Membranes M Kusnezoff, Fraunhofer Institute for Ceramic Technologies and Systems, Dresden, Germany & 2009 Elsevier B.V. All rights reserved. Introduc...

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Membranes M Kusnezoff, Fraunhofer Institute for Ceramic Technologies and Systems, Dresden, Germany & 2009 Elsevier B.V. All rights reserved.

Introduction High-temperature fuel cells such as the solid oxide fuel cell (SOFC) is an important type of fuel cell that can be used as an alternative to the conventional heat engine owing to the high efficiency of energy conversion. This type of fuel cell, first invented by Grove (1839) and Scho¨nbein (1838), has undergone a major transformation due to technology improvements and is close to industrial application. The main feature of the SOFC is the utilization of a solid electrolyte, which is an ionic conductor of oxygen ions or protons, for the manufacturing of membrane–electrode assembly (MEA). Different types of solid oxides have been proposed as the electrolyte material for SOFC membranes such as zirconia stabilized with yttria or scandia, ceria stabilized with gadolinia or samaria, and lanthanum strontium magnesium gallate as oxygen ionic conductors, as well as barium/strontium cerates as protonic conductors. The operational temperature of SOFC membrane ranges between 550 and 1000 1C depending on the electrolyte type and thickness as well as electrodes used. Industrial membranes have an oxygen-ion-conducting electrolyte and work in the temperature range of 700– 900 1C to achieve the highest possible power density and a reasonable long-term stability. The main advantage of the SOFC membrane is that every type of fuel, (e.g., fuels containing hydrogen, carbon monoxide, methane, as well as ammoniacal fuels) can be utilized because of the high operating temperature, assuming that no carbon (soot) formation takes place under operating conditions. The acceptable impurity level of typical fuel contaminants, such as sulfur, chlorine, and silicon, depends strongly on the operating temperature, the anode material, and the fuel gas composition. The ongoing cell development is mainly focused on the reduction of ohmic and polarization losses of the membrane to achieve a high power density and efficiency



for the SOFC generator (called stack). Beyond the power density, the long-term mechanical and chemical stability and robustness of the system cycles are important issues for membrane development.

Membrane–Electrode Assembly: Principal Constitution and Electrochemical Reactions Basically, an SOFC corresponds to electrochemical transfer of oxygen from the air side to the fuel side. The driving force for this oxygen transfer is the difference in the oxygen partial pressure between two sides of the electrolyte membrane. On the air side (cathode), the oxygen is present in a molecular form; on the fuel side (anode), there is no molecular oxygen. The electrical energy, EF, which is the product of the cell potential E (volts) and the Faraday constant F, liberated while current passes through an external load is equal to the free energy DG of the corresponding reaction between the fuel (hydrogen, carbon monoxide, methane) and air (oxygen). In contrast to a battery, the electrode of the fuel cell does not contain any energy storage medium and acts only as a catalyst for partial electrochemical reactions on the air and fuel sides. The electrochemical cell consists of gastight thin electrolyte (5–200 mm) and two electrodes: an anode (fuel side) and a cathode (air side). Both electrodes are strong catalysts based on ionically and electronically conductive composites. The cathode/anode consists of at least two layers: thin electrochemically active cathode/anode and current collector layer on top of it (see Figure 1). The cathodic polarization losses in the electrochemically active layer are defined by oxygen transport through the porous electrode and the oxygen reduction kinetics at the electrochemical active sites, which are

(La, Sr)MnO3 (La, Sr)MnO3−SZ SZ SZ



Figure 1 Principal constitution of membrane–electrode assembly (MEA) for solid oxide fuel cell (SOFC).


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called ‘three-phase boundaries’ owing to the coexistence of three phases: electronic conductor and ionic conductor as solid phases and molecular oxygen as the gas phase. In the cathode, the reduction of molecular oxygen takes place: O2 þ 4e -2Oel 2

Usually, the cathodic polarization resistance increases on decreasing the oxygen partial pressure in the oxidant and the operation temperature. The anodic polarization resistance depends on the fuel gas composition. With preconverted natural gas as a fuel, the following reactions take place in the anode: 1a: Conversion : H2 þ Oel 2  2e -H2 O 1b: Conversion : CO þ Oel 2  2e -CO2 2: Reforming : CH4 þ H2 O-CO þ 3H2 3: Water shift reaction : CO þ H2 O-CO2 þ H2

In the presence of significant amounts of hydrogen (420 vol%) and water vapor (46 vol%) in the fuel, the oxidation of hydrogen to water (1a) is the reaction most likely to take place in the anode. In this case, instead of the electrochemical oxidation of carbon monoxide (1b), a water shift reaction (3) resulting in hydrogen and carbon dioxide production in the gas phase takes place. The total resistance of the MEA is the sum of the electrolyte resistance (ohmic loss arising from the ionic resistivity in the solid electrolyte) and the polarization resistances of the electrodes. To separate the impact of the different contributions to the electrochemical resistance, impedance spectroscopy must be used. In contrast to a battery, the total cell resistance (RCell) cannot be directly determined from the I–V characteristics. The reason for this is fuel utilization, which takes place in passing current through the cell and changes the open-circuit potential of the cell along the electrode. For this reason, the value of the cell resistance obtained from the I–V curve should be estimated using the mean opencircuit voltage (OCV) between the inlet and the outlet of the cell by using the following formula: RCell ¼

Inlet Outlet ðUˆ Nernst þ Uˆ Nernst Þ=2  UCell Itot

The specific cell resistance normalized to electrode area (A) also called area specific resistance (ASR, O cm2) ASR ¼

RCell A

is generally used as a measure of cell performance. The cell performance can be significantly improved by utilization of novel materials and technologies for electrode and electrolyte manufacturing.


Cell Materials Electrolyte The electrolyte for SOFC should be gastight and have high ionic and negligible electronic conductivity. The doped zirconia (ZrO2) is well known as a material with pure ionic conductivity and was originally used as an electrolyte material for SOFC. Zirconia powders contain, as a role, small amounts of hafnia (0.9–2.5 wt% depending on powder manufacturing process and raw materials used), which is often included into zirconia balance in the chemical composition. The dopants with di- or trivalent cations such as calcium oxide CaO), manganese oxide (MgO), yttrium oxide (Y2O3), scandium oxide (Sc2O3), or others are used to stabilize tetragonal or cubic phase in the wide temperature range and to avoid phase transformations during thermal cycling. The tetragonal phase can be stabilized by low-level doping (3 mol% yttrium oxide in zirconium dioxide (ZrO2); for cubic phase stabilization, a higher dopant concentration (48 mol% (yttrium oxide and zirconium dioxide) should be used. The dopants take the position of Zr4þ ions in the lattice. To maintain the electrical neutrality of the lattice, the negatively charged oxygen vacancies are built: ZrO2

Y2 O3 - Vo¨ þ 3Oo þ 2Y0Zr

The presence of oxygen vacancies makes the ionic conductivity in the solid electrolyte possible. The ionic conductivity that results from the oxygen transport in the anionic lattice depends on concentration of the vaccancies and their mobility. The maximum conductivity depends strongly on the kind of dopant and its concentration. The concentration of the vacancies is proportional to the dopant concentration, and the mobility of vacancies depends on the temperature, dopant concentration, and art of the dopant used. The mobility of vacancies decreases by the doping of zirconium dioxide over some critical dopant concentration level owing to interaction between the vacancies in the lattice. The ionic conductivity of cubic phase is nevertheless always much higher than that of the tetragonal phase, and the reduction of conductivity owing to decrease of mobility takes place first by overdoping of cubic phase. Using scandium oxide, higher conductivity values in comparison with other dopants (i.e., yttrium oxide) can be obtained. The ionic radius of Sc3þ (0.88 A˚) is much closer to the ionic radius of Zr4þ (0.84 A˚) than that of Y3þ (1.115 A˚), which is used as a standard dopant, causing only small distortions of zirconia lattice and resulting in high mobility of vacancies. Zirconia codoped with scandia and yttria leads to conductivity values in between those of yttria-doped zirconia (YSZ) and scandia-doped zirconia (ScSZ) electrolyte. The ionic conductivity depends on the raw materials used for electrolyte production (purity and particle size), impurities (such as silicon dioxide (SiO2), aluminium


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oxide (Al2O3), and ferric oxide (Fe2O3), and others), and manufacturing process (sintering temperature and atmosphere), which all define the microstructure of the sintered electrolyte (porosity, grain size, number and composition of grain boundaries, etc.). Because of high sintering temperature of pure stabilized zirconia, the nanopowders or/and sintering aids are used to sinter it close to their theoretical density (porosity o2%) or at least to get almost gastight microstructure (porosity o6%) at temperatures p1400 1C. Various additives can be used to reduce the sintering temperature and/or to increase the mechanical strength of the electrolyte. The effect of dopant concentration and different impurities on ionic conductivity of the electrolyte can be better explained with zirconia-based electrolytes. Influence of Dopant Concentration The maximum ionic conductivity in zirconia-based systems is observed when the concentration of acceptortype dopants is close to the value necessary to stabilize the cubic fluorite structure. The highest conductivity levels that can be obtained by doping with different acceptor dopants are well known (Figure 2). Further increase of dopant concentration decreases the ionic conductivity owing to an increase in the association of the oxygen vacancies and dopant cations into complex defects of lower mobility compared with simple oxygen vacancies. This tendency increases with the increase in the difference between the host and the dopant cation radii (see Figure 2). Similar phenomena take place in numerous fluorite and perovskite structures. The maximum ionic transport for zirconium dioxide is observed by doping with Sc3þ. The attempts to increase the



Migration enthalpy ΔH m (kJ mol−1)

σ1273 K (S cm−1)









90 Yb3+


Er 3+


Dy3+ 3+

Y 70



Zr 4+ ΔH m


60 0.85




Association enthalpy ΔH A (kJ mol−1)



Dopant cation radius (Å)

Figure 2 Maximum conductivity in ZrO2–Ln2O3 systems at 1000 1C and the oxygen-ion migration and association enthalpies vs radius of Ln3þ cations. Reproduced from Arachi Y, Sakai H, Yamamoto O, Takeda Y, and Imanishai N (1999) Electrical conductivity of the ZrO2–Ln2O3(Ln ¼ lanthanides) system. Solid State Ionics 121(1–4): 133–139.

conductivity of cubic scandia-stabilized zirconia by codoping with ytterbia or other earth metal cations have not led to worthwhile improvement. Compared to other solid electrolytes, stabilized zirconia exhibits the lowest electronic contribution to total conductivity in the relevant oxygen partial pressure range. Influence of Impurities Most of the zirconia powders used for electrolyte manufacturing contain small amounts of impurities. The origin of impurities lies in the processing and purification of raw materials used for zirconium dioxide and yttrium dioxide recovery. The dominant impurities are silicon dioxide, magnetite (Fe3O4), and aluminum oxide. These oxides have very low solubility in both zirconium dioxide and yttrium dioxide. Sometimes, alumina, silica, or iron oxides are added to the electrolyte powder as sintering additives. During sintering, the oxide with low solubility in zirconium dioxide stays at the grain boundaries and on the surface, influencing the mechanical, electrical, and chemical properties of the material.


During sintering of Y2O3–ZrO2 material, the silicates or glassy phases with melting point in the temperature range of 1000–1500 1C can be formed at the grain boundaries. The physical and chemical properties of the glassy phase as well as wetting of zirconium dioxide grains depend strongly on the composition and sintering conditions (peak temperature, sintering atmosphere, and annealing conditions after sintering). Especially, the position of the glassy phase in the sample (surface or bulk of the sample) is very sensitive to the peak temperature and annealing history. There are three general possibilities for glassy phase distribution in the electrolyte: 1. The glassy phase wets the grain boundaries during the sintering process and is placed between the grains of zirconium dioxide. The glass is homogeneously distributed in the electrolyte. 2. The glassy phase dewets the grain boundaries and is situated in the knuckles of the grains. In this case, the glassy phase is inhomogeneously distributed in the electrolyte. 3. The glassy phase moves from the interior of the material to the surface. The segregation of the impurities at the grain boundaries and distribution of the glassy phase are a complex function of the amount, art, and concentration of different impurities. The silica-containing phases such as 3Al2O3  2SiO2, Na2O  SiO2  Y2O3, SiO2  Y2O3, and Mn2O3  SiO2  Y2O3 are reported in the literature.

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The segregation of silicon dioxide at the grain boundaries influences the ionic conductivity, and the presence of silicon dioxide on the electrolyte surface can cause deterioration of the electrode polarization resistance. The ionic conductivity of the electrolyte depends on grain and grain boundary conductivity of the material. As a role, the ionic conductivity at 300–450 1C is very sensitive to the presence of second phase and increases with the increase in the silica content in the YSZ powder. Also small amounts of silica (150 ppm) have a pronounced effect on the grain boundary conductivity of pure YSZ. Sometimes, the influence of silica addition on the bulk conductivity of YSZ grains has been observed. This effect is due to the migration of yttria outside of grains toward the grain boundaries with formation of Y2O3–SiO2 phases. The reduced yttria content of the grains causes, in this case, a change of grain conductivity.


Alumina is one of the most important additives to zirconia for enhancing the mechanical properties of the electrolyte material. The addition of aluminum oxide to 8YSZ and 10YSZ has been used in precommercial products to enhance the mechanical strength of electrolyte-supported MEAs. The addition of small amounts of aluminum oxide (0.25–1 wt%) has small influence on the mechanical properties but sometimes enhances the conductivity of the electrolyte. The reason for this is the interaction with silicon dioxide impurities in original powders used for electrolyte manufacturing and reduction of silicon dioxide concentration in the grain boundaries. The interaction results in the formation of Al2O3–SiO2 phase, which dewets the grain boundaries and moves to the knuckles, or in the wetting of aluminum oxide particles with silicon dioxide and purification of grain boundaries between YSZ particles. Further addition of alumina (1–3 wt%) has, as a role, no influence on ionic conductivity. With increasing amounts of aluminum oxide (5–30 wt%), the ionic conductivity decreases and the mechanical strength increases. At the same time, the grain size of YSZ decreases and the grain boundary length in the sintered material increases. The decrease of ionic conductivity is due to the presence of isolating particles in the microstructure and the rise in grain boundary resistance owing to increased grain boundary length. The addition of aluminum oxide has negative effects on the polarization resistance of standard electrodes (YSZ/NiO and YSZ/(La,Sr)MnO3 (LSM)) owing to the reaction of alumina with nickel monoxide (NiO) and LSM. The surface of the aluminum oxide-containing electrolyte should be coated with a gastight barrier layer to avoid interaction with the electrode material.



The solubility limit of ferric oxide in zirconium dioxide at 1200 1C is B2.1%. The addition of small amounts of iron oxide (up to 1 mol%) does not influence strongly the conductivity of pure YSZ. The sintering temperature of the electrolyte, however, can be reduced. The addition of X2 mol% of ferric oxide leads to the reduction of ionic conductivity. Interaction with Electrode Materials (Mn2O3, La2O3, SrO, CeO2, NiO) Mn2O3, La2O3, SrO

Because of low solubility of lanthanum and strontium in YSZ (o2 at% lanthanum and 0 at% strontium at temperatures below 1500 1C), the diffusion of lathanum and strontium from LSM into YSZ and the corresponding influence on the interface stability are quite limited. Instead of the formation of solid solution with these oxides, the formation of poor ionic-conducting and electronically isolating secondary phases such as La2Zr2O7 and SrZrO3 takes place. It is however a different situation for the manganese diffusion, owing to relatively high solubility of manganese in YSZ. The Mn3þ ions have small ionic radii (0.58 A˚) compared with that of Zr4þ (0.84 A˚) and have good solubility in the zirconium dioxide lattice. The solubility limit of manganese in doped zirconia depends on temperature, oxygen partial pressure, and the composition of stabilized zirconia (art of dopant and its concentration). For 8 mol% yttrium oxide-doped zirconia, the solubility of 5 mol% manganese at 1000 1C, 8 mol% at 1400 1C, and 12 mol% at 1200 1C in air is measured. The decrease of oxygen partial pressure increases the solubility of manganese in the electrolyte. The interdiffusion of manganese from stoichiometric cathode material (LSM) causes partial decomposition of perovskite and formation of free La2O3 and afterward La2Zr2O7 by reaction with zirconium oxide during sintering of electrodes, which deteriorate the cathode polarization resistance. To avoid this effect, the lanthanum understoichiometric perovskites are used. In moderate concentrations (up to solubility limit), manganese(III) oxide acts as a sintering additive and induces grain growth and simultaneously reduces the porosity of zirconia. However, the conductivity of stabilized cubic zirconia decreases and the activation energy as well as the total conductivity increase by the addition of small amounts of manganese(III) oxide. Nevertheless, the ionic transport number for 8YSZ with 4% Mn is 40.99 for pO2 ¼ 105  1015 Pa. CeO2

Doped ceria is also often used in the electrodes of SOFC on the cathode and anode sides. Especially, the barrier


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layers of gadolinia-doped ceria are used for electrode fabrication. Ceria has unlimited solubility in YSZ, and solid solutions in this system are widely investigated. The transport properties of mixed oxides such as [(ZrO2) 0.85(YO1.5)0.15]1x[(CeO2)0.8(YO1.5)0.2]x , [(ZrO2)0.85(YO1.5)0.15]1x- [(CeO2)0.8 (GdO1.5)0.2]x , ([ZrO2]1x[CeO2]x)0.85(YO1.5)0.15, [(ZrO2)0.85 (YO1.5)0.15]1x  [CeO2]x, and Zr0.75 Ce0.08Y0.17O1.92 are known. The conductivity of 8YSZ, as a role, decreases by the addition of ceria (see Table 1). The ionic conductivity of the mixed oxide [(ZrO2)0.85(YO1.5)0.15]1x–[(CeO2)1y(ReO1.5)y]x depends on YO1.5 and ReO1.5 content. At the same zirconia content, the electrolyte with higher rhenium content has higher conductivity. In the cerium(IV) oxide concentration range x ¼ 0.4  0.6, a maximum reduction of ionic conductivity is observed. The interlayers with the composition of Zr0.38Ce0.37Y0.072Gd0.185O1.88 at the interface [(ZrO2)0.85(YO1.5)0.15]/[(CeO2)0.8(GdO1.5)0.2] are observed after sintering at 1400 1C/4 h. The conductivity of this phase is rather low and therefore the formation of mixed oxide leads to poor performance of the cell, especially at low temperatures. The interdiffusion of cerium, yttrium, and rhenium depends strongly on the type of ReO1.5 used for ceria stabilization and can be reduced by lowering the sintering temperature of the ceria interlayer. Ceria has the highest diffusivity among other species in the solid solution. The electronic conductivity in stabilized zirconia doped with ceria increases strongly under reducing conditions (pO2 ¼ 106  1014 atm), owing to cerium reduction from Ce4þ to Ce3þ. The ionic transport number even at low ceria concentrations in YSZ (x ¼ 0.1) can be o0.99. NiO

The solubility of nickel monoxide in YSZ is very low and depends strongly on the temperature (1.2 mol% at 1200 1C, 1.4 mol% at 1350 1C, and 1.5 mol% at 1500 1C). The ionic conductivity of 8YSZ decreases slightly by the addition of small amounts of nickel monoxide (1 and 2 mol%). Doping with nickel monoxide stabilizes the cubic phase in 8YSZ and reduces the degradation of ionic conductivity with time. Long-Term Stability of Ionic Conductivity The main factors influencing the deterioration of ionic conductivity in the YSZ are the operation temperature, phase stability (cubic/tetragonal/monoclinic), and impurity interdiffusion. The strongest degradation is observed for electrolytes with time-dependent phase content such as 8YSZ. The degradation of ionic conductivity of 8YSZ is a strong function of deviation of yttria content in the 8YSZ powders (Table 2).

The highest ionic conductivity for 8YSZ can be achieved at yttria concentration of about 7.6–8 mol%. However, the degradation of ionic conductivity for this composition is much higher when compared with yttria contents more than 8.5 mol%. The reason for degradation at lower yttria content in zirconia is the cubic/tetragonal phase transformation at the operation temperature. The codoping of electrolyte by interdiffusion of manganese, iron, and nickel can change the phase content in the electrolyte. For example, the codoping of 8YSZ stabilizes the cubic structure. Other Aspects Even at relatively low operating temperatures (600– 750 1C), the zirconia-based electrolytes can be used as electrolyte material in the form of thin gastight layers. The thickness of the electrolyte layer sintered or plasma sprayed on a porous support varies between 5 and 30 mm. The mechanical strength of the support is of great importance for application in the SOFC. The mechanical strength of ceramics is generally defined by defect size, form and distribution, and by the resistance against crack propagation described by fracture toughness (KIC). The mechanical strength depends strongly on manufacturing process and increases by reducing the grain and defect size in the microstructure. The highest strength can be obtained for tetragonal zirconium dioxide stabilized near the monoclinic/tetragonal morphotropic phase boundary (i.e., 3YSZ, 3ScSZ) and is achieved by electrolyte support made of these materials. The reason for this behavior is that the phase transformation from tetragonal to monoclinic phase reduces the mechanical stress at the crack tip and stops crack propagation resulting in high KIC values. Unfortunately, the mechanical stability of the cubic phase with highest ionic conductivity is much lower than that of the tertragonal phase (see Table 3) and therefore the substrates made of these materials have lower toughness. Alternative Electrolytes Ceria-based

Alternative electrolytes for SOFC are doped ceria and lanthanum gallate because of their high ionic conductivities (see Figure 3). The general properties of doped ceria (CeO2) are similar to those of zirconia. The advantage of ceria stabilized with trivalent cations such as yttrium, samarium, or gadolinium is the higher ionic conductivity in comparison with that of stabilized zirconia. Among ceria-based solid solutions with the composition Ce1xRexO2x/2, the highest level of ionic conductivity is achieved for gadolinium and samarium (x ¼ 0.10.2). However, the high ntype electronic conductivity in the reducing atmosphere caused by the reduction of Ce4þ to Ce3þ in the lattice makes it impossible to use this material as electrolyte at

Ts (1C) 1700 1C/3 h in air 1600 1C/10 h in air 1500 1C/8 h in air

1600 1C/10 h in air

1700 1C/3 h in air

1600 1C/10 h in air

1500 1C/8 h in air

1600 1C/10 h in air

1700 1C/3 h in air

1600 1C/10 h in air

Zr0.85Y0.15O1.96 Zr0.85Y0.15O1.96 x ¼ 0 Zr0.74Ce0.073Y0.15Gd0.017O1.91 x ¼ 0.09

Zr0.765Ce0.085Y0.15O1.96 x ¼ 0.1

Zr0.68Ce0.16Y0.16O1.948 x ¼ 0.2

Zr0.68Ce0.17Y0.15O1.96 x ¼ 0.2

Zr0.64Ce0.18Y0.13Gd0.044O1.91 x ¼ 0.23

Zr0.595Ce0.255Y0.15O1.96 x ¼ 0.3

Zr0.51Ce0.32Y0.17O1.936 x ¼ 0.4

Zr0.51Ce0.34Y0.15O1.96 x ¼ 0.4

Ionic conductivity of mixed oxide in the system CGO–YSZ


Table 1

TZ-8Y (Tosoh) TZ-8Y (Tosoh) TZ-8Y (Tosoh) Ce(NO3)3 Y(NO3)3 TZ-8Y (Tosoh) CeO2 (Kojundo) Y2O3 (AC&T Co.) TZ-8Y (Tosoh) Ce(NO3)3 Y(NO3)3 TZ-8Y (Tosoh) CeO2 (Kojundo) Y2O3 (AC&T Co.) TZ-8Y (Tosoh) Ce(NO3)3 Y(NO3)3 TZ-8Y (Tosoh) CeO2 (Kojundo) Y2O3 (AC&T Co.) TZ-8Y (Tosoh) Ce(NO3)3 Y(NO3)3 TZ-8Y (Tosoh) CeO2 (Kojundo) Y2O3 (AC&T Co.)

Raw materials used

99.9 99.9 99.9

99.9 n/s n/s

99.9 99.9 99.99

99.9 n/s n/s

99.9 99.9 99.99

99.9 n/s n/s

99.9 99.9 99.99

99.9 99.9 99.9 n/s n/s

Purity (%)

















0.2 0.141

s1000 1Ca (S cm1)

>95 96–97

Density (%)

(Continued )









0.94 0.89

Eab (eV)

Fuel Cells – Solid Oxide Fuel Cells | Membranes 39

1700 1C/3 h in air

1600 1C/10 h in air

1700 1C/3 h in air

1500 1C/8 h in air

1700 1C/3 h in air

1500 1C/8 h in air

1500 1C/8 h in air

1700 1C/3 h in air

Zr0.425Ce0.4Y0.175O1.93 x ¼ 0.5

Zr0.425Ce0.425Y0.15O1.96 x ¼ 0.5

Zr0.34Ce0.48Y0.18O1.924 x ¼ 0.6

Zr0.23Ce0.58Y0.038Gd0.146O1.91 x ¼ 0.73

Zr0.17Ce0.64Y0.19O1.912 x ¼ 0.8

Zr0.1Ce0.7Y0.017Gd0.176O1.9 x ¼ 0.88

Zr0.05Ce0.75Y0.008Gd0.188O1.9 x ¼ 0.94



Ionic conductivity in air. Activation energy for ionic conductivity in air at T ¼ 700–1000 1C.

1500 1C/8 h in air

Zr0.45Ce0.38Y0.073Gd0.096O1.92 x ¼ 0.48


Ts (1C)



Table 1

TZ-8Y (Tosoh) Ce(NO3)3 Y(NO3)3 TZ-8Y (Tosoh) Ce(NO3)3 Y(NO3)3 TZ-8Y (Tosoh) CeO2 (Kojundo) Y2O3 (AC&T Co.) TZ-8Y (Tosoh) Ce(NO3)3 Y(NO3)3 TZ-8Y (Tosoh) Ce(NO3)3 Y(NO3)3 TZ-8Y (Tosoh) Ce(NO3)3 Y(NO3)3 TZ-8Y (Tosoh) Ce(NO3)3 Y(NO3)3 TZ-8Y (Tosoh) Ce(NO3)3 Y(NO3)3 TZ-8Y (Tosoh) Ce(NO3)3 Y(NO3)3

Raw materials used

99.9 n/s n/s

99.9 n/s n/s

99.9 n/s n/s

99.9 n/s n/s

99.9 n/s n/s

99.9 n/s n/s

99.9 99.9 99.99

99.9 n/s n/s

99.9 n/s n/s

Purity (%)










Density (%)










s1000 1Ca (S cm1)










Eab (eV)

40 Fuel Cells – Solid Oxide Fuel Cells | Membranes

Fuel Cells – Solid Oxide Fuel Cells | Membranes

Table 2


Ionic conductivity at 1000 1C before (s1000 1C) and after annealing (s*1000 1C) in air


Y2O3 (mol%)

Density % dth

Ts (1C)


s1000 1C (S cm1)

s*1000 1C (S cm1)


Tosoh TZ-8Y YSZ7.7 YSZ8.5 YSZ9.5 Nissan Chemical Daiichi Toray Scimarec Co. MEL8Y

7.8 7.7 8.5 9.7 7.63 7.96 8.15 8.54 8.85

99 97–99 97–99 97–99 99.2 99.5 97.6 97.7 92.8

1500 1C/2 h 1500 1C/2 h 1500 1C/2 h 1500 1C/2 h 1500 1C/4 h 1500 1C/4 h 1600 1C/3 h 1500 1C/4 h 1500 1C/4 h

ISa IS IS IS dcb dc IS dc dc

0.156 0.16 0.149 0.107 0.178 0.177 0.16 0.165 0.135

0.1375 0.137 0.145 0.105 0.142 0.156 – 0.151 0.131

Disc Disc Disc Disc Bar Bar Disc Bar Bar

a b

Impedance spectroscopy data. Direct current (dc) four-point measurement data.

Table 3 Mechanical properties at room temperature and ionic conductivity of yttria-doped zirconia (YSZ) Property 1

Ionic conductivity at 850 1C (S cm ) Bending strength (MPa) Fracture toughness (MPa m1/2)




0.03 1200 3.3

0.05 480 6.3

0.08 232 1.6

T 4700 1C. The reduction of Ce4þ also causes expansion of the lattice owing to additional oxygen loss, which can lead to mechanical failure. The mechanical strength of ceria-based self-supported electrolytes is very low, and only metal-supported cells (MSCs) or cathode-supported cells/anode-supported cells (ASCs) with thin layers from doped ceria are mechanically stable. The thermal expansion coefficient of ceria-based electrolyte is higher than that of stabilized zirconia and much closer to the thermal expansion coefficient (TEC) of ferritic alloys used for planar stack manufacturing, which makes ceria attractive for metal-supported MEAs. The problems of electronic conductivity of ceria can be partially solved by a combination with other solid electrolytes such as stabilized zirconia. However, the performance of the electrolyte can deteriorate owing to the formation of reaction products (ceria-based electrolytes have higher sintering temperature when compared with zirconia-based electrolytes) or microcracks because of the differences in thermal expansion of the electrolytes. The sintering temperature of ceria-based electrolytes can be efficiently reduced by the addition of small amounts of transition metal oxides. Doped ceria is relatively chemically inert toward the standard electrode materials. The best performance can be achieved using Ni/ceria anodes and (La,Sr)(Co,Fe)O3 (LSCF) cathodes even at temperatures below 600 1C. LaGaO3-based

The perovskite-type solid solutions based on lanthanum gallates (La1xSrxGa1yMg yO3d or LSMG) have higher ionic conductivity than ceria- and zirconia-based electrolytes and also low activation energy. Compared with

ceria-based electrolytes, the electrolytic domain of LSMG is extended substantially to lower oxygen partial pressures. High oxygen ionic conductivity of lanthanum gallate (LaGaO3) is achieved by substituting lanthanum with strontium, calcium, and barium and incorporating divalent metal cation Mg2þ into the gallium sublattice in order to increase the oxygen vacancy concentration. Strontium cation leads to minimal lattice distortion among the alkaline earth elements giving the maximum oxygen ion mobility and as a consequence a higher ionic conductivity. For the LSGM solid solution, the maximum ionic conductivity is achieved at x ¼ 0.10.2 and y ¼ 0.150.2. Introduction of small amounts of cations such as cobalt and iron into gallium sites can increase the ionic transport in LSGM, and produces only a small increase in electronic conductivity. The main problem with this class of materials is the phase instability at high temperatures (at T 4800 1C, gallium is released from perovskite lattice, especially under reducing conditions). The formation of stable secondary phases and high reactivity with perovskite electrodes under oxidizing conditions as well as with metallic nickel under reducing conditions are other disadvantages of this material. The mechanical properties of LSMG are comparable with those of cubic zirconia (8YSZ); therefore, the electrolyte-supported cells (ESCs) can also be manufactured. The electrodes for this material should be carefully selected to avoid interdiffusion and reactions at the interface. Especially, iron and copper cause electronic conductivity in LSMG, reducing the ionic transport number of the electrolyte. The high cost of gallium and the toxicity of LSMG electrolytes are the main barriers to the commercial use of LSMG electrolytes. Cathode The most common cathode materials used for SOFC are LSM and (La,Ca)MnO3 (LCM) perovskites. Because of their good catalytic activity for oxygen reduction and their chemical and thermomechanical compatibility with doped zirconium dioxide electrolyte, these materials can


Fuel Cells – Solid Oxide Fuel Cells | Membranes Temperature (°C) 800 700




300 Bipolar plate material

Stainless steel Cr-Fe (Y2O3), Inconel−Al2O3 La (Ca) CrO3

log  (S cm−1)


1500 μm

R0 = L/ = 0.15 Ω cm2

Bi2 V


150 μm

0.9 Cu 0.1 O 5.35




0.9 S


2 )0

.9 (Y 2O 3 )0


15 μm



Self-supported electrolytes

0.9 G

.1 G


Supported electrolytes


.8 M

.1 O 1.9

1.5 μm



.2 O 2.8




0.15 μm








1000/T (K−1)

Figure 3 Conductivity of electrolyte materials and possibility of their use at different operation temperatures in the solid oxide fuel cell (SOFC) area specific resistance (ASR of electrolyte r0.15 O cm2).

O2(g) h+






h+ h+

n tio ac Re t tpb a


(h )

O2− VO



sio iffu lk d +


Micropores h+ O2(g)

Gas diffusion





h+ h+

sio n O x (a d

Gr ain b dif oun fus da ion ry


(h+) (Ln,A)MTO3+δ YSZ


VO VO (Surface) conductivity of electrolyte

VO ≈1 μm

Figure 4 Reaction paths in the cathode. YSZ, yttria-doped zirconia.

be directly applied as SOFC cathodes in the temperature range of 800–1000 1C. The polarization resistance reported for LSM cathodes sintered at 1300 1C is in the range of 1.5–19.8 O cm2 at 950 1C. The electrochemical reactions take place at the threephase boundary electronic conductor/ionic conductor/air (i.e., LSM/YSZ/air). In the porous electrode, the oxygen transport in the gas phase takes place in the pores of the cathode. The oxygen is adsorbed on the surface of the electronic conductor (LSM), and the adsorbed species diffuse to the interface electronic/ionic (LSM/YSZ)

conductor. At this interface, the adsorbed oxygen species go into the electrolyte lattice as oxygen ions (Figure 4). The LSM perovskites have very low ionic conductivity in comparison with doped zirconia, and the cathode containing a mixture of LSM and YSZ allows us to increase the length of the three-phase boundary and the ionic conductivity of the electrode. In agreement with this, it is found that the most efficient way to lower the polarization resistance of the cathode is to use a mixture of perovskite and electrolyte powders as an electrochemical active interlayer referred to as composite cathode.

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Furthermore, it is shown that the cathodes with increasing perovskite content from the electrolyte/cathode interface to the cathode surface (graded cathode) have the lowest polarization resistance. However, the highly graded structures produced by screen printing are complex and, therefore, costly to fabricate. The polarization resistance of the composite cathodes depends on the powder processing route, the structure of the layers (thickness, porosity), the coating technology (screen printing, slurry spraying), and the sintering procedure (temperature, cofiring). It is suggested that the interaction between electrolyte and perovskite resulting in the formation of La2Zr2O7 or SrZrO3 second phases is very sensitive to the processing procedures and crucial for electrocatalytic properties of the electrode. The formation of second phases can be avoided for La1xSrxMnO3d with a strontium content of 0.15–0.35 by use of La understoichiometry. For cobalt- and iron-containing perovskites, the intermediate cerium(IV) dioxide layer is generally used to avoid the reaction between electrode and electrolyte material. The sintering temperature of the cathode is the most important parameter for their electrochemical properties. The upper limit of the sintering temperature is limited to 1350 1C because of La2Zr2O7 formation due to the chemical reaction between ZrO2 and La0.75Sr0.2MnO3 at higher temperatures. The lower limit is defined by the adhesion of the cathode layers to the electrolyte. Anode In the anode, a large number of reactions take place, which are influenced by the anode composition and microstructure. The most important reactions are hydrogen conversion to water, water shift reaction with carbon monoxide, and internal reforming of hydrocarbons. The anode composition should simultaneously allow supply of fuel gas and evacuation of water and carbon dioxide, removal of electrons resulting from anodic electrochemical reaction, and supply of oxygen ions to the reaction sites (see Figure 5).


The most commonly used anode material is nickel. Because the thermal expansion coefficient of nickel is much higher (18  106 K1) than that of the electrolyte, the nickel/ceramic mixtures called cermets are used to assure thermomechanical compatibility and good adhesion of the anode to the electrolyte. In the cermet, the ionic conductivity is realized by ceramic component (stabilized zirconia or ceria) and the electronic conductivity by nickel. The addition of ceramic phase to the nickel also increases the number of electrochemically active sites (three-phase boundary) and the agglomeration of metallic nickel particles in the anode. The microstructure of the anode is crucial for obtaining high power density and long-term stability. The nickel and ceramic particles should build a network of electronic and ionic paths that allows a continuous change from ionic to electronic type of carriers from the electrode/electrolyte interface to the electrode/current collector interface, respectively. At the same time, the porosity of the electrode should be high enough for fuel supply and removal of products of electrochemical reaction (water and carbon dioxide). By building such three-dimensional network, the three-phase boundary length is moved from the electrolyte/electrode interface into the depth of the anode. The standard material for anode is Ni/YSZ cermet. Other metallic phases or nickel-based alloys ((Ni/X)/YSZ mixtures with X ¼ Fe, Cu, Co, or Cr) are added to increase the electronic conductivity and electrochemical activity of the anode. Nevertheless, pure nickel is mostly used as a metallic component of the cermet. To increase the ionic conductivity, the Gd/Sm-doped ceria or scandium-doped zirconia is used as the ceramic phase. The addition of ceria in this case is very effective due to high ionic and electronic conductivity under reducing conditions. Doped ceria acts as an active site for hydrogen conversion to water and helps to sustain the electronic conductivity in the anode in the vicinity of nickel. Care should be taken in combining ceria in the anode and stabilized zirconia as electrolyte. At high sintering temperatures, the interdiffusion of ceria into zirconia takes place, which leads to

Ni YSZ TBP Oxygen ions Electrons Reaction products (H2O, CO2)

Electrolyte Figure 5 Reaction paths in the cermet anode. YSZ, yttria-doped zirconia.

Fuel (H2, CO)


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the formation of intermediate layers of poor ionic conductivity and increase of polarization resistance. The main mechanism for anode degradation is the agglomeration of nickel particles during operation. As a consequence, the catalytically active surface and the electronic conductivity in the anode are reduced. To prevent nickel agglomeration, the wetting of nickel particles by doped zirconia or ceria is used. The activities in the development of cermet anodes are concentrated on the optimization of redox stability (stability against repetitive reduction/oxidation cycles), sulfur tolerance, and steam reforming activity. It is generally possible to produce redox-stable anodes for the ESCs. For ASCs, the problems connected with redox stability of the cermet-based substrates have not been successfully solved. The electrolyte cracking owing to substrate expansion during the redox cycle is the main reason for the destruction of the ASC. It was found that the irreversible expansion of the substrate during redox cycling can be reduced by using coarse YSZ and NiO powders as well as increasing the sintering temperature. Nevertheless, the reversible expansion of the substrate (420%) is always observed after full oxidation. The anode containing ScSZ or gadolinia-doped ceria as the ceramic phase is more sulfur tolerant than the one with YSZ.

Lifetime-Determining Processes The long-term operation at high current densities (4300 mA cm2) is the most critical point for the application of SOFCs as well as for the stack and system design. The inhomogeneous current distribution over the single cell in the stack, especially during self-sustained thermal operation, causes local current density overload of most electrochemical active cell areas, which can lead to accelerated cell degradation. The aging of the cell properties during operation is generally not desired. The main degradation effects take place in the anode. The reasons for the degradation of anode properties are (1) nickel agglomeration, (2) mechanical stresses induced by thermal and redox cycles, (3) contamination owing to the presence of impurities in the fuel, and (4) carbon formation. The nickel agglomeration during operation and redox cycling generally cannot be avoided. The agglomeration can be reduced by using ZrO2/NiO composites with special morphology. Generally, it is observed that increasing the nickel content in the anode increases the

agglomeration rate. However, the agglomeration of nickel, up to some extent, does not influence the polarization resistance of the electrode. Especially, thermal stresses induced by redox and thermal cycles are responsible for degradation rates of the cell. It is possible to withstand the cycling by optimizing the electrode microstructure and softening the parameters for thermal and redox cycle by operating SOFC system under reasonable conditions. The wide variation of fuels such as hydrogen, carbon monoxide, hydrocarbons, alcohols, and synthesis gases from natural gas, biogas, and petroleum makes them attractive as energy sources for utilization in SOFC. However, direct usage of different fuels also causes degradation of the cell power due to contaminants. The major contaminants that are known to poison the SOFC anodes are sulfur compounds – added to the natural gas as an odorant for safety reasons or naturally existing, for example, in biogas. Previous studies have indicated that hydrogen sulfide (H2S) among sulfur compounds is the most stable compound at the operating temperature of SOFC that causes power degradation. Many studies have also shown degradation of the cell power when the cell is exposed to H2S. These studies indicate that the sulfur poisoning increases as the H2S concentration increases and the temperature decreases. The formation of coke may occur by several mechanisms and it is difficult to identify the major source from the analysis of the products, because most coke deposits are very similar. The coke may contain 1. soot produced in the gas phase, 2. ordered or disordered carbon, produced on an inert surface, 3. ordered or disordered carbon, produced on a catalytic surface, and 4. condensed high molecular weight aromatic compounds (tar). The catalyzed formation of coke can occur on metals and on metal oxides/sulfides. Coke formation on metals is considered to be more complex. In studies of coke formation on nickel, the reactions have been suggested to involve hydrocarbons adsorbing on the surface of nickel and reacting to produce carbon atoms/groups of atoms. These may stay on the surface to encapsulate the metal (and deactivate the catalyst) or may dissolve in the nickel and migrate to growth centers such as grain boundaries. As a result, at least some of the activity of the nickel remains but the growth of the carbonaceous columns will lead to blockage. Two main factors can affect the rate of coke formation: 1. gas composition and temperature (thermodynamics in equilibrium as a driving force for coke formation), and 2. reaction rate (kinetics of reactions).

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The crucial parameter for carbon formation is reaction rate. The reaction rate highly depends on the temperature and the catalyst used. Although carbon formation at lower temperatures is thermodynamically more likely to happen than at higher temperatures, it occurs first at temperatures higher than 4450 1C because of very slow reaction rates. The nickel acts as a catalyst for carbon production. The mechanism of carbon particle growth on nickel is similar to that of carbon nanotube growth, which is well known from the literature. First, the dissolution of the released carbon in the nickel particles occurs. Then at some point, the particles become saturated with carbon. Later, the carbon particles grow and cause damage to the anode microstructure. In this way, the anode can be damaged and hence the carbon formation should be avoided. The stability of the cathode/electrolyte interface with respect to polarization and current density is the topic of numerous investigations. The changes in the polarization resistance are mainly attributed to the destruction of La2Zr2O7 nano-interlayer between LSM and YSZ particles or the creation of the additional oxygen vacancies in the bulk and at the surface of the perovskite. Most investigations in this field were performed on single-layer cathodes using the symmetrical cells, the sputtered dense microelectrodes, or MEAs under high cathodic overpotential (Zco  0.2 V) to obtain reasonable current densities. The following effects were observed at the cathode/electrolyte interface during current load:

The process can proceed according to the following scheme:

(1) reduction of the polarization resistance/activation of the electrode, (2) nanopore formation at the cathode/electrolyte interface, (3) destruction of the intermediate La2Zr2O7 layer by use of stoichiometric LSM, and (4) rapid degradation of polarization resistance by cyclic current load of single-layer cathodes.

The reversible reduction of the perovskite possibly takes place at lower current densities (p300 mA cm2) and has no effect on the microstructure resulting in electrode activation and relaxation of polarization resistance during standby operation without current. The further increase of the electrode polarization during operation leads to the decomposition of perovskite and irreversible electrode activation. The critical overpotential value for the nanostructuring of the interface depends on the interaction between perovskite and zirconia particles as well as on the particle size and number of zirconia particles surrounding the perovskite particle in the microstructure. The effect of irreversible electrode activation by high current density can be used for activation of composite cathode by high current density and nanostructuring of the cathode/electrolyte interface.

The nanopore formation at the cathode/electrolyte interface in the composite cathodes is observed under current densities X450 mA cm2. It is difficult to explain the electrochemical reduction of the uLSM ((La,Sr)MnO3 with deficit of lanthanum (lanthanum-understoichiometric, i.e., La0.65Sr0.3MnO3)) perovskite only because of the cathodic overpotential under operating conditions. Probably, the extensive interaction between uLSM and YSZ particles such as manganese interdiffusion causes transport of oxygen vacancies from YSZ to uLSM, inducing vacancies enrichment in the perovskite material. To compensate the increased vacancy concentration under electrical load, partial decomposition of the perovskite material takes place inducing pore formation and probably strontium oxide (SrO) segregation on the surface and change of manganese oxidation state in the perovskite structure.

Step 1: Reaction of uLSM with YSZ during sintering (manganese has an oxidation state 3 þ in the zirconium dioxide lattice (the ionic radius of Mn3þ (0.54 A˚) is low in comparison with the ionic radius of Zr4þ (0.84 A˚), causing a relative good solubility of MnO1.5 in zirconium dioxide)) La0:75 Sr0:2 Mn3þ x Mn4þ 1x O3d La0:75 Sr0:2 ðMn3þ x Mn4þ 1x Þ1y O31:5yd þ ðMnO1:5 ÞYSZ y

Step 2a: Reversible perovskite reduction by electrochemical overpotential La0:75 Sr0:2 ðMn3þ x Mn4þ 1x Þ1y O31:5yd La0:75 Sr0:2 ðMn3þ xþ2z Mn4þ 1x2z Þ1y O31:5yxzd þ zOoðYSZÞ

Step 2b: Irreversible perovskite reduction by electrochemical overpotential La0:75 Sr0:2 ðMn3þ x Mn4þ 1x Þ1y O2:9251:5yd -


La0:75 Sr0:2z ðMn3þ x Mn4þ 1x Þ1y O31:5yxzd þ zSrO La0:75 Sr0:2 ðMn3þ x Mn4þ 1x Þ1y O2:9251:5yd -

La0:75z Sr0:2 ðMn3þ x Mn4þ 1x Þ1y O31:5yx1:5zd þ zLaO1:5

Step 3: Reaction of free LaO1.5 or strontium oxide with zirconium dioxide with zirconate formation

Different Membrane–Electrode Assembly Types There are two general types of electrochemical single cells for SOFC systems: tubular and planar (Figure 6). The tubular cell is based on a tubular substrate and is the oldest design for cell manufacturing originating from the zirconia tubes used as oxygen sensors. The advantage


Fuel Cells – Solid Oxide Fuel Cells | Membranes Interconnection

Electrolyte Air electrode

Air flow

Fuel flow

Fuel electrode

Figure 6 Tubular and planar cell design.

Cathode interconnection





( Tubular SOFC

l electrode Anode)

) High power density SOFC

Figure 7 New designs combining features of planar and tubular geometries.

of a tubular cell is the very good mechanical stability and therefore the ability to attain very long cells (up to 1800 mm in length) delivering power up to 200 W per cell. Also fewer cells are needed for large generators. The main disadvantage is too low a power density. Tubular cells with diameters from 2 to 22 mm are reported in the literature. The bending of cells during manufacturing and operation are the main technical problems that must be solved for successful tubular stack operation. The planar cell has some advantages: a high power density and good potential for automated manufacturing using technologies such as screen printing, tape casting, and wet-powder spraying. The mechanical stability of the planar cell is always an issue and has been successfully demonstrated only by using partially stabilized zirconia as the electrolyte material. Currently, cells with lateral dimensions up to 200 mm can be produced in a partially automated process. The planarity of the cells during manufacturing, sealing, contact with the current collectors, and mechanical stability are the main issues that must be solved for successful integration into a stack. New designs combining the features of planar and tubular cells are now under development (Figure 7).

Inside the tubular and planar designs, there exist further cell types depending on the support material used (see Table 4). Generally, the electrolyte, anode, or cathode can be used as the support for depositing the other functional layers (Figures 8(a)–8(c)). Also porous metal or ceramics (Figure 8(d)) can be applied as the support for all layers. The ASCs can operate at lower temperatures with higher power densities compared with the ESCs. Obviously, the operational temperature of the ASCs is 100 1C lower than that of ESCs, and the ESCs can overcome their deficits in the power density if a thin electrolyte with a high ionic conductivity at a higher operating temperature is used. The disadvantage of the ASC is the instability to the redox cycle (the cell breaks exposing the reducing and oxidizing atmospheres (full oxidation of nickel) owing to the volume expansion of the Ni/YSZ substrate during the cycle) (Figure 9). At present, the cathode-supported cells and ASCs are the most promising for a tubular design, and ESCs and ASCs are the main cell types for the planar design. Metal-supported cells are now under development for both planar and tubular cell types. The advantage of this

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Table 4


Current density of different cell types in H2:H2O ¼ 1:1 reported in the literature

Cell type

T0 ( 1C)


Cell voltage (V)

Power density (W cm2)

uf (%)

j (A cm2)


800 850 900 700 750 800 850 900 950

0.936 0.924 0.912 0.960 0.948 0.936 0.924 0.912 0.893

0.7 0.7 0.7 0.7 0.7 0.7 0.65 0.65 0.65

0.315 0.525 0.700 0.280 0.455 0.595 0.195 0.292 0.390

23 28 36 23 23 25 80 80 80

0.450 0.750 1.000 0.400 0.650 0.850 0.300 0.450 0.600


Cathode supported Tubular OCV, open-circuit voltage.






Ni/8YSZ, Ni/CeGdO


8YSZ, 5YSZ, 10Sc1CeSZ





Figure 8 Classification of cells in order to support material. LSM, (La,Sr)MnO3; YSZ, yttria-doped zirconia.


1200 954 °C, j0.7 V = 1317 mA cm−2, R = 0.14 Ω cm−2 912 °C, j0.7 V = 1111 mA cm−2, R = 0.19 Ω cm−2 1000

850 °C, j0.7 V = 706 mA cm−2, R = 0.32 Ω cm−2 0.85

Voltage (V)

800 0.80 600 0.75 400 0.70

Power density (mW cm−2)




10Sc1CeSZ cell 0.60

0 0







Current density (mA cm−2)

Figure 9 I–V curve and power density. Mosch S, Trofimenko N, Kusnezoff M, Kellner M, and Betz T (2007) Electrochemical and microstructural characterization of the solid oxide fuel cell anode prepared by co-precipitation. In: Proceedings of the 10th International Symposium on Solid Oxide Fuel Cells (SOFC-X), vol. 7, issue 1, pp. 1547–1553, Nara, Japan, June. ESC Transactions: Pennington, NJ.


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Ceramic mass production Extrusion Substrate Drying Sintering Cell production Paste / slurry / spray powder Functional layers

Screen printing / WPS / thermal spray Sintering / annealing

Figure 10 Technological processes for tubular cell manufacturing.

Slurry manufacturing Tape casting Substrate Cutting Sintering Cell production Paste / slurry manufacturing Electrodes / electrolyte

Screen printing / WPS Sintering

Figure 11 Technological processes for planar cell manufacturing.

cell type is the possibility of low-cost production. The consequence is operation at a reduced temperature (550– 700 1C) because of steel oxidation on the fuel side. Because of the use of a porous metal as the substrate, it will be probably difficult to achieve long-term stability over 10 000 h of operation.

Manufacturing of the Membrane– Electrode Assembly The tubular and planar type cells have different technological roots for their production but both are based on well-known ceramic processing. Figure 10 exhibits a schematic way to produce a tubular cell. The main steps involved in this process are extrusion, sintering of the substrate, and coating of the substrate with functional layers using thermal spraying or a wet powder spraying process. The planar cells are made using tape casting and sintering of the substrate and screen printing or wet powder spraying of functional layers followed by sintering steps (Figure 11). The production of planar cells by thermal spraying is also possible and is well known from the literature.

The conventional technologies for the manufacturing of planar ASCs, ESCs, and MSCs differ in the order and the number of sintering steps and applied layers. Generally, the lowest costs can be achieved by using a minimum of sintering steps. Currently, at least two sintering steps for every technology are required for cell production.

Nomenclature Symbols and Units A ASR dth E

F Itot j KIC L pO2 R R0

electrode area area specific resistance (O cm2) theoretical density open-circuit cell voltage of the cell (equivalent to the Nernst potential of fuel cell) Faraday constant current current density (mA cm  2) fracture toughness electrolyte thickness oxygen partial pressure resistance max. allowed electrolyte area specific resistance

Fuel Cells – Solid Oxide Fuel Cells | Membranes RCell T T0 uf UCell Inlet Uˆ Nernst Outlet Uˆ Nernst DG DHA DHm r

cell resistance temperature (1C) operating temperature fuel utilization (%) cell voltage during operation (V) Nernst potential at the cell inlet (V) Nernst potential at the cell outlet (V) free Gibbs energy of reaction association enthalpy migration enthalpy conductivity


Anode-supported cell area specific resistance gadolinia doped ceria direct current electrolyte-supported cell (La,Ca)MnO3 (La,Sr)(Fe,Co)O3 (La,Sr)FeO3 (La,Sr)MnO3 La1  xSrxGa1  yMgyO3  d membrane–electrode assembly metal-supported cell open-circuit voltage scandia-doped zirconia solid oxide fuel cell stabilized zirconia three phase boundary thermal expansion coefficient La-understoichiometric (La,Sr)MnO3 yttria-doped zirconia

See also: Applications – Stationary: Fuel Cells; Batteries and Fuel Cells: Efficiency; Lifetime; Electrochemical Theory: Thermodynamics; Electrodes: Porous Electrodes; Electrolytes: Solid: Oxygen Ions; Fuel Cells – Overview: Modeling; Fuel Cells – Solid Oxide Fuel Cells: Anodes; Cathodes; LifeLimiting Considerations; Micro Cells; Measurement Methods: Electrochemical: Impedance Spectroscopy.

Further Reading Appel CC and Bonanos N (1999) Structural and electrical characterisation of silica-containing yttria-stabilised zirconia. Journal of the European Ceramic Society 19: 847--851. Arachi Y, Sakai H, Yamamoto O, Takeda Y, and Imanishai N (1999) Electrical conductivity of the ZrO2–Ln2O3 (Ln ¼ lanthanides) system. Solid State Ionics 121: 133--139. Badwal SPS (1992) Zirconia-based solid electrolytes: Microstructure, stability and ionic conductivity. Solid State Ionics 52: 23--32. Badwal SPS and Drennan J (1992) Microstructure/conductivity relationship in the scandia–zirconia system. Solid State Ionics 53–56: 769--776.


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Fuel Cells – Solid Oxide Fuel Cells | Membranes

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