecl z?zz hl0leell uxmmnmm
air fuel
fuel
Fig. 7. A schematic drawing of the stacked fuel cell and interconnect. the V-shaped configuratlon shown and mounted on a ceramic cell support structure. The films making up the SOFC would be deposited directly onto this substrate. After removal of the polymer, the interconnect sheet would be added and the cells stacked. The fnterconnect would be a Ni sheet coated on the side exposed to oxygen with a conducting oxfde with low oxygen d ffusfvity to prevent oxidation of the Ni. iaCrOS, used successfully In existing SOFCs '-' would be a suitable choice. A separate sputter deposftfon system would be used to deiosft the oxide layer on the Ni sheet to form the interconnect. 5.
DISCuSSIDll
The thin-film SOFC overpotentials estimated fn section 1II.B indicate that for operatlon near 75O'C. the cell power density and efficiency would be as good or better than existing cells. Furthermore, it may be possible to increase the useful cell Tc range slightly. and hence fmprove cell performance, by capping the Ag cathode wfth a mfxedconducting oxide in order to reduce evaporation (see section 1I.C). An all-thfn-film deslgn employing a Ag cathode is thus not only feasible from an electrochemical pofnt of view, but may provide Improved efficiency and power density over existing SOFCs. The thin-film approach offers the advantages of monolithic cell design, reduced materials costs, higher power-to-weight and power-to-volume ratios relative tf existing SOFCs. Estimates were made for cell stacks as proposed by the Argonne group. The main materials cost is for the Ni (taken to be $ZO/kg) structural supports and interconnect, with a fraction due to the Ag (taken to be $250/kg) films and connectfons. A very conservative estimate gives a materialsScost cjf< $2O/kW. A power-to-weight ratio of > 2 kW/kg and a power-to-volume rati? > 10 kW/m are estimated for this type of cell stack using a power density of 0.5 W/cm . The lower T, value of the present SOFC design offers a few potential advantages. First, less insulation would be required for cell stacks to be thermally self-sufflcfent. a 'ons such as synthesis reactions, lower temperatures will Second, in non fuel cell appl$$,$3 Third. the low T value and low thermal mass of the provide improved selectivity. thin-film SOFC suggests the possibility that a short-ctrcufted cell can heat itself rapidly from an "idling" temperature to T,. A simple calculation based on a cell current limited by the oxygen electrode resistance (as expected for T, < 700°C based on the data presented above) and assuming a 5 pm total cell thickness yields a 1 degree/set heating rate at 400°C. and 10 degree/set at 450'. In transportation applications Involving intenaittant use, a well-insulated cell stack could thus be maintained at 400°C by burning very small quantities of fuel and be able to reach operating temperature in -1 min. The above discussion suggests that a thin-fflm approach to SOFC design merits further investigation. Considerable research and development would be required to make thls a viable technology. First, the fabrication processes described above and the resulting films would have to be characterized in detail. Second, detailed fnvestlgatfons of the
New solid oxidefuelcell design
9
mechanical stability of the thin-film structure would be needed. One issue is that the structure have sufficient toughness to withstand any gas pressure gradients or impacts that This requirement will impose limits on the Nf mesh it would be subjected to in normal use. dimensions and on electrolyte film thickness. In addition, a partially-stabf#red Zr02 could be electrolyte, known to exhibit much better toughness than fully-stabilized ZrO e partiallyused to improve mechanical stability. The lower fonic conductivity of t8' stabilized ZrO2 would not significantly degrade cell characteristics since they are not limited by electrolyte resistance. Another important issue is the stability in thermal cycling applications. While the thermal expansion coefficients of the Nf components, ceramic cell supports (presumably made of stabilized ZrO ), and_.,e;ectrolytesare well yatched, Ag has a relatively large mismatch with the oxides 2i1.9x10 / C for Ag VS. 1.2~10 /OC for YSZ). While the extent of the problems posed by this mismatch are not known, the First, the fracture toughness of ceramic-based following points should be noted. multilayer structyges with different thermal expansion coefficients improves as the layers favoring the present thin-fflm design. Second, it may be possible to are made thinner, minimize thermal strains by minfmiring the temperature cycling range as described above. Thfrd, most of the thermal strains will taken up in the porous, low elastic modulus Agcermet layer. High temperature fatigue in the Ag would thus be the likely thermal cycling failure mechanism. Acknowlecf~nts - We thank the Donors of The Petroleum Research Fund, administered by the American Chemical Society, for support of this research. REFERENCES J.T. Brown, Energy 11, 209 (1986). A.O. Isenberg. Solid State Ionics 3/4. 431 (1981). J.T. Brown and S.E. Veyo, Mod. Power Syst. 6, 43 (1986). C.S. Tedmon, H.S. Spacil. and S.P. Mitoff, 7. Electrochem. Sot. 116, 1171 (1969). 0. Yamamoto, Y. Takeda. R. Kanno, and M. Noda, Solid State Ionfcn2, 241 (1987). J. Sasaki. J. Mfzuasaki. S. Yamauchf, and K. Fueki, Solid State IosiTcs 3/4, 531 (1981). 70 (1983). 7. M.J. Verkerk, M.W.J. Hamnfnk, and A.J. Burggraaf, J. Electrochem. Sot. m, 8. T. Suzuki, T. Yamazaki, K. Kaku, and M. Ikegami, Solid State Ionics I5,xl (1985). 9. R. Accorsi and E. Bermann, J. Electrochem. Sot. 127. 804 (1980). L.J. Olmer. J.C. Viguie. and E.J.L. Schouler, Sam State Ionfcs I, 23 (1982). ::: T. Kudo and H. Obayashi, J. Electrochem. Sot. 123, 415 (1976). H.L. Tuller and A.S. Nowfck, J. Electrochem. Sot. 122, 255 (1975). :23:A.J.A. Winnubst, A.H.A. Scharenborg. aud A.J. Burggraaf, Solid State Ionfcs 14, 319 (1984). Kflner, P.F. Dennis, A.E. McHale, M. Van Hemert, and A.J. 14. 8.C.H. Steele, J.A. Burggraaf, Solid State Ionfcs 18 19, 1038 (1986). 15. S. Majumdar, T. Claar, and 8. I+andermeyer, J. Am. Ceram. Sot. 69, 628 (1986). 8.C.H. Steele, Solid State Ionfcs I& 391 (1984). :;: T. Kudo and H. Obayashi, J. Electrochem. Sot. 122, 142 (1975). of Electronic Ceramics, p. 143, L.L. Hench and D.8. Dove eds., 18. J. Patterson, in b Marcel Dekker, New YorkT971) T.L. Markfn, Power Sources Q, i. 583, D.H. Collins ed., Oriel, New York (1973). Ch. 3, S.M. Sze ed., McGraw-Hill, New York (1983). ::: A.C. Adams, inITechnolosy, L.B. Welsh, and F.R. Szofran, J. Vat. Scf. Technol. I& 177 21. J.E. Greene, R.mlinger, (1977). K.E. Petersen, Proc. IEEE, 70, 420 (1982). f3: K.E. Petersen and C.R. GuarSireri,J. Appl. Phys. 50, 6761 (1979). 8. Abeles, P. Sheng, M.D. Coutts, and Y. Arie, Adv. Phys. 2. 407 (1975). Ht: S.K. Ghandhi, in VLSI Fabrication Princi les, Wiley, New York (1983). 26. See, for example,TE. Sundgren,-EX&- o ansson, A. Rockett, S.A. Barnett, and J.E. Greene, in Chemistry and Physics of Hard Coatings, p. 95, B. Sproul, J.E. Greene, and J.A. Thornton eds., AKNew York-m 27. J.A. Thornton, in Thin Film Processes, p. 76. J.L. Vossen and W. Kern,Academfc, New York (1978). F. Jones, J. Vat. Sci. Technol. A 5, 3088 (1988). z: T.M. Gur. I.D. Raistrick, and R.A. Huggfns, J. Electrochem. Sot. 127, 2620 (1980). 30. J.E. Greene, S.A. Bamett, J.-E. Sundgren, and A. Rockett, in _Ion Beam Assisted Film Growth, Ch. 5, T. Itoh ed., Elsevier, Tokyo (1988). Martin, J. Mater. Sci. 21, 1 (1986). and references contained therein. 31. mP.J. T.M. Gur and R.A. Huggins, SolidState Ionics 5, 563 (1981). R.D. Farr and C.6. :32:C.T. Sigal and C.G. Vayenas, Solid State Ionics 5, 567 (1981); Vayenas, J. Electrochem. Sot. 127. 1478 (1981). 34. T. Yoshida, T. Hoshfna, I. Mukxawa, and S. Sakurada, J. Electrochem. Sot. 136, 2604 (1989).
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