Science of Ceramic Interfaces II J. Nowotny (Editor) �9 1994 Elsevier Science B.V. All rights reserved.

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INTERFACES IN ZIRCONIA BASED ELECTROCHEMICAL SYSTEMS AND THEIR INFLUENCE ON ELECTRICAL PROPERTIES

S.ES. Badwal and J. Drennan

CSIRO, Division of Materials Science & Technology Private Bag 33, Rosebank MDC, Clayton 3169, Victoria, Australia

Ceramic Fuel Cells Ltd Monash Science and Technology Park 710 Blackburn Road, Clayton, Victoria 3168

Abstract The interfaces play a critical role in establishing both electrical and electrochemical pro-

perties of zirconia based electrolyte cells. Various types of interfaces which exist in such systems have been described. Apart from gas/electrode and electrode/electrolyte interfaces, the role of interfaces within the grains (coherent growth of precipitates, compositional variations within grains and second phase inclusions) and at grain boundaries (intermediate phase formation from the matrix, intermediate phase of the impurity type, and inclusions and pores) of polycrystalline electrolyte materials has been elucidated. The effect of inter- faces on electrical and electrochemical properties has been discussed.

1. INTRODUCTION

Zirconia based materials, apart from their use in mechanical applications, are commonly used in solid oxide fuel cells for power generation; in oxygen sensors for combustion, pollution and process control; oxygen pumps for generation of oxygen or oxygen atmo- sphere control; and electrochemical reactors for production of chemicals [1-5]. The inter- faces play a crucial role in establishing both electrical and electrochemical properties of zirconia-based electrolyte cells. In broad terms, processes contributing to overall charge transport through a solid electrolyte device consist of electrode reactions at both electrodes and ionic transport within the solid electrolyte. The electrode reactions take place at the electrode/electrolyte interface. The types of interfaces which contribute to electrode reac- tions are: gas/electrode, electrode/electrolyte, gas/electrolyte and three phase boundary between the gas phase, electrode and electrolyte. Impurity segregation can significantly modify the external surface of the electrolyte. The efficiency of electrode reactions apart from material properties and electrode morphology therefore depends significantly on the state of external surface of the electrolyte. Following the oxygen exchange reaction, ionic transport takes place within the solid electrolyte. For practical reasons, polycrystalline materials are invariably used and this further increases the number of interfaces within the system. Ionic transport can be divided into two main components: through the

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grains and across grain boundaries. The grains in the ceramic may have inclusions, pores, precipitates of other phases, and compositional inhomogeneities. They all create additional interfaces which can influence ionic transport within the grains. Grain boundar- ies may have secondary phases, inclusions, pores and enhanced segregation of matrix components. Depending upon the conducting properties of these phases and their quant- ity and location, the effect on the electrical or electrochemical performance of the device may be quite significant. Any interface which hinders mass transport as a result of applied potential is detrimental to the performance of an electrochemical device. In this paper, the role of different types of interfaces which exist in zirconia based electrolyte cells has been considered and their influence on electrical and electrochemical properties has been discus- sed. The main sections under which interface properties have been discussed are: electrode reactions at the electrode/electrolyte interface; oxygen-ion transport within the solid elec- trolyte (interfaces within grains); and grain boundaries.

2. GENERAL COMMENTS

A number of figures in this chapter are impedance diagrams. For readers not familiar with the topic a simple explanation is given below. In an electrochemical cell, as discussed above, overall mass transport consists of a number of different processes at the electrode/electrolyte interface and within the solid electrolyte. At the electrode/electrolyte interface, typical electrode processes are adsorption/dissociation, diffusion and charge transfer and within the electrolyte the main processes involve oxygen-ion migration through the lattice and across grain boundaries. Each of these processes has a different time con- stant associated with it and occurs over a different time or frequency domain and can be simulated by a network of resistors and capacitors. Figure 1 shows a simple electrical equivalent circuit of a solid electrolyte cell. The response of such a system in the fre- quency domain when represented in the impedance plane consists of a number of arcs each corresponding to a different process. This is shown schematically in Figure 1. Electrode processes, in general, have a large time constant and invariably relax over the low frequency domain followed by ion migration across grain boundaries and bulk diffusion of ions in the crystal lattice. One or more arcs are usually observed for electrode reactions in the complex impedance plane as shown in Figure 1 [6-10]. For a polycrystalline electro- lyte, two arcs are commonly observed. The low frequency electrolyte arc is associated with the migration of oxygen ions across grain boundaries and build-up of space charge layers near the grain boundary surface region. In a single crystal or grain, this arc is not ob- served. The high frequency electrolyte arc is due to dipole relaxation within the bulk of grains. The difference of intercepts of each arc on the real impedance or x-axis defines the resistance or rate of that particular process. The time constant for each process can be determined from the frequency at the top of each arc. Due to selection of measurement conditions and availability of the frequency range or instrumentation, all arcs due to elec- trode or electrolyte processes may not be seen or only partly seen in the impedance plane. However, this does not imply that a process is not contributing to the overall mass transport in a solid electrolyte cell. Measurements over a wide frequency, temperature and gas concentration are required to fully elucidate the role of rate limiting processes. For more detailed explanation of impedance or admittance diagrams, the readers, for example, should read [7].

73

m

N I

(a) C R

R R

(b)

Cgb Cel

t

Rgb Rel

Frequency I

I

Cad

Rad

Electrode T~ Electrolyte

Tgb

R~ R f + Rg b R~ + Rg b+ Rel + Rad I

Z

Figure 1. A simple schematic impedance plane representation of an electrochemical cell. The cell response is shown as consisting of two electrode (adsorption (Rd, C~d) and charge transfer (Re1, C1)) and two electrolyte (grain boundary (R_, C .) and lattice (R t Cd) processes and is not a true representation for every electrode/~ectr~yte system.

3. ELECTRODE REACTIONS AT ELECTRODE/ELECTROLYTE INTERFACES

In zirconia based electrolyte systems the oxygen reduction reaction take place at the cathode/electrolyte interface. The oxygen ions thus produced transport through the electro- lyte and discharge at the anode/electrolyte interface. The discharge reaction may be simple evolution of oxygen (in the oxygen pumping mode) or it may involve complex mechan- isms where discharging oxygen ions react directly or indirectly with fuel species (fuel cell operation) or chemicals (electrochemical reactor) [1-4, 11-15]. The state of both cathode/electrolyte and anode/electrolyte interfaces and physical, chemical and thermal compatibility of the electrode and electrolyte interfaces are critical for long term and satis- factory performance of the device.

At the cathode oxygen reduction takes place through the normal processes of gaseous diffusion, adsorption/dissociation at the electrode or electrolyte surface, diffusion of ad- sorbed electroactive species to the reaction site and charge transfer reactions. The adsorp- tion/dissociation reaction may proceed via simultaneous transfer of charge at the gas/elec- trode interface. At the anode/electrolyte interface, oxygen ions discharge releasing elec- trons to the electrode in the normal mode of operation (e.g. oxygen pumping mode). However, in the presence of oxidisable gases (e.g. CO, H E, CH 4, other gaseous chemicals), the reaction mechanism is complex and may involve adsorption/dissociation of fuel species at active sites on the electrode or electrolyte surface [14,15]. The electroche-

74

mical processes taking place in a solid electrolyte cell are shown schematically in Figure 2. The nature of the electrode material, its electronic and ionic conductivity, surface area and porosity, play a significant role. For many electrode materials which have a low solu- bility for oxygen or which have low oxygen-ion mobility, the role of the three phase boundary (tpb) between the gas phase, electrode and electrolyte is critical [16,17]. The constraints imposed by the existence of tpb reduce electrode/electrolyte contact area and requires the electrode microstructure to be carefully optimised. For this reason electrode materials with high oxygen-ion conductivity are preferred as the requirement of tpb is considerably reduced. Also the adherence of the electrode material to the electrolyte is critical for better contact between electrode and electrolyte. The electrodes must have minimum thermal expansion mismatch with the electrolyte and other cell components. A large mismatch can lead to delamination or cracking of electrode layers causing loss of the contact area and poor definition of the interface. The electrode morphology also plays a

Figure 2. A schematic of electrochemical processes occurring in a solid electrolyte cell.

significant role in defining reaction rates and can not be ignored. The electrode pore structure must allow reactive gases to diffuse through to reaction sites and allow unreacted gases and reaction products to escape freely without any hindrance.

The main types of interfaces, baring the physical appearance (microstructure, porosity and surface area of the electrode) are: relatively clean interfaces (no impurity segregation or an intermediate phase formation); existence of an interphase due to segregation; exis- tence of an interphase due to reaction/diffusion between cell components; and in the case of multi phase electrodes, interfaces between components of a composite electrode. Various types of interfaces which can have an affect on the electrode performance are schematically depicted in Figure 3. The role of electrode morphology and different types of interfaces

75

have been discussed below.

3.1 Electrode Morphology The microstructure, porosity and surface area of the electrode define the contact area

between gas/electrode and electrode/electrolyte interfaces and play a crucial role in defin- ing electrode kinetic rates both in the presence and absence of an interphase. The electrodes which have low oxygen-ion conductivity or low solubility for oxygen require optimum porosity for gas diffusion so that the maximum number of triple phase boundaries between gas, electrode and electrolyte are available. Noble metals (Ag, Au, Pt, Pd), especially Pt, are good catalytic electrodes for oxygen reduction, exhibit good stability in both oxidising and reducing environments and have been commonly used in oxygen sensors, oxygen pumps and fuel cells. However, these metals undergo sintering and grain growth when exposed to temperatures in the vicinity of 700 - 1000~ This leads to a substantial reduc- tion in the gas/electrode interface and three phase contact area between gas/electrode/elec- trolyte (essential for oxygen charge transfer reaction) [18,19]. Figure 4 shows the micro- structure of Pt paste and Pt electrodes sputter deposited on zirconia-yttria electrolyte and

@

iiiiiiiiiii il |

l iiiiiiiiiiii [ / 1 1 I I / 1 t

| | @

i ii!i so,acoE'~176176176

Electrolyte

Figure 3. Various types of interfaces in solid electrolyte cells which affect electrode per- formance. (1) Segregation of electrolyte components; (2) impurity segregation at the sur- face of the electrolyte; (3) reaction products between electrode and electrolyte phases; (4) non-contact region (dead zone) for air or fuel electrode reactions; and (5) gas/electrode and clean electrode/electrolyte interfaces.

given heat treatment at different temperatures in air. The grain growth and reduction in the porosity and three phase contact area with increasing temperature is clearly obvious and has a substantial effect on the interfacial electrode resistance. The effect of the microstruc- ture on the electrode resistance is shown in Figure 5 for platinum electrodes which have low solubility for oxygen and require the presence of a large three phase contact boundary area between electrode, electrolyte and the gas phase. With decreasing electrode surface area and three phase boundary (due to the high temperature treatment), both the electrode

76

Figure 4. Microstructure of sputter deposited Pt: (a) & (b), and Pt paste: (c) & (d), elec- trodes on zirconia-yttria substrate and heat treated in air at 600~ (50h): (a) & (c) or 900~ (2h): (b) & (d).

resistance and the time constant associated with the electrode processes increase and differ- ent processes become rate limiting as shown in Figure 5 by the appearance of another arc on the low frequency end of the impedance spectrum. This low frequency arc is a function of the electrode surface area and is generally attributed to slow adsorption/dissociation of

77

oxygen at the electrode surface [20,21]. In general, oxide electrodes are more refractory and undergo less sintering and grain

growth compared with metals. Nevertheless degradation in the electrode performance does occur with time. The common cause of degradation in oxide electrodes with low oxygen- ion conduction, in the absence of a chemical reaction with the electrolyte phase, is also sintering and grain growth leading to a reduction in the three phase boundary contact area. For example, Sr doped LaMnO~ is used as the air electrode in solid oxide fuel cells. Al- though the electronic conductivxty is high (about 190-200Scm 1 for (La o ~Sr o 2)MnO3 at 1000~ the oxygen-ion diffusion is extremely slow and the presence-ofthree phase boundary region and large porosity are crucial to low overpotential losses [22,23]. How- ever, if oxygen electrodes (cathodes) were available with high bulk oxygen-ion diffusion rates as well as fast surface exchange kinetics at the gas/electrode interface, the role of three phase boundary and porosity would not be as critical and the contact area between electrode and electrolyte can be enhanced by using dense electrode materials. For such electrodes, oxygen exchange with the gas phase would take place at the gas/electrode inter- face followed by oxygen-ion transport within the electrode and transfer at the electrode/electrolyte interface. Several materials with the perovskite and related structures have high oxygen-ion diffusion rates and high oxygen-ion conductivity [24] and for these materials, the rate of the oxygen transfer reaction is expected to be established by the sur- face exchange properties of the electrode material. Carter etal [25] and Steele etal [26]

r

E r

e,,

b~ !

r

b~ !

6000

0 4000

2000

I (c) Pt paste 6926 1 0~ . . . . . o . . . . . o o

o o o o ~

F ,ioo o+ooo~176176 o ~ / . 2 ~ ~ ~ ~ 0 m

o ~"V.~AA.

I ~ I oOm_./t i ] 6000 12000 18000

100% 02 600 ~

. "0 , / lOOm O o

o o

Oo

o

lOre Ooo l m

I 2 4 0 0 0 : 3 0 0 0 0

sputtered Pt

(c) o o O O o o o �9 o o O o

o o O o o ~ �9 �9 �9 �9 -- o � 9 ~ 1 4 9 o

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�9 I

F �9 . . . . . . . (b) 1.25m

i ~ lOOm lOOm ~ l I I

2000 4000 6000 8000

-Z' ~cm2

Figure 5. Effect of electrode surface area and three phase boundary on the electrode resis- tance and behaviour for two Pt/zirconia-yttria/Pt ceils. The data were recorded at 600~ in 100% 02, after heat treatment of the cell at: (a) 600~ (b) 750~ and (c) 900~ The numbers on the arcs are frequencies in Hz.

78

have measured bulk diffusion coefficients for oxygen ions in the electrode and surface exchange coefficients at the gas/electrode interface for both lanthanum manganite and lanthanum cobaltite perovskite electrodes doped at A- and B-sites and have shown that for doped LaCoO 3 electrode materials, oxygen-ion diffusion rates are 4-6 orders of magnitude higher than those observed in lanthanum manganites with similar doping levels. These authors have demonstrated that although the surface exchange coefficients in cobaltite based perovskites are higher by one to three orders of magnitude compared with the manganites, the oxygen flux through the material is limited by the oxygen exchange coefficient at the gas/electrode interface. In lanthanum cobaltites, the net oxygen flux through the material (or the electrode reaction rate) is controlled by the surface exchange reaction whereas in manganites the process is bulk diffusion controlled. Thus it appears that optimisation of the microstructure is required for both types of materials.

Similarly for Ni/zirconia-yttria cermet, which is the common fuel electrode material used in solid oxide fuel cells, the electrode microstructure plays a dominant role in determining the electrode performance at the interface. Not only fuel gases (H 2, CO) need to diffuse in to the electrode/electrolyte interface, but products of reaction (steam and CO~) are required to diffuse out as well. About 50% or higher porosity is required and pore size and distri- bution are crucial to the performance of the electrode. Zirconia-yttria is added to Ni to reduce the thermal expansion mismatch at the interface and to slow down sintering and grain growth of the Ni particles. Therefore, the distribution and the amount of Ni and zirconia-yttria phases need to be carefully controlled in the electrode to avoid performance degradation of the cell at the anode/electrolyte interface. The thermal expansion of Ni metal is significantly higher than that of zirconia-yttria. This combined with poor wetting at the Ni/electrolyte interface can lead to debonding or delamination at the interface, poor microstructure and higher overpotential losses [27,28].

3.2 Relatively clean interfaces The electrode/electrolyte interfaces free of impurities and other phases can truly define

the electrode kinetic behaviour of the system and are generally preferred unless the forma- tion of another phase somehow facilitates electrode reactions (see below). However, for all practical purposes clean surfaces or interfaces are rarely encountered. Even in the absence of segregated impurities and secondary phases, particulates in the gas stream, unwanted reaction products (coke deposition in fuel ceils in carbon containing atmospheres) and impurities in the gas phase (sulphur containing compounds, metal vapours etc) can signific- antly affect the performance of electrodes. Hydrogen embrittlement occurring in metallic electrodes is another common problem affecting the electrode behaviour. For example, the adsorption of very small (ppm levels) quantities of fuel gases even in oxidising environ- ments (in the reference air stream for example) is known to poison the electrode and has been shown to adversely affect the performance of oxygen sensors especially at lower temperatures [29-31]. In potentiometric sensors, large deviations from the Nernstian beha- viour are commonly observed for Pt electrodes below about 400 - 500~ as shown in Figure 6. This type of behaviour is related to the presence of trace levels of CO and hydrocarbons in the gas stream and their adsorption on the surface of Pt electrodes giving rise to mixed potentials [29,31,32]. Sensors prepared with platinum electrodes with a low electrode surface area (heat treated at high temperatures: above about 800~ show much larger non-Nernstian behaviour [30] possibly due to high sensitivity of the reduced surface area electrodes to trace levels of combustibles in the gas stream. When Pt elec- trodes are replaced with materials such as nonstoichiometric oxides (e.g. (U,M)O2• ~) with apparently low adsorption for these gases or which facilitate reaction between trace mapur-

79

>

E uS

Pf paste 6082

45

30

15

l , _ A , I . . , . . . . . . I . L ~ . . . .

370 410 450 490 530 570 610

T, ~

Figure 6. Non-ideal behaviour of Pt electrodes in potentiometric sensors at low tempera- tures. Air versus 5.01~ in N 2. After the heat treatment of electrodes at (O) 600 ~ (zx) 750~ and (<>) 900~ in air. 15igure courtesy of F.T. Ciacchi.

ity gases and oxygen, the non-ideal behaviour is considerably reduced [30]. Thus even for clean electrode/electrolyte interfaces, electrode morphology and interference from external factors can substantially modify the electrode performance.

3.3 Existence of interphase due to segregation The starting powders used to make zirconia ceramics invariably contain small amounts of

impurities. These impurities are either deliberately added as sintering aids, or are present as contamination picked up during the processing or are present because of the excessive cost of cleaning the raw materials. During the sintering of zirconia ceramics, complex chemical reactions occur between adsorbed or segregated impurities and often involve bulk matrix components leading to the formation of glassy phases [33-35]. These phases are quite mobile (some with a melting point below 1000~ and move around rapidly in grain boundaries and tend to migrate to the external surface during sintering [36-37]. Even if the surfaces of zirconia ceramics are cleaned free of impurities by chemical etching or a mechanical grinding process, during subsequent heat treatments, annealing or use of the ceramics at high temperatures (800-1200~ more impurity phases migrate to the external surface from the grain boundary network. Hughes and Badwal, from x-ray photoelectron spectroscopy studies performed on zirconia-yttria ceramics with different impurity content have shown that most of the impurities which are present in the starting powders migrate to the external surface [38,39] during sintering or subsequent annealing. In addition, enhanced segregation of dopant (yttrium) is commonly observed at the sintered surface. Figure 7 shows the impurity segregation at the external surface of a 3 tool% Y.O 3 - ZrO. ceramic (previously sintered at 1500~ ground and polished) annealed at 1200~ for 50h. ZDepend_ ing on the quantity and type of impurities, significant surface layers can build up. Ceramic electrolytes prepared from commercial powders have been shown to have texture at the

80

L .

N

I-- O

Z e,

U3

0.30

0.25

0.20

0.15

0.10

0.05

3 tool% Y203-Z r02

0.00 I ~ , , i , , , ,

- 1 0 90 190 290 390

Sput ter ing Time, Sec

1200~

Figure 7. Yttrium and impurity segregation measured by x-ray photoelectron spectroscopy at the external surface of a 3 tool% Y203 - ZrO 2 ceramic (the sintered surface cleaned by grinding and polishing after sintering at 1500~ annealed at 1200~ for 50 h in air. O- Y/Zr; A- Si/Zr; O- Na/Zr; v- Ti/Zr ratios. Figure courtesy of A.E. Hughes.

sintered surface which differs from the bulk [40]. Figure 8 shows scanning electron micrographs taken of the sintered (1500~ 4h) surface of two 8 tool% Y20. - ZrO 2 cera- mics. The presence of the glassy phase at the external surface is clear. Tl~e amount of the glassy phase at the external surface is a function of the impurity content, types of impurities present and sintering conditions (temperature, time, gas atmosphere). During the sintering process as the impurities migrate to the external surface, some of these (e.g. alkali metal oxides) have high vapour pressures at sintering temperatures and volatilise rapidly from the external surface. Ciacchi et al [40] have reported the precipitation of Y203 at the as-sintered surface of some zirconia-yttria ceramics especially those with a high yttna content. This was attributed to the migration of a Y-rich glassy phase to the external surface followed by its decomposition on volatilisation of some components. Similar impurity migration to the external surface has also been reported by other authors [41,42]. For example, Chaim et al [41] reported migration of an iron-rich silicate phase which also contained significant amounts of yttrium, to the external surface of ZrO 2 - 4 wt% Y203 (a tetragonal phase) on annealing pre-sintered ceramic specimens. This glassy phase follow- ing migration to the external surface reacted with surface grains leading to dissolution of the tetragonal zirconia phase grains and re-precipitation of rather larger YEO3-rich cubic zirconia grains containing about 10 wt% Y203. The glassy phase was then nearly depleted of Y203, C a t and Na20 and was very much enriched in Fe203 but was reported to be present in large quantities at the external surface coveting the heat treated surface. The depletion of Na20 occurred from the glassy phase probably due to volatilisation and those of Y203 , C a t due to their dissolution in the zirconia grains [41].

Drennan and Hannink [43] have also reported the migration of the glassy phase to the external surface in the case of MgO- partially stabilised zirconia, a tough and strong

81

Figure 8. Scanning electron micrographs of the external surface of two sintered (1500~ 4h) ceramics of 8 tool% Y203 - ZrO 2 composition showing the glassy phase segregation (dark areas).

ceramic with high wear resistance properties. They have discussed the beneficial effects of this migration in terms of bulk mechanical properties. The existence of SiO2-based impurity phase at grain boundaries leads to the formation of MgzSiO 4 thus depleting stabiliser from the surface of grains resulting in precipitation of excessive amounts of monoclinic phase at the grain boundaries which like the glassy phase itself has poor mechanical properties. On addition of a small quantity of SrO (in the form of SrCO 3) to starting powders, a low melting glassy phase was formed at interfaces between grains (grain boundaries) which during the sintering of the ceramic migrated to the external surface leading to cleaning of the grain boundaries and counteracting the formation of MgESiO 4. Significant improve- ments to bulk mechanical properties were reported [43] due to control of the grain boun- dary chemistry.

One of the striking feature of this impurity migration to the external surface is the grain growth and the phase redistribution they can cause at the external surface. Thus the micro- structure and phases present at the external surface of the ceramic may be substantially different from those in the bulk. Chaim et al [41] reported that migration of the impurity phase to the external surface caused substantial grain growth of the external grains in addition to redistributing its contents. As shown in Figure 9, exaggerated grain growth has been reported in a zirconia-yttria ceramic (with 3 tool% YzO3) heated in contact with a glassy phase [44]. The constituents of the glassy phase were the same as that present in the bulk of the ceramic (equivalent of impurity contents in the starting powder). The glassy phase apparently helps redistribute the solute. The larger grains in zirconia-yttria ceramics are usually associated with the presence of a cubic phase with higher amounts of Y203 and smaller tetragonal grains with a lower concentration of Y.O 3 [45,46]. Thus the grains at the external surface (in Figure 9 and in the results reported" by Chaim et al [41]) are rich in YzO3 and are close to the cubic phase although the bulk phase in both cases has a tetragonal structure.

The segregated phases often have low conductivity, minimum activity to oxygen charge transfer reactions and therefore impede electrode reactions. Depending upon the amount of

82

Figure 9. Scanning electron micrographs showing exaggerated grain growth in a zirconia- yttria ceramic induced by heating the sintered ceramic in contact with the glassy phase at 1400~ in air). (a) Surface (--,) directly in contact with the glassy phase; and (b) micro- structure in the bulk of the ceramic. Note different magnification for (a) and (b).

segregated phases present, the effect on the electrode kinetic behaviour can be substantial. Up to an order of magnitude higher electrode resistance has been observed as a result of contamination of the external surface during sintering [47]. Since the segregation of impur- ity phases to the external surface of the electrolyte continues for a long period of time, stability of the electrode/electrolyte interface and long term stability of the zirconia- based solid electrolyte devices (e.g. fuel cells, oxygen pumps) is questionable if electrolyte mater- ials with high impurity contents are used.

Another source of impurity segregation at the external surface of zirconia ceramics is the substrate which is placed in contact with green density or sintered ceramics during sinter- ing and subsequent heat treatments. Any impurities present in the substrate material can segregate on the external surface of zirconia ceramics. This is a common problem in tech- nology. Steele [48] has reported that during the sintering of zirconia wafers, SiO2 contami- nation occurred from alumina boards (96% Al~O 3 with significant quantity of SiO 2) placed on top of zirconia wafers to keep them flat. St.eele reported more than an order of magni- tude higher electrode resistance at the electrode/electrolyte interface for contaminated elec- trolyte sheets.

3.4 Existence of interphase due to reaction/diffusion Chemical stability of electrode materials and electrode/electrolyte interfaces in the opera-

ting environment is of paramount importance in solid state electrochemical devices. A stable interface is desirable for long term stability of the device. However, if the interface corrodes as a result of oxidation or reduction of the electrode material or reactions be- tween electrode and electrolyte materials or interdiffusion of species from one phase into the other, the consequences for the interfacial electrochemical reactions may be severe. The nature of the intermediate phase and its amount and distribution will then determine the electrode kinetic properties. Several examples have been reported in the literature where formation of intermediate layers between the electrode and electrolyte either during the cell

83

operation or as a result of cell fabrication have lead to serious degradation or poor cell performance. Some of these examples are discussed below.

Many metals are used as electrode materials in zirconia-based solid electrolyte cells. The interface between a metal and the zirconia electrolyte is characterised by the presence or absence of an intermediate layer of a metal oxide which in turn depends on the thermody- namic stability of the electrode and the intermediate layer. As mentioned previously, Pd has high catalytic activity to oxygen charge transfer. However, formation of a thin inter- mediate layer of PdO at the interface between Pd and stabilised zirconia has a large influ-

4.0

3.5-

3.0--

N -r

o 2.5-

2.0- O _J

1.5-

1.0-

6o0

PdO Pd

I 1 I J 7oo 800 900 ~o0o

TEMPERATURE,~

Figure 10. Low frequency (10 Hz) impedance of a Pd/YSZ/Pd cell as a function of tem- perature in 100% oxygen.

ence on the oxygen exchange kinetics and leads to almost blocking of the oxygen reduction reaction. The formation of the PdO layer is dictated purely by thermodynamic requirements and is dependent on temperature and oxygen partial pressure. For example above 870~ in pure oxygen Pd exists as a noble metal but below this temperature Pd oxidises to PdO. Figure 10 shows electrode impedance (monitored at a low frequency) versus temperature in pure oxygen. The temperature at which a sudden decrease in the electrode impedance occurs corresponds to the thermodynamic decomposition temperature of PdO in a given oxygen partial pressure [49,50].

Another example worthy of discussion is the stability of the interfaces between perovskite based electrode systems and the zirconia electrolyte during cell fabrication and subsequent device operation in constrained thermodynamic environments. Perovskites of the formula .ABO3+ x with substitution at A-site or both A- and B-sites are commonly used as electrodes in oxygen sensors, oxygen separation membranes, steam electrolysis and solid oxide fuel cells [51-54]. Sr doped lanthanum manganite (LaMnO 3) materials are commonly used as electrodes on the air side because of their close thermal expansion match with the electro-

84

lyte, high electronic conductivity, mixed electronic/ionic conduction, ease of fabrication and the potential they offer to tailor their properties to specific needs. Although these cathode materials are reasonably stable in contact with zirconia-based electrolytes at the solid oxide fuel cell operating temperature of 1000~ on heat treatment of the (La,Sr)MnO3/doped zirconia interface above about 1250-1300~ an intermediate layer of a pyrochlore La2Zr207 and/or SrZrO 3 is formed [55-61] The conductivity of La2Zr207 or SrZrO 3 is several orders of magnitude lower than that'of the electrode. The intermediate phase not only acts as a barrier layer for oxygen transfer reaction, it also contributes to resistive voltage losses at the interface. Thus formation of an intermediate layer during cell fabrication or during cell operation over a period of time will lead to cell perfor- mance degradation. Yokokawa et al [58,59] have performed thermodynamic calculations on the phase stability of several perovskites in contact with zirconia electrolytes. The most probable reaction in the case of LaMnO 3 is:

LaMnO 3 + (2x)ZrO 2 + (3x/2)O 2 = Lal.2xMnO3 + (x)La2Zr207.

In this reaction a small amount of La20. reacts with ZrO 2 to form La2Zr207 and as the La deficiency at A-site increases the drivin~ force for the formation of the pyrochlore phase subsides. According to Yokokawa et al [58,59], LaMnO 3 sufficiently deficient (10-15% deficiency) at the A-site should be inert towards zirconia. However, as the La deficiency increases at the A-site, the Mn activity increases at the B-site. Dissolution of Mn into the zirconia matrix under these conditions has been discussed. However, solubility of Mn 3+ and Mn 4+ is relatively low in zirconia in air or pure oxygen and Milliken et al [60] have observed an extensive manganese diffusion along grain boundaries in zirconia once the interface is heated above 1200~

For SrO doped LaMnO 3, the A-site deficiency limit is affected by the La/Sr ratio and depending on the doping level, both SrZrO 3 and La2Zr.O 7 can exist in equilibrium with the 4 perovskite electrode [58,59]. The schematic of interfaclal reaction between perovskite elec- trodes and zirconia electrolytes is given in Figure 11.

A number of other electrode materials with doping at the B-site have also been proposed as potential electrodes for the oxygen reduction reaction. Sr doped LaCoO 3 is a much better electrode in an electrochemical sense than Sr doped LaMnO 3 but it reacts even more vigorously with zirconia. Ivers-Tiff6e et al [61] have reported that for (La,Sr)(Mn,Co)O 3 electrodes, the amount and ratio of the secondary phases (SrZrO 3, La2Zr207 and cobalt oxide) formed between mixtures of electrode and electrolyte powder compacts (heat treated at 1300~ was a function of La:Sr and Mn:Co ratios. Co-free materials showed a small quantity of SrZrO 3 phase and the amount of this phase increased with the SrO content in the perovskite. For Co-containing compositions, both SrZrO. and La..Zr207 phases were

0 , Z . observed in large quantities with the amount of the former increasing w~th increased Sr doping at the A-site. For (La,Sr)CoO 3, more than 50% of the electrode consisted of S r Z r O 3, La2Zr207 and cobalt oxide. Ivers-Tiff6e et al [61] have also reported that at the electrode/electrolyte interface, the overpotential losses increased dramatically with increas- ing heat treatment temperature of the interface for Co-containing electrode compositions whereas the effect of increasing heat treatment temperature (from 1150 to 1300~ on electrode overpotential losses was relatively small. Nevertheless, the electrochemical characteristics of the interface can change if reasonable quantities of one or more secondary phases are formed. The electrode reaction then will be dictated by the physical and electro- chemical nature of the intermediate layer.

85

Trends for intermediate phase formation for

AyBO 3 (A = La, B = Mn, T > 1200oC)

y > 0.9 La2Zr207 (I')

y < 0.9 La2Zr207 ($)

Mn ~ in zirconia La2Zr207 (1')

Sr substitution at A-site

Sr:La rat io ( t ) SrZrO 3 ( f )

Co substitution at B-site

Co :Mn rat io ( t ) La2Zr207 ( f ), SrZrO 3 ( t ), CoO ( t ).

Figure 11. A schematic of the LSM/zirconia interface and possible interface reactions for different conditions.

Another example of interface reactions is that of electrodes based on stabilised uranium oxide. These nonstoichiometric oxides of the general formula (U,M)O2~:x (M = Sc, Y, Dy, Pr) and having the fluorite-type structure are good electrodes for use m oxygen sensors and fuel cells [62-65]. However, some of these solid solutions when heated above about 1200~ in contact with zirconia-based electrolytes react to form intermediate phases such as (U, Zr)O2~_x and/or (U,M,Zr)O. . . These intermediate phases may have different cata-

; l : X

lytic activity~to oxygen exchange ~dnetics and once formed at the interface can alter the rate of oxygen transfer reaction. Badwal et al [66] have reported that the formation of an inter- mediate phase between (U,M)O z electrodes and zirconia-based electrolytes is a function of the type of dopant used to sta:~]lise uranium dioxide. In the case of (U, Sc)O2• x elec- trodes heated above 1200~ in contact with a zirconia-yttria electrolyte, the formation of an intermediate phase was clearly observed as shown in Figure 12. On the contrary (U,Y)Oz• x electrodes of similar composition showed much higher stability under similar heat treatment conditions [66] The (U, Sc)O_ fluorite solid solutions show very high

. " 2 X

volattlity for U in the form of UO 3 and the mecffanism for the intermediate phase formation appears to be through vapour phase transport rather than through the direct solid state reaction between electrode and electrolyte phases. In the case of (U, Sc)O z_ solid solu- tions, the loss of uranium in (U, Sc)O~_ also leads to the precipitation ~J~ Sc203 and decomposition of the fluorite phase, and~Xer cause for the loss of catalytic activity of the electrode.

86

Figure 12. (a) Scanning electron micrograph of the interface between zirconia-yttria (1) and (U, Sc)O2§ x (3) heated at 1400~ for 15 h in air showing the formation of a dense inter- mediate phase (2) at the interface. (b) X-ray line scan across the electrolyte/electrode inter- face showing diffuse region.

3.5 Interfaces between components of a composite electrode Composite electrodes are commonly used in solid state electrochemical devices based on

zirconia for several different reasons. These may include enhancing the catalytic activity of the electrode at the interface, change in the thermal or mechanical properties of the mater- ial, control of sintering and grain growth kinetics. For example Badwal et al [67,68] have reported that composite electrodes consisting of a noble metal (such as Pt) and some semi- conducting metal oxides show much superior electrode kinetic behaviour for the oxygen exchange reaction compared with either Pt or the metal oxide electrode alone. About an order of magnitude reduction in the values of the time constant and the resistance associa- ted with the electrode process have been reported for mixtures of Pt with either (U, Sc)O2§ x or CrNbO 4 materials [67,68]. This is shown in Figure 13 for mixtures of Pt and CrNbO 4 electrodes. Such electrodes gave a very good performance in potentiometric sensors ana have enabled the operating temperature to be lowered to below 400~

As mentioned earlier, zirconia doped with yttria or a similar dopant is added to the Ni fuel electrode to reduce thermal expansion mismatch between the composite electrode and the electrolyte, to reduce sintering and grain growth of Ni particles and to control electrode morphology [15,27,28,69]. It is not entirely clear if the addition of zirconia to Ni directly enhances electrode reaction rates by increasing the electrode/electrolyte/gas contact area boundary. The method of Ni/zirconia composite electrode preparation and distribution and size of both Ni and zirconia particles are important to achieve high conductivity, lower overpotential losses as well for long term stability of the electrode performance [70-72]. In

87

Electrode CrNb04 +Pt Electrode ~ CrNb04 + Pt

4.0 600o C 0.5 ~ 600Oc

3.5

o

" 2 . 5 ~

9 2.0 9 - 0 . 5

1.5

1.0 I I I I 1 I .. -1.0 I 1 I I I 1 0 20 40 60 80 1 O0 0 20 40 60 80 1 O0

Wt% Pt02 wt% Pt02

Figure 13. The effect of Pt addition on the electrode resistance and time constant of CrNbO 4 electrode. The data were recorded at 600~ PtO 2 decomposes to Pt on heat treatment at 600~

any case, the interface between Ni and zirconia plays a critical role and the existence of an oxide near the Ni particles appears to facilitate the hydrogen oxidation reaction. Contradic- tory mechanisms suggesting the dominant role of the electrode or the electrolyte have been proposed. It has also been reported that the presence of a small amount of water vapour (2-3%) in the hydrogen fuel considerably reduces overpotential losses at the Ni/zirconia interface [ 15]. The hypotheses proposed to explain this behaviour are numerous and unsatisfactory. Recently Jiang and Badwal [73] have suggested that the preferred mechanism for hydrogen oxidation is via its adsorption on the Ni surface which is facilita- ted in the presence of water vapour. However, in dry hydrogen, hydrogen adsorption takes place predominantly at the electrolyte surface and electron migration on the electrolyte surface is the rate limiting step. Other composite electrodes proposed for the fuel oxidation reaction are (i) a mixture of Ru and zirconia which is a better electrode than Ni/zirconia cermet because of the higher melting point of Ru compared with Ni (low sintering and grain growth), higher steam reforming activity and lower carbon deposition rates [74]; and (ii) zirconia-yttria-titania mixed conducting materials [75]. Mixed conducting elec- trodes prepared by the addition of a transition metal to zirconia-based electrolytes have attracted interest because of the possibility of an increase in the reaction interface between the electrode and the electrolyte and a consequent reduction in the polarisation losses. Also the stability of such an electrode/electrolyte interface would be good due to the expected low sintering and grain growth kinetics for the refractory oxide and closely matching thermal expansion coefficient (dewetting similar to that observed at the Ni/zirconia inter- face is less likely to occur) with other cell components [75-77].

88

4. INTERFACES WITHIN THE ELECTROLYTE GRAINS

In this section interfaces within the electrolyte grains have been considered. In homogen- eous single phase zirconia-based electrolyte materials, in the absence of inclusions or precipitates of another phase, the ion transport properties are determined by the level of dopant (vacancy concentration), type of dopant, interaction between vacancies and dopant cations and the temperature range over which measurements are made. In such materials ion transport takes place freely without any hindrance from internal interfaces. However, in practical electrolyte systems, the diffusing anions must negotiate a series of barriers as it transports through the zirconia-based electrolyte grains. Depending on the system under study, the types of barrier can be broadly categorised into four distinct groupings:

(a) (b) (c) (d)

Precipitation of a second phase within the grains. Compositional variations across grains. Second phase inclusions. Microdomains/ordering of cation or anion sublattice.

In each of the cases, the interfaces described produce a deviation from the idealised pic- ture of an electrolyte consisting of a homogeneous single phase material with a high con- centration of randomly distributed vacancies providing diffusion pathways.

4.1 Precipitation of a second phase within the grains In zirconia based systems it is the general case that the optimum electrolyte which com-

bines both high ionic conductivity with high mechanical strength is not a single phase material but is found in the two phase field where it is possible to produce the desired combination of phases and microstructure. Moreover, the maximum conductivity (above 800~ in zirconia-based systems occurs for a dopant concentration close to the phase boundary in the two phase field. The maximum in ionic conductivity has been reported at about 8 tool% M203 in ZrO - M O. (M= Y, Yb, Sc) and 12 -13 mo1% CaO in ZrO 2 -

2 2 CaO (3,78,79]. In all such systems, the thermodynamic requirements of phase diagrams would require some growth of precipitates of another phase. Additionally, in order to produce a homogeneous solute distribution, high temperatures well above the operating temperature of the cell are often required for processing. However, the subsequent cooling invariably results in the samples passing into two phase regions where even though the kinetics are such that cation redistribution is insignificant, very often over very localised regions of a few atomic units evidence of ordering or minor inhomogeneities can be seen. Further to this is the operating temperature of the solid electrolyte device. Although the kinetics of phase redistribution are slow, the temperature regime in which the system finds itself is usually well inside a two phase region and invariably phase separation begins albeit on an extremely fine scale due to thermodynamic constraints. Figure 14 schematically shows this situation for an idealised system in which the optimum aliovalent cation content for which maximum ionic conduction occurs is close to the phase boundary between the single and two phases regimes (for detailed phase diagrams see Yoshimura [80]).

The developing precipitates in a homogeneous matrix of electrolyte grains will definitely contain very little stabilising dopant. However, if some of the dopant remains in the deve- loping precipitate, this will lead to some conductivity which is expected to be much lower than that in the surrounding matrix An example of phase redistribution in a ZrO 2 - Y..O 3 system sintered in the single phase field and subsequently annealed in the two phase fielt~ is

89

o 0

(1) I - -

L._

(9 O.

E

3 0 0 0

2 0 0 0

1000

I I I 2 4 6

4

I 8

c

Sintering t~rrTp~r~tdr~

_ _ _Fuel cell op_er_ati_on

Technologically _ important composition range

I 1 I I 10 12 14 16

mol % stabi l iser

Figure 14. A schematic of parts of the phase diagram showing the regions of technological interest.

shown in Figure 15. The so called 'tweed structure' has developed in a specimen of 8 tool% YzO3 stabilised zirconia after prolonged heat treatment at 1000~ This structure is caused by the development of coherent precipitates of the tetragonal var- iant of zirconia within the cubic grains. There is a resultant decrease in the conductivity of this material as a consequence of this so-called "ageing" because of the precipitation of the poorly conducting tetragonal zirconia. The conductivity degradation as a function of time at 1000~ is shown in Figure 16 for 8 tool% Yz O. - ZrO 2.

A further example was reported by Steele et a~ [81] for the system CaO-ZrO z. In this case the developing microstructure as a function of aging time at specific temperatures (i.e. solid solution - > tweed structure - > tetragonal precipitates - > monoclinic precipitates) could be clearly followed, even in this very complicated system, by conductivity measure- ments. The varying departure from the idealised case of randomly distributed vacancies for this system leads to a continually deteriorating conductivity as the nature and intensity of interfaces between various phases within the grains of the electrolyte changes.

The situation is even further complicated in some zirconia-based systems by the thermo- dynamic requirements of the phase rule and the slow kinetics of phase redistribution. The necessary redistribution of dopant cations dictated by the development of a two phase region in combination with low diffusion rates of the cation species can lead to inhomogen- eous accumulations of cations at precipitate interfaces. The consequence of this will be a reduction in the concentration of vacancies available for conduction. An extreme example of this is in the case of Mg-PSZ (magnesia partially stabilised zirconia) in which the deve- lopment of coherent tetragonal precipitates gives rise to solute rich regions around these precipitates which have been shown to accumulate into ordered regions of ~5-phase [82]. Figure 17 shows the dark field electron micrograph in which the 6-phase (an ordered phase

90

Figure 15. The "tweed structure" observed in a yttria-stabilised zirconia after prolonged ageing. Electron diffraction patterns recorded from this sample show tetragonal symmetry.

I E , - 0 I C

0-18 i 8 mol% Y203 - ZrO 2

0.17 k

0.16

0.15

0.14

0.15

0.12 L . . . . 0 1000 2000

1000~

-'~176176 I | ., I ,

3000 4000 5000

Time, min

Figure 16. Conductivity degradation in 8 mol% Y203 - ZrO 2 at 1000~ as a function of time resulting from the composition being in the two phase field. Diagram courtesy of ET. Ciacchi.

91

Figure 17. Dark field electron microscope image using the ~-phase reflections observed in electron diffraction patterns of this area. The light regions show the development of ~- phase domains at the interface of the growing tetragonal precipitates. (Micrograph courte- sy of R.H.J. Hannink).

with a higher solute content) precipitates are highlighted. The evidence for this occurring in Mg-PSZ comes from the analysis of diffuse scatter in electron diffraction experiments in combination with analytical electron microscopy [83,84). Diffuse scatter in similarly aged specimens of other stabilised zirconias is often observed, i.e. Sc~O3-ZrO 2 [85] and Y O - ZrO 2 [86] and although these systems have not been fully examined as in the case o~2~g - PSZ, it is not an outlandish assumption made that a similar phenomenon is occurring. Some ordering is taking place, and with increased formation of ordered phases with a higher dopant content more vacancies become unavailable for conduction.

In the Y203 -.ZrO 2 [86-88] and Sc203 - ZrO 2 [85] systems, which have the most impor- tant properties m connection with their use in the solid oxide fuel cell, it has been shown that rapidly quenching specimens which have been fired within the cubic phase field pro- duces a metastable phase known now as the t'-phase. This phase contains a high concentra- tion of solute and the rapid quench prevents the partitioning of the material into the equi-

92

librium two phase condition. The t'- phase is characterised by the occurrence of large twins running through the grains of the material. The conductivity of this phase is high but it rapidly deteriorates with annealing at moderate temperatures [89]. The t'-phase is seen to decompose into a cubic phase with precipitation of the equilibrium tetragonal phase which is usually observed. The tetragonal phase is very low in solute and hence will have a low concentration of vacancies available for conduction. This is illustrated in Figure 18 which shows the changing microstructure of a 7.75 mol% SCzO 3 - ZrO z as a function of annealing and the corresponding deterioration in the conductivity.

A further example of phase changes which can effect the conducting properties of these materials occurs when the microstructure of the ceramic itself prevents the system attaining its equilibrium state. In such cases the system is metastable and external factors such as mechanical stress can trigger a transformation to the lower energy state with detrimental effects on the conductivity. An example of this is the doped - TZP (Tetragonal Zirconia Polycrystals) materials which have a very well defined microstructure. For the material to remain tetragonal and not transform to the monoclinic state, the grain size must not exceed a critical size (about 0.5~m). The grain size nevertheless is dependent on the solute con- tent, type and amount of impurities present, solute homogeneity and sintering conditions. If some damage is externally introduced then the transformation to the monoclinic phase can spontaneously occur with an accompanying volume expansion. The monoclinic phase is intrinsically a poor conductor and in addition an accompanying volume expansion gives rise to the formation of insulating microcracks. The sequence of events can be seen in Figure 19 in which the tetragonal grain is seen to have transformed in this case as a consequence of beam heating in the electron microscope (Figure 19).

ZrO z doped with 2-3 mol% Y20. has tetragonal structure, and is one of the most widely �9

studied potenttal electrolyte material for use in oxygen sensors and fuel cells [90,91]. For Y-TZP flexural bending strength in excess of 1 GPa has been reported [2]. Although these materials have lower conductivity than the 8 mol% YzO3 - ZrO.. composition at 1000~ at lower temperatures (below about 400 - 500~ the conductivi~ty is comparable or better [78]. However, precipitation of poorly conducting phases has also been reported in 3 tool% YzO3 - ZrO 2 ceramics [92]. This composition is in the two phase field over a wide tem- perature range [80] and on annealing sintered and rapidly cooled ceramics in the 800- 1200~ temperature range, solute redistribution leading to the presence of several variants of the tetragonal phase within the same grain and precipitation of monoclinic zirconia was observed [90]. The development of this type of microstructure leads to a substantial reduction in the intragrain conductivity as shown in Figure 20 [92].

Doped tetragonal zirconia (TZP) materials containing 2-3 tool% Y203 undergo a phase transformation to the less conducting monoclinic phase when annealed in the presence of moisture in the low temperature range of 100 - 400~ The transformation starts from the external surface and moves inwards and it is somewhat dependent on the microstructure of the ceramic. The phase transformation has serious consequences for the mechanical integr- ity of the ceramic [2]. Also Badwal and Nardella [93] have reported the formation of m- ZrO 2 on the anodic side (only) of the tetragonal zirconia electrolyte in complete cells on current passage (Figure 21). The amount of m-ZrO 2 formed is a function of the current density, time of current passage, the temperature of cell operation and the ceramic micro- structure. No m-ZrO 2 was detected at the anodic side of the electrolyte at temperatures above 500~ This behaviour can have a significant effect on the electrode performance and was attributed [93] to the existence of space charge layers near the interface leading to substantial reduction in the oxygen vacancy concentration which otherwise is responsible for the stability of the tetragonal phase.

93

T E ,..0 It2

0.32

0.30

0.28

0.26

0.24

0.22

1 0 0 0 ~ " ,

~

%

-... %

"Oo

0.20 , ' 0 1000

O0 O O OOoOoQ

000Oo0 O 00 OIO@ @

OOOoo00oO

I , I

2000 3000

Time, min

e ~ 1 7 6

~ 1 7 6 1 7 6 c o o ,

J I t

4000 5000

Figure 18. Bright field transmission electron micrographs of the sintered 7.75 tool% Sc20~ - ZrO 2 ceramic (prepared by coprecipitation) before (a) and (b) after annealing at 1000~ for 2000h in air. Also shown is the conductivity degradation for this material (freshly sin- tered) as a function of time change [3].

4.2 Compositional variations. Compositional variations can occur in zirconia based systems as a result of a number of

situations which may include: processing difficulties which produce an inhomogeneous dis-

94

Figure 19. Bright field electron micrographs showing (arrowed) the partial transformation of a Y-TZP grain.

tribution of cations; the tendency for produced metastable systems to disproportionate over a period of time at elevated temperature; and by the presence of contaminants which react preferentially with some component of the ceramic causing localised concentration gradients and creation of pseudo interfaces. These are not real interfaces in the sense that there is no precipitation of another phase. However, the system is not homogeneous with respect to solute distribution and regions of different composition have different electrical conductivities and overall ion transport through the grain involves oxygen ion exchange between these regions.

95

33000

E 22000 0

N I 11000

0 0

3 tool% Y203-Zr02 350~

A A A A A A Z $ & A A A A A

AA~ OAO O O O O O O O O Oo AAA& A .AO Or,

11 bOO 22000 3 3 b O O 44000 as000

Z' f~cm 9

Figure 20. Impedance spectra of a 3 tool% Y.O 3 - ZrO 2 ceramic: (O) as-sintered (1500~ 4h in air) and quenched in air; (zx) after annealing at 1200~ for 50 h in air. The impedance data were recorded at 350~

60 [ I=1.4mA cm 2 �9

50

4o

~ 20

lo

0 25 50 75 100 125

Time, min

Figure 21. The effect of current passage in a electrode/electrolyte/electrode cell on m- ZrO 2 formation on the anodic side of a 3 mol% Y203 - ZrO z electrolyte.

96

Improvements in processing techniques associated with the improved powder technology have gone a long way to ensure that problems of inhomogenous mixing are rare. In zirco- nia based systems this was most important since inhomogeneities in the cation distribution are difficult to remove. It is a feature of the fluorite system that the cation sublattice is most stable in comparison with the mobile anion sublattice. Extreme examples of this problem have been reported in the Sc203 - ZrO 2 system [94-96] where the kinetics of cation redi- stribution are known to be very slow probably due to similarity in the size of Zr 4+ and Sc 3§ cations. The resulting microstructure is not optimised for maximum conductivity with localised regions of high resistivity causing blocking of the diffusing anions. Badwal and Drennan [97] in studying the Y20 - Sc203 - ZrO 2 system reported wide compositional variations (Sc/Y ratios) from grain t~ grain and also within the same grain. An example is shown in Figure 22. This type of inhomogeneous distribution produces additional interfaces for ionic transfer and materials with less than optimum ionic conductivity.

Figure 22. Scanning electron micrograph of 50 wt% (4.7 mo1% Sc203 - ZrO 2) + 50 wt% A1203 and electron probe micro analysis results showing inhomogeneous distribution of Sc within a zirconia grain. Sc/Zr = 7.0(1); 3.0(2); 3.0(3); 6.0(4); 5.5(5); 6.0(6); and 8.0(7).

In a number of systems which are of technological importance in solid electrolyte devices the optimum compositions for maximum ionic conductivity, as reported earlier, are usually close to the phase boundary between single and two phase regions. Therefore, localised variation in composition may be brought about by the system having to be operated in

97

temperature regimes in which the equilibrium situation is a two phase mix. An excellent example of this is described by Chaim et al [98] in which disproportionation can occur and even if this is over atomically small regions the process removes available vacancies from the conducting process.

A further possibility first outlined by Heyne [99] and later discussed by Steele and Buffer [100] is the formation of space charge regions developing at the surface or the grain boundary of the ceramic. These form as a result of the large difference in mobility between the cation and anion species and the resulting difficulty the system has in rapidly attaining the equilibrium situation. Local variations in the concentration of vacancies will effectively produce barriers to the diffusing anion.

Contamination of the starting powders by glassy impurities can produce situations where the microstructure can be dramatically altered. In general, impurities are found to accumu- late at the grain boundary of the ceramic and will in general react preferentially with one of the components of the system. Examples of this can be seen in the yttria-zirconia [101- 105] system where it has been shown that accumulations of secondary grain boun- dary phase rich in Si t . aid in the redistribution of the solute in the system giving rise to inhomogeneous grain g~rowth and large compositional gradients. An example is shown in Figure 23. These large variations in microstructure once again interfere with the continuity of conducting pathways, producing additional interfaces which result in inferior materials for use in electrochemical devices.

4.3 Second phase inclusions. Generally speaking, any restriction of the pathway of the diffusing anions through the

ionic conductor will result in the degradation of the conductivity. Second phase inclusions of a non-conducting species will have two detrimental effects. Firstly, the second phase effectively dilutes the volume of conducting matrix and secondly, the insulating phase constricts the direct pathway for the diffusing ions. However, subtleties can arise. It has been reported that the addition of alumina to yttria stabilised zirconia can have a beneficial effect on the total conductivity of the ceramic at low temperatures [106-108]. The proposed mechanism involves the gettering of the contaminants such as silica so that the addition effectively cleans the grain boundary regions opening up the pathways for conduction (see section 5.4). Nevertheless the effect of insulating phase addition is normally detrimental to the intragrain conductivity. It should be noted that the dilution of the conducting species has still occurred and more restricted pathways have been introduced by the addition. As a result of second phase addition, the net intragrain conductivity (ignoring grain boundary resistivity effects), usually is lower than can be calculated taking into account the volume fraction of the insulating phase due to isolation of some grains of the conducting phase from the conducting pathways in addition to constriction of current lines. This is shown in Figure 24 [109].

One of the other features of second phase addition such as alumina is its effect on the grain size. The addition of alumina to ZrO 2 - Y203 leads to a significant reduction in the grain size with a consequent increase in the interfaces between grains (grain boundary surface area. As shown in Figure 25, the effect is clearly noticeable in the case of fully stabilised ZrO 2 - Y203 [110]. In general, if defect sizes are minimised in combination with a reduction in the grain size, an increase in the mechanical strength is observed.

4.4 Microdomain formation and ordering A subtle effect is observed in stabilised zirconia systems which manifests itself in the

appearance of diffuse scatter in diffraction experiments. Although these effects have been

98

Figure 23. Bright field micrographs showing the inhomogeneous grain growth in a speci- men of Y-TZE The yttria content of the large grain was observed to be close to 9 mol%, approximately three times the yttria content of the surrounding grains.

observed and studied for over thirty years, there still exists doubt as to the exact cause of the effect. All workers agree that some ordering process is taking place and with the order- ing comes the inevitable reduction in ionic conductivity. For the system CaO-ZrO 2 (within the composition region Ca Zr. O. where x lies in the range 0.1-0.2) Allpress and Rossell [111] provided very ~lea~Xev~lence that the observed diffuse scattering was as a result of the formation of microdomains of the order of 3nm in diameter within the grains of the material. The microdomains were shown to have clear structural integrity corre- sponding to the ordering observed in the compound which occurs in the CaO-ZrO 2 system, CaZraO 9. In the case of Y203-ZrO. the situation is less clear. A microdomain model has been suggested [ 112,113] but the difficulty with this model is that diffuse scatter can be observed over a wide range of stoichiometry but no compound exists to which the micro- domain model can be applied. More recently alternative suggestions [114, 115] invoke the concept of the modulated structure where intermediate compositions between the end members of the solid solution range show a regular commensurable modulation in very

99

60000

50000

E 4OO0O U

>" 30000

.,I--

| 20000

I0000

4500C

/

/ j.J"

j o ~

0 I I I I I 0 I0 20 30 4 0 50 6 0

wt. % A1203

Figure 24. The effect of alumina addition on the resistivity of 5.9 mol% 5C203 - ZrO 2 ceramics. The data were recorded at 450~ [30].

specific crystallographic orientations. Irrespective of the mechanism, an ordering process is clearly taking place which is limiting the availability of anions for the conduction proces- ses.

Figure 25. The effect of alumina addition on the grain size distribution in 8 mol% Y203 - ZrO z ceramic sintered at 1500~ for 4h. (a) No alumina addition and (b) with 10 wt% alumina addition.

100

5. INTERFACES BETWEEN GRAINS (GRAIN BOUNDARIES)

In most applications polycrystaUine materials are invariably used. The polycrystalline nature of ceramic materials means that the interfacial region between the grains have a profound influence on the physical properties of the material. In terms of the electrical properties, the boundary region provides a barrier to the transport of oxygen ions in zirco- nia - based solid electrolytes. The grain boundary interfaces can be divided into a number of different categories similar to the interfaces within the grains discussed in the previous section. Four major types of grain boundaries encountered are"

(a) Phase free (or clean) boundaries. (b) Intermediate phase formation at grain boundaries from the matrix components. (c) Intermediate phases of the impurity type. (d) Inclusions.

5.1 Phase free boundaries In terms of the electrical properties, the boundaries which are free from any secondary

phases or segregation of impurities should be the ones with minimum hindrance to the ionic conducting species. However, this would require complete matching of the atomic inter- faces which in practice is not observed. Phenomena such as lattice mismatching, disloca- tions and residual stresses observed commonly in metals are also observed in ceramic systems although their analysis and characterisation is still in the early stages of study due to difficulties in preparing representative samples and the added crystallographic complexi- ties. However, imperfections, lattice mismatch, and dislocations are all observed in cera- mic systems and no doubt will receive more attention as the structures are better under- stood. Nevertheless, these are some of the interfaces present in clean materials which will influence the ionic transport properties across grain boundaries. Moreover, the presence of pores at grain boundaries or imperfect contacts between grains will cause constriction of current pathways leading to a decrease in the overall conductivity. Also enhanced segrega- tion of bulk phase matrix components near the grain boundary region can alter the nature of space charge layers. In zirconia-based systems, it is extremely rare to come across grain boundaries free of secondary phases and/or segregated impurities. Apart from single grains or crystals with low angle grain boundaries, there is no demonstrated case of so- called clean grain boundaries.

5.2 Intermediate Phase Formation from the Matrix In the case of complex zirconia - based systems which are maintained at temperatures in

which phase changes are taking place, the grain boundary region can provide the nucleation site for the development of secondary phases. The classic example of this is the MgO-PSZ system which when maintained below 1400~ develops a decomposition product around the perimeter of each grain. The decomposition reaction is simply the MgO-PSZ decomposing into zirconia (monoclinic form ) and free MgO. This reaction nucleates at the boundaries and develops towards the centre of the grains as can be clearly seen in Figure 26 which shows the grains outlined by the decomposition region. Obviously the effect on the conductivity of the formation of this insulating layer is dramatic. With increase in the amount of m-ZrO 2 at grain boundaries, a substantial increase in the grain boundary resis- tance is observed as shown in Figure 27 (78).

Attempts to reduce the amount of decomposition with time have been successful by con- trolling the chemistry of the grain boundary phase. The addition of SrO to the system has

101

Figure 26. Optical micrograph of the decomposition region around the grains of over aged Mg-PSZ. The white regions consist of monoclinic zirconia and free MgO.

3 0 0 0 0 0 . . . . . . . . . . . . .

E ~ o o e o o e o o e e o c

C~ lOOk �9 , , . " �9

" k l , , 150000 ... " " ee e

I .o~OOk 1Ok

~i~oo ~ l k I i

0 150000 300000 450000

IB o

Mg-PSZ 450~

�9 l k

�9 %e% e ~ %..

~,.

600'000 z s o o o o

Z', Q c m

Figure 27. Impedance spectra (recorded at 450~ in air) for 3.4 wt% MgO - ZrO 2 (MgO- PSZ) before ((3) and after (o) annealing at the modest ageing temperature of 1000~ for l l0h. A substantial increase in the low frequency grain boundary arc appears to have occurred.

102

been shown by Drennan and Hannink [43] to favourably alter the grain boundary decompo- sition with subsequent improvements to changes to the mechanical properties of the cera- mic. Although MgO stabilised zirconia are not suitable materials for most applications requiting high ionic conduction apart from their possible use as immersion sensors for molten metals, measurements of conductivity have been made [116] and the influence of SrO additions examined. These appear-to support the idea that the addition effectively reduces grain boundary decomposition. Although this is an extreme case of intermediate phase formation, it serves as an example of the situation which may occur when nucleation points such as grain boundaries and pores provide the regions where solid state reactions can rearrange the local chemistry of these metastable systems.

5.3 Intermediate phases of the impurity type Important advances have been made in the preparation of zirconia powders. Contamina-

tion levels are controlled down to the 10's of ppm and additives deliberately added to aid in sintering are controlled precisely. In general most contaminants or sintering additives accumulate in the grain boundary region. The control of the location and chemistry of this grain boundary phase becomes paramount in optimising the ionic conduction of the cera- mic. Two competing phenomenon take place. Homogeneous, high density, flaw free ceramics are necessary properties of any electrolyte since with present device designs the electrolyte must maintain integrity whilst being thin enough to have minimum resistance. A means of obtaining these properties is to use sintering aids which in tum produce liquid phases to facilitate homogeneous grain growth and assist in the elimination of porosity. The difficulty arises when the liquid phase formed remains along the grain boundary providing a most effective insulating barrier. In such cases the beneficial effects of the presence of ,the liquid phase on the densification processes are negated by the insulating grain boundary phase. Examples of the type of liquid phase location are shown in Figure 28. In Figure 28 (a), the grain boundary phase can be seen outlining the grains of yttria stabilised zirconia, totally wetting any interface region. Altematively, in (b) the triple point regions show accumulations whilst in (c) the total dewetting of the boundary results in isolated accumula- tions along the boundary. The type of grain boundary phase formed is a function of the sintering temperature, gas atmosphere and chemical nature of impurities [39].

The electrical behaviour of these different grain boundary microstructures, as expected, is quite different. The grain boundary phases which wet the grain surfaces well have the most detrimental effect on the conductivity across grain boundaries. Figure 29 shows examples of different contributions to the grain boundary resistivity for different grain boundary microstructures.

Attempts have been made to influence the ionic conductivity of the grain boundary phase so that minimum disruption to the conduction is obtained. For example Verkerk et al [117] added Bi203 to yttria stabilised zirconia to aid in densification but also to influence the con- ducting properties of the boundary since bismuth oxides are known to have high ionic conductivity at elevated temperature. The addition succeeded in assisting densification, however, the conductivity was seen to decrease as a result of unwanted interfacial reac- tions. So far attempts to enhance grain boundary conduction by adding Bi203 have failed.

The grain boundary phases are quite mobile and their location can be altered by post- sinter heat treatments. Some examples are discussed below.

Altering the cooling rate of the fired specimen of yttria stabilised zirconia has been found by Badwal and Drennan [ 118, 119] to affect the grain boundary resistivity. They suggested that the effects observed were a result of the changing nature of the wetting behaviour of

103

Figure 28. A series of micrographs showing the various distributions of grain boundary phase in yttria zirconia. (a) Distribution evenly along the grains; (b) in the triple point regions; (c) in pockets along the boundary.

the grain boundary phase. At elevated temperatures the liquid phase was observed to dewet the ceramic boundary and rapid cooling effectively froze-in this structure. Alternatively slow cooling allowed the grain boundary phase to rewet the boundary region consequently reducing the total conductivity.

External pressure [ 120] on the sintered ceramic at elevated temperatures (in the vicinity of 1200~ can also force the grain boundary phase to move away from the conduction paths in the direction of pressing and move into parallel paths thus occupying regions it would not under normal conditions. In effect the application of anisotropic pressure to force the liquid phase into regions parallel to the applied force has lead to increase in the ionic conductivity along the pressure direction. The grain boundary resistivity was also seen to decrease as a function of increased applied stress. Similarly heat treatments of the sintered ceramic lead to migration of the grain boundary phase to the external surface thus decreas- ing their impact at grain boundaries [47].

This shows that the grain boundary microstructure of ionic conducting materials can be engineered and the nature of interfaces changed to optimise conditions.

5.4 Inclusions Since the present generation of commercially supplied powder materials are of a high

quality we will ignore the cases of accidental contamination by foreign items. It is suffi- cient to say that any non conducting inclusion present along grain boundaries will in gen- eral impede ionic transport. However, the exception has been found to occur with the addi- tion of A1203 to yttria stabilised zirconia. The effect here is subtle. The addition of the nonconducting species does reduce the ionic conductivity through the grains as would be

104

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o o

o o

o o

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0 0 0 0 0 0 0 0 0 0 0 0 0

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0 150000 ,500000 450000 600000 750000

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Figure 29. Impedance spectra showing different contribution of the grain boundary resistiv- ity (arc on the right hand side) for different grain boundary microstructure. (a) Relatively clean grain boundaries; (b) grain boundary microstructure with grain boundary phase present at a number of interfaces between grains; and (c) extremely wetting grain boundary phase microstructure. The data were recorded at 350~

expected but for certain conditions, the total conductivity of the system is increased because the alumina changes the grain boundary conductivity. Figure 30 shows impedance spectra of two samples of Y-TZP with differing amounts of added alumina. A clear reduction in the total conductivity of the system is seen to occur for the specimen containing 10 wt% AlzO 3. The explanation for this is that the wetting of the grain boundary phase at the AlzO3/yttria-zirconia interface is favoured over the yttria-zirconia/yttria-zirconia interfaces (Figure 31). Evidence for this comes from detailed microstructural studies combined with electrical measurements [106,121]. The actual amount of A1203 will depend on the distribu- tion of the alumina and the relative size of the Y-TZP and the alumina particles. As the concentration of A1203 increases above a critical amount then the blocking nature of the alumina becomes dominant and the pathways for conduction are greatly reduced.

105

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, I

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o" ~176 ; oeo.

"o coo

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Figure 30. Impedance diagrams (recorded at 400~ for a Y-TZP ceramic (a) without and (b) with the addition of 10 wt% alumina during powder processing. A substantial decrease in the grain boundary resistivity and a small increase in the intragrain resistivity has oc- curred.

Some recent work [122] has shown that inclusions, pores and even microcracking can be investigated using complex impedance measurements and information on how these particu- lar phenomena interfere with the ionic conductivity is becoming available.

Previously we have cited the example of how the use of external pressure during sintering has been used to influence the location of the grain boundary phase. A further interesting case is where deliberately added second phase inclusions provide another means of influ- encing the electrical properties of the material when external pressure is applied during sintering. This has been shown by Drennan et al [123] for Y-TZP/A1203 composites where the application of external pressure by hot pressing gives rise to a microstructure in which the A1203 particles appear to accumulate in rafts perpendicular to the pressing direction. As a consequence of the formation of this anisotropic microstructure impedance measurements show a 25-30% increase in grain boundary resistance measured parallel to the pressing direction and a 9% increase in the bulk conductivity.

6. CONCLUSIONS

We have described a number of the interfacial phenomena which occur in systems which incorporate zirconia based electrolytes and have attempted to categorise them according to the specific region of activity in solid state electrochemical cells. The understanding of the

106

Figure 31. Micrographs showing the interfacial region between Y-TZP grains and alumina particles where accumulations of a secondary grain boundary phase has occurred.

processes that occur at various interfaces and their effect on the electrical and electroche- mical properties of the materials becomes an important part of the development of success- ful devices especially those that must continue to perform over long periods of time. By obtaining an understanding of the processes it is possible to engineer the various micro structures involved so that optimum properties can be achieved. The study of interfacial properties will continue to be a most important field as we move into the age of the fuel cell and associated electrochemical de~,ices.

7. ACKNOWLEDGEMENTS

Authors wish to thank Nathasha Rockelman and Kylie Crane for assistance with prepara- tion of micrographs and Dr M.J. Bannister for reviewing this manuscript.

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