Ishihara T. (Ed.) Perovskite Oxide for Solid Oxide Fuel Cells (Fuel Cells and Hydrogen Energy)

Подождите немного. Документ загружается.

8.3 Perovskite Chemistry

Perovskite oxide materials possess the general stoichiometry ABO

. Convention-

ally, the A cation is larger than the B cation. In the archetype, the A cation has an

oxidation state of +2 and the B cation has the oxidation state +4. These materials

comprise three different ionic species, each with its own equilibrium defect con-

centration due to three different activation energies for defect formation, which,

combined with the constraint of electroneutrality, make for diverse and potentially

useful defect chemistry, particularly when considering electronic, hole, and ionic

conduction under atmospheres of different oxygen partial pressures [13].

Perovskite oxides generally exhibit excellent thermal and mechanical stability.

They remain stable well above 10008C, and therefore the temperatures of SOFC

operation do not present a problem. This condition is in contrast with the Ni-

cermet, where nickel sintering and agglomeration are potential hazards. It is also

important that the anode exhibits good physical and mechanical compatibility

with the dense electrolyte layer onto which it is deposited. The electrolyte choice is

usually YSZ, although there is interest in alternative electrolyte materials such as

lanthanum gallate. Examination of unit-cell parameters and thermal expansion

coefficients reveals that perovskites generally show good compatibility with the

electrolyte material. Several perovskites have also been shown to be stable in the

kind of reducing environment encountered at the fuel electrode, and specific

perovskites have exhibited the levels of ionic and electronic conductivity neces-

sary to be realistically viable as an SOFC anode.

Perovskites containing a transition metal are of particular interest because of

the availability of multiple oxidation states, which facilitate elect rocatalytic

processes and provide mechanisms for electronic conductivity. For example,

under reducing atmospheres the transition metal ions cha nge to lower oxidation

states, effectively freeing up electrons to pass current. Typical examples of such

species are titanium, niobium, and vanadium. SrNbO

has an electronic con-

ductivity of 10

Scm

1

under reducing conditions [14], and Petric reported a

conductivity of 10

1

for SrVO

under similar conditions [15]. Unfortu-

nately, it was also found that these compo unds could not be fabricated in air,

ELECTROLYTE/ANODE INTERFACE

2–

–

Fig. 8.1 Enhancement of

electrode reactive surface

area for fuel oxidation in an

SOFC anode by mixed

conducting materials

8 Perovskite Oxide Anodes for SOFCs 169

and their stability under fuel electrode conditions was questionable. As well as

the increased electronic conductivity, there may also be an increase in the oxide

ion contribution to conductivity due to the formation of oxygen vacancies on

reduction, according to Eq. (8.2):

$ 2e

þ V



(8:2)

By raising the intrinsic oxygen vacancy concentrationinthismanner,thereisan

increase in available hopping sites. The increase in vacant sites facilitates oxygen

transport through the crystal and hence raises the potential for oxide ion conduc-

tivity. The coordination of these oxide vacancies to the B-site ion is particularly

important in determining mobility. Generally the B-site ion is six coordinate,

octahedral. If fivefold or fourfold coordinates are also well known for a particular

transition metal ion, then an obvious mechanism for oxide ion conduction exists;

however, this would not be the case for B-site ions such as Cr(III), which strongly

prefer sixfold coordination; this is illustrated in the chart shown in Fig. 8.2.

8.4 Doping, Nonstoichiometry, and Conductivity

Defect concentrations and hence defect chemistry of perovskites can be con-

trolled and tailored significantly by dopin g. Figure 8.3 shows schematically the

various possibilities.

tet

Fuel ?

Fig. 8.2 Comparison between different perovskite B-site ions comparing stability under fuel

conditions and ability to reduce coordination number to allow vacancy oxide ion conductivity

of transition metals in perovskite oxides

ABO

–δ

ABO

3n+1

1–x

…..

0.6

0.3

Fig. 8.3 Possible domains of

perovskite

nonstoichiometry

170 J.T.S. Irvine

By substitution of parent cations with similar-sized cations of different

valence, defects can be introduced into the structure. Oxygen ion vacancies or

interstitials can be generated by substitution of B-site ions with cations of

lower or higher valence, respectively, producing compounds of stoichiometry

(1–x)

(3 –/+d)

. A-site vacancies can be introduced by substitution of A-site

ions with cations of higher valence, giving compounds of stoichiometry of

(1–x–)

. Substitution of A-site ions with lower-valence cations results in

oxygen vacancy formation giving compounds of stoichiometry A

(1–x)

(3–d)

The effects of such extrinsic defect concentrations on both carrier concentration

and conductivity mechanism are discussed next.

Oxygen ion vacancy concentrations can be increased by partial substitution

of the tetravalent B-site ions with lower-valency cations, as shown in Eq. (8.3):

þ B

þ MO $ M

þ V



þ BO

(8:3)

is a divalent cation, and V



is an induced oxygen ion vacancy. It is

expected that these additional vacant sites facilitate oxygen transport through

the crystal by increasing the number of potential carriers.

The most important example of a perovskite that exhibits high oxide ion

conductivity when doped is lanthanum gallate, LaGaO

. Slater et al. [16] per-

formed neutron diffraction and conductivity studies on La

0.9

0.1

0.2

2.85

and found it to exhibit significantly different structure and properties when com-

pared with the undoped compound. In this case, both the A-site and the B-site have

been doped to create oxygen vacancies, the resulting material possessing an oxide

ion conductivity of 6.6  10

2

at 1000 K. Figure 8.4 shows schematically structural

changes that may account for the dependence of activation energy on temperature

involving the tilting of GaO

octahedra. The oxide ion conductivity attained is

25°C 1000°C

[001]

Fig. 8.4 Comparison between structure of La

0.9

0.1

0.8

0.2

2.85

at ambient temperature

and 10008C illustrating significant changes in octahedral tilting with temperature: views down

[110]

, La atoms in black, O atoms, and MO

octahedra shown [16]

8 Perovskite Oxide Anodes for SOFCs 171

higher than YSZ at the same temperature and has sparked considerable ongoing

research into its application as a SOFC electrolyte [16–18].

Strontium titanate is an archetypal example of perovskite and exhibits a wide

range of defect chemistry that aptly illustrate the different factors that may

influence electronic conductivity. The effects of changing oxygen partial pres-

sure upon undoped and differently doped strontium titanates are shown in

Fig. 8.5. Important aspects are the extended p-type behaviou of the B-site-

doped sample at higher P

, the p-type behavior of the undoped sample at

higher P

, and the very high n-type conductivity of the A-site-deficient sample

at lower P

values. Equation 8.4 shows how such a doped material is likely to

exhibit p-type conductivity at the expense of ionic conductivity in high oxygen

partial pressures. The ambient oxygen atoms fill the positively charged vacan-

cies, generating a pair of holes in accordance with electroneutrality.

þ V



$ O

þ 2h



(8:4)

Equations (8.3) and (8.4) combine to give Eq. (8.5):

$ M

þ 2h



(8:5)

The p-type behavior of the undoped sample at high P

has various possible

explanations, the most likely being simply that equilibrium intrinsic oxygen

–4.50

–4.00

–3.50

–3.00

–2.50

–2.00

–1.50

–1.00

–0.50

0.00

0.00–5.00–10.00–15.00–20.00–25.00

log (P

)

log (conductivity)

5% Bsite mg @ 835°C

undoped @ 930°C

Slope = – 0.210

Slope = –0.246

Slope = 0.201

0.4

TiO

Fig. 8.5 SrTiO

, conductivity variation for different doping scenarios with oxygen partial

pressure at 9308C [13]

172 J.T.S. Irvine

vacancies are filled by ambient oxygen atoms generating holes. The high n-type

conductivity at low P

seen for the A-site-deficient sample can be explained

with reference to the equilibria described by Eqs. (8.6) and (8.7).

þ V



$ O

(8:6)

þ A

$ AO þ V



þ V

(8:7)

The large value of V

achieved by doping reduces the number of intrinsic

Schottky defects pushing Eq. (8.7) to the left and hence reducing V



. For a

given P

, Eq. (8.6) shifts left to oppose the change, facilitating the removal of

lattice oxygen by hydrogen and the associated generation of free electrons.

Figure 8.6 shows the temperature dependence of the n-type conduction in a

strontium titanate, which indicates metallic behavior [12].

8.5 Perovskite Anode Materials

Perovskite oxides can accommodate a large content of oxygen vacancies; hence,

some perovski tes are good oxygen ionic conductors. The small B site in the

perovskite allows first-row transition elements to be introduced in the lattice.

These elements exhibit multivalency under different conditions, which may be a

source of high electronic conductivity. Good ionic and mixed conductivity is

thus found in several perovskite oxides. As alread y mentioned, such mixed

conductivity is beneficial to electrode performance. P-type perovskite materi-

als are widely considered for SOFC and other applications [19]. Mixed con-

ducting perovskites, such as La

1–x

MnO

with modest oxide ionic conduc-

tivity or La

1–x

with quite high oxide ionic conductivity, have

been used as SOFC cathode materials [8, 20]. La

1–x

CrO

(M ¼Ca, Sr), a

purely electronic conductor, has also been widely used as the interconnector

for SOFCs [8].

0.05

0.1

0.15

0.2

0.25

1000800600400200

temperature (C)

Resistivity

Fig. 8.6 Resistivity (in

ohm-cm) vs. temperature for

A-site-deficient

0.875

0.75

0.25

(3–d)

exhibiting metallic

conductivity [12]

8 Perovskite Oxide Anodes for SOFCs 173

Perovskites have also been widely investigated as potential SOFC anode

materials. Among these materials, chromites and titanates are promising [21,

22]. Interesting results have been obtained with lanthanum strontium titanates

[23] and especially cerium-doped lanthanum strontium titanate [24]; however, it

is now thought that the cerium-doped anodes are in fact two phases consisting

of a ceria–perovskite assemblage [24].

It was also reported that Y-doped SrTiO

exhibits high electrical conduction

under SOFC anodic conditions [25–27]. For example, the optimized composition

of Sr

0.86

0.08

TiO

3–d

exhibits a conductivity of 82 S/cm at a P

of 10

–19

atm at

8008C. However, the sample was pre-reduced in pure argon or 7% H

/Ar at

14008C before cond uctivity measurements. It is supposed that the conductivity

of the materials would be significantly lower if the sample were only reduced below

10008CinthiscaselessTi

4þ

was reduced to Ti

3þ

, which is the source of the high

electronic conductivity. The high-temperature pre-reduction process for such tita-

nates makes it diff icul t to co-fir e the anode and cathode. The conductivity of

0.86

0.08

0.9

0.1

is only about 1–2 S/cm when reduced in situ in 5% H

9008C [28]. No phase changes were found for a mixture of Y-doped SrTiO

(SYT)

with YSZ or LSGM on calcining at 14008C for 10 h, indicating good chemical

compatibility between the SYT and electrolyte materials. The conductivity of

SrTiO

in a reducing atmosphere can also be improved by replacing titanium

with some niobium. For charge compen sation, the strontium content at the A

site should decrease. Good electrical conductivity was observed for Sr

1–x

1–x/

3–d

(x  0.4) [29] on reduction in low oxygen partial pressure, with a

maximum for the sample with x ¼0.25, s ¼5.6 S/cm at 9308C(P

¼10

18

atm).

Lanthanum strontium titanates are usually treated in the literature as simple

cubic perovskites, although the presence of extra oxygen beyond the ABO

stoichiometry plays a critical role in both the structure and the electrochemical

properties, as summarized in Fig. 8.1. The lower members of the La

n-4

3n+2

series, n < 7, are layered phases, having oxygen-rich planes in the

form of crystallographic shears joining consecutive blocks. These planes become

more sporadic with increasing n (i.e., decreasing the oxygen content) until they

are not a crystallographic feature, rendering local oxygen-rich defects randomly

distributed within a perovskite framework, n > 11 [30, 31]. The presence of

such disordered defects appears to strongly affect the redox characteristics of

the oxide, as indicated by marked effects on conductivity induced by mild

reduction (Fig. 8.7). Unfortunately, although the materials from this lantha-

num strontium titanate oxygen excess series are much easier to reduce, and

hence exhibit much higher electronic conductivity than their oxygen stoichio-

metric analogues, they do not exhibit very good electrochemical performance

[32]. This detriment is attributed to the inflexibility of the coordination

demands of titanium, w hich strongly prefers octahedral coordination in the

perovskite environment.

To make the B-site co ordination more flexible and hence to improve electro-

catalytic performance, Mn and Ga wer e introduced to replace Ti in La

38–z

-based fuel electrodes. Mn supports p-type conduction in oxidizing

conditions and has been previously shown to promote electroreduction under

174 J.T.S. Irvine

SOFC conditions [33]. Furthermore, Mn is known to accept lower coordination

numbers in perovskites [34], especially for Mn

3þ

, and thus it may facilitate

oxide ion migration. Similarly, Ga is well known to adopt lower coordination

than octahedral in perovskite-related oxides. The possibility of mixed ionic/

0.25

∞

0.20

0.15

0.10

0.05

0.00

–5.00

0.1 0.2 0.3 0.4 0.50

–4.00

–3.00

–2.00

–1.00

1/n

Log σ\Scm

–1

Sc series

Local

defects

Extended defects

Layered

n>12

n<12

= 12

5nm

4nm

Fig. 8.7 High Resolution Transmission Electron Microscopy (HRTEM) images (top)ofsam-

ples from the ‘‘La

n–4

3n+2

’’ series varying from ordered extended planar oxygen excess

defects (1; n ¼5) through random layers of extended defects (2; n ¼8) to disordered extended

defects (3; n ¼12). The middle picture shows the location of these phases on the composition

map with 1/n plotted against oxygen excess, d, in perovskite unit ABO

3+d

.Thebottom diagram

shows defect electronic conductivity (s/cm

1

) of grain component as determined by ac

impedance spectroscopy on samples quenched from 13008C in air, also plotted against d [31]

8 Perovskite Oxide Anodes for SOFCs 175

electronic condu ction is very important because it would allow the electro-

oxidation process to move away from the three-phase electrode–electrolyte–gas

interface onto the anode surface, with considerable catalytic enhancement.

The anode polarization resistance was measured using three-electrode geo-

metry. By optimization of electrode microstructure, polarization resistances in

wet H

were 0.12 O cm

in wet H

, 1.5 O cm

in wet 5% H

, and a remarkably

low value, 0.36 O cm

, in wet CH

, at 9508C. These polarization resistances were

attained after about 24 h in fuel conditions; initial polarisation resistances were

two to three times higher. This long time period to achieve equilibration is fairly

typical for donor-doped strontium titanates that are not cation vacancy com-

pensated, and we attribute this to reorgani zation of a complex defect structure.

The open circuit voltages (OCVs) matched the value predicted by the Nernst

equation, 0.97 and 1.13 V at 9508C, for wet 5% H

and wet H

, respectively. The

OCVs in wet CH

, for a one-layer 50:50 YSZ:LSTMG anode, were: 1.39 V at

9508C, 1.32 V at 9008C, and 1.36 V at 8508C. These values were reproducible

after 2 days of testing in wet 5% H

, wet H

and wet CH

[35].

Although SrVO

shows excellent electronic conductivity of 1000 S/cm at

8008C and an oxygen partial pressure of 10

20

atm, it is unstable under a more

oxidizing atmosphere [36]. The conductivity in a reducing atmosphere may

drop rapidly if strontium at the A site is partially or completely replaced by

lanthanum, and neither is the stability in oxidizing condition improved. A

material that has both adequate high-temperature conductivity in a reducing

atmosphere and redox stability has not been found in the vanadates. In addi-

tion, the reduction process of SrVO

is rather slow once the material is oxidized.

The perovskite oxide La

0.6

0.4

0.2

0.8

was proposed as an SOFC

anode at intermediate temperatures (5508–7008C) [37]. The stability of these

materials under fuel atmosphere is in doubt, however, even at such low tem-

peratures. SrFeCo

0.5

exhibits both high electronic and ionic conductivities in

air and is applicable for ceramic membrane used for gas separation [38]. It has

been reported that mixed conductors SrFeCo

0.5

, SrCo

0.8

0.2

3–d

and

0.6

0.4

0.8

0.2

3–d

may be used as SOFC anode materials. Although

there remain some questions about the exact structure of SrFeCo

0.5

, these

are generally thought to be related perovski te/brownmillerite intergrowths of

the Grenier type [39]. The performance was not ideal when using only these

mixed conductors as an anode; however, the performance was impr oved when

these mixed conductors were used as an interlayer between an Ni-YSZ anode

and YSZ electrolyte. The long-term stability would again be a problem because

SrFeCo

0.5

is unstable under anodic conditions [40].

As stated above, LaCrO

-based materials have been investigated as inter-

connect materials for SOFCs [8]; however, they are also potential anode mate-

rials for SOFCs due to their relatively good stability in both reducing and

oxidizing atmospheres at high temperatures [41]. The reported polarization

resistance using these materials is too high for efficient SOFC operation,

although significant improvements have been achieved using low-level doping

of the B site. As no significant weight loss was observed when LaCrO

was

176 J.T.S. Irvine

exposed to a reducing atmosphere (P

¼10

21

atm at 10008C) [42], this

indicates that chromium strongly retains its sixfold coordination. Indeed,

III

is well known to strongly prefer sixfold coordination in its chemistry;

thus, it is difficult to introduce the oxygen vacancies that are required for

oxygen ion conduction into the LaCrO

lattice.

When the B sites are doped by other multivalence trans ition elements that do

tolerate reduced oxygen coordination, such as Mn, Fe, Co, Ni, and Cu, oxygen

vacancies may be generated at the B-site dopa nts in a reducing atmosphere at

high temperature. Thus, a significant degree of B-site dopant is required to

generate a percolation path for oxygen vacancies to achieve high oxygen ion

conductivity. Quite a lot of attention has been focused on 3% replacement of Cr

by V, and although methane cracking seems to be avoided [43], the polarization

resistance is still of the order of 10 O cm [2, 21, 38, 44]. The introduction of other

transition elements into the B site of La

1–x

1–y

(M ¼Mn, Fe, Co, Ni)

has been shown to improve the catal ytic properties for methane reforming [45].

Of the various dopants, nicke l seems to be the most success ful, and the lowest

polarization resistances have be en reported for 10% Ni-substituted lanthanum

chromite [46]; however, other workers have found nickel evolution from 10%

Ni-doped lanthanum chromites in fuel conditions [47]. Certainly nickel oxides

would not be stable in fuel atmospheres, and although the nickel may be

stabilized by the lattice in the higher oxidation state, there will always be the

suspicion that the activity of nickel-doped perovskites is the result of surface

evolution of nickel metal and hence questions about long-term stability. A

composite anode of 5% Ni with a 50:50 mixture of La

0.8

0.2

0.8

0.2

and Ce

0.9

0.1

1.95

was successfully used for SOFCs with different fuels [48].

The combination of both titanium and chromium at the B site of a perovskite

has also been investigated [49, 50]. The highest conductivity in 5% H

/Ar of 5 S/

cm at 10008C was observed with composition La

0.4

0.6

0.2

0.8

3–d

; how-

ever, the catalytic effect of these materials for oxidization of hydrogen at the

anode is possibly not ideal, because a large anode polarization resistance was

observed when La

0.7

0.3

0.8

0.2

was applied as the SOFC anode [49].

8.6 A(B,B

Perovskites

As perovskites with one cation occu pying the B site have not generally yielded

good enough properties for efficient anode operation in SOFCs, an extensive

series of studies has considered the possibility of enhancing performance by

using two different B-site ions, both with concentration in excess of percolation

limit (i.e., >30%). The objective is to obtain complementary functionality from

appropriate cation combinations, hopefully without seriously degrading the

good properties induced by the individual ions. Not surprisingly, many of the

tested combinations did compromise properties, but in some important

instances good complementary functionality has been achieved.

8 Perovskite Oxide Anodes for SOFCs 177

Early studies focused on double perovskites with niobium and a first-row

transition metal or main group ion occupying the B site. With Nb and Mn

occupying the B site, electronic conductivity is fairly low, probably reflecting

the rock salt-type ordering of the B cations [51]. Using Cu and Nb somewhat

improves conductivity in air, but in reducing conditions copper metal is

exolved and conductivity is impaired, as the resultant perovskite is more

resistive and the copper does not form a conducting network [52]. Using Ga

with Nb again results in an ordered superstucture that impairs electronic

conductivity [53].

Improved performance has been obtained with complex perovskites based

upon Cr and Mn at the B sites forming compositions (La,Sr)Cr

1–x

3–d

[54].

Previous workers have focused upon doped lanthanum chromite, where doping

is used in the solid-state chemical sense of up to 20% dopant on the B site,

usually 5% or 10%. (La

0.75

0.25

)Cr

0.5

(LSCM) exhibits comparable

electrochemical performance to Ni–YSZ cermets. The electrode polarization

resistance approaches 0.2 O cm

at 9008C in 97% H

/3% H

O. Very good

performance is achieved for methane oxidation without using excess steam . The

anode is stable in both fuel and air conditions and shows stable electrode

performance in metha ne. Thus, both redox stability and operation in low-

steam hydrocarbons have been demonstrated, overcoming two of the major

limitations of the current generat ion of nickel zirconia cermet SOFC anodes.

Catalytic studies of LSCM demonstrate that it is primarily a direct oxidation

catalyst for metha ne oxidation as opposed to a reforming catalyst [55], with the

redox chemistry involving the Mn–O–Mn bonds [56]. Although oxygen ion

mobility is low in the oxidized state, the diffusion coefficient for oxide ions in

reduced LSCM is comparable to yttria-stabilized zirconia [57].

Another important double perovskite is Sr

MgMoO

6–d

, which has recently

been shown to offer good performance, with power densities of 0.84 W/cm

and 0.44 W/cm

in CH

at 8008C, and good sulfur tolerance [58]. The

molybdenum-containing double perovskite was initially prepared at 12008C

in flowing 5% H

and then deposited on top of a lanthanum ceria buffer layer

before testing [59].

8.7 Tungsten Bronze Anode Materials

Tungsten bronze-type materials have also been investigated as potential SOFC

anodes. The tungst en bronze structure can be obtained from the perovskite by

rotation of some of Ti/NbO

octahedra: 40% of the A sites (A2 sites) are

increased in size from tetracapped square prisms to pentacapped pentagonal

prisms, 40% remai n essentially unchanged, and the remaining 40% of the sites

is de creased in size (Fig. 8.8). The formula may be written as A

0.6

when the

small-size A sites are left empty. The distortion of the octahedra means that

some B–O bonds are extended and some are shorter than the average. The

178 J.T.S. Irvine