Structural evolution of Sc-containing zirconia electrolytes

SOLID STATE I3 SEWER

Solid Stale lonics

IONIG

100 (1997) 2X9-295

Structural evolution of SC-containing zirconia electrolytes F. Tietza9*, W. Fischer”, Th. Hauberh, G. Mariotto’ “Forschungszentnrm Jiilich G&H, Institut fiir Werkstofle dcr Energietechnik (IWE), D-5242.5 Jiilich, Germany hForschungszentrun~ Jiilirh GmhH. lnstitut fiir ~nergieverfahrm.v!echnik (IEV). D-52425 Jiillich. Germajzy ‘Istiiuto Nazionair per la Fisica dclla Mareria (INFM) and Dipartimmto di Fisica, C’nivrrsitci di Trento, I-38050 Pow.

Italy

Rcccived 13 March 1997; accepted 4 June IYY7

Abstract A powder of nominal composition (ZrOz)U~886(Sc20~),, ,,,,(AIZO ,),,,., was synthesized by spray drying with of testing the performance of solid oxide fuel cells containing scandia-stabilked rirconia (SC!%) as electrolyte. resulting from calcination and sintering at different temperatures were investigated by XRD, impedance spectroscopy. At sintering temperatures of l200-1400°C nearly equal amounts of cubic and rhombohedral dctcctcd, whereas heat treatment higher than 1500°C Icd to a cubic single-phase material. A preliminary reaction of phase formation is proposed with respect to the various results depending on heat treatment. Keywords:

Solid oxide fuel cell: Zirconia-scandia

electrolyte;

X-ray diffraction:

1. Introduction Planar solid oxide fuel cells (SOFCs), actually operating at 900-lOOO”C, have shown serious drawbacks such as high degradation rates, espcciahy of anode materials [l], the use of high-cost alloys for interconnectors [2]: and incrcascd corrosion of components [3]. Therefore an operating temperature level around 800°C is envisaged to diminish thcsc problems as a compromise between lowcrcd power density and better long-term stability. To maintain power density similar to that at 900-1000°C other ceramic materials arc needed with superior electrochemical properties. Among these the solid electrolyte plays a crucial role for a planar electrolytcsupported SOFC and has to show higher ionic *Corresponding author: Tel.: : 4Y 2461 2461 6lS700: e-mail: f.tietz~h-juelich.de

61.5007;

0167-2738/97/$17.(X) C I997 Elsevier Science PII SO167-2738(97)00356-l

fax:

+49

H.V. All rights reserved.

Impedance

spectroscopy;

the purpose The phases and Raman ScSZ were mechanism

Raman spectroscopy

conductivity than 8 mol% yttria-stabilized zirconia (8YSZ). Three materials have been identified as potential alternatives during the last few years: (i) Sc-stabihzed zirconia 14-71, (ii) alkaline-earth- or rare-earthcontaining cerias [8,9] and (iii) lanthanum-gallatebased perovskitcs [lo]. From the manufacturer’s point of view, SC-substituted zirconia is the most interesting candidate because it has chemical and thermomechanical properties very similar to 8 mol% yttria-stabilized zirconia (SYSZ) and its interfacial interaction with electrode materials may be rcgardcd as already investigated [I II. The results of Yamamoto and co-workers [5,12,13] have indicated that zirconia with IO-12 mol% Sc,O, shows no degradation after exposure for several thousands of hours at 1000°C. Furthermore, small additions of AI,O, suppress the cubic # rhombohcdral phase transition usually occurring at 60%660°C [6,14, I.51 for

290

F. Tiea et 01. I Solid Stow Ionics 100 (1997) 289-295

electrolytes with such concentrations of stabilizer. Although different compositions are electrically characterized [4,6,7,14] and first cell tests showed promising power density even at 800°C [7,14.16], the mechanism of phase stabilization by the Al,O, has not yet been clarified. To also test this material in an anode-supported SOFC [ 171, a powder with the composition (Z~,),.,,,(Sc,0,),.,0,,(A1203)0.0, has been sytlthcsized [ 211. Additionally, in order to understand the influence of the alumina doping several experimental techniques have been used to investigate the initial phase evolution of this material from a rather amorphous state to the fully SC-stabilized zirconia. Here WC report first results obtained by X-ray diffraction (XRD)? Raman and impedance spectroscopy.

2. Experimental Starting materials for powder preparation were ZrOCl, . 8Hz0, Al(NO,), .9H,O and Sc,O.,, (all purchased from Merck). ZrOCI, . 8Hz0 was first dissolved in distilled water. After adding a concentralcd ammonia solution the precipitate was filtered and washed with distilled water until no chloride ions could be detected in the elucnt using silver nitrate solution. Then the solid was dissolved in concentrated nitric acid and the solution heated and stirred until a clear solution was obtained. The scandia and the aluminium nitrate were dissolved in hot concentrated nitric acid and distilled water, rcspcctively. The three solutions were combined for a stock solution, which was spray-dried at 172°C. The resulting white powder particles had the typical shape of hollow spheres with particlc diameters of 2- IS Frn (see Fig. 1). The raw powder was calcined at 900°C for 5 h resulting in a weight loss of approximately 50%. The calcined powder was uniaxially pressed at 100 MPa into pellets of 25 mm in diameter and 1 mm in thickness. The specimens were then air-sintcrcd at several temperatures between loo0 and 1650°C typically for 5 h and slowly cooled down in the furnace. Longer annealing or rapid cooling did not show any signilicant influence on phase formation. X-ray powder diffraction (XRD) measurements

Fig. I. Typical powder particles of spray-dried ScSZ, magnilication: (a) IOWX

and (h) 5000X

The hollow spheres consist of

very thin scales stuck together like puff paste.

were performed with a Siemens D5000 diffractomcter in Bragg-Brentano geometry in the range 20 = 20”-70” (step width 0.01”). Cu Ko radiation and a diffracted beam monochromator were applied in all casts. Two types of XRD mcasurcmcnts wet-c carried out: (i) measurements at room temperature after heat treatment and (ii) in-situ measurements at clcvatcd tcmpcraturcs. For the high-temperature insitu investigations a Buehler HT attachment was used. The sample powder was scdimcntcd onto a d.c.-heated PC tape (surrounded by a radiation screen for temperature homogenization), and a Pt- 1OPtRh thermocouple was spot-welded at the rear side for measurement and control of the temperature. The temperature was increased with a slope of 1 K s-l; before starting an XRD in-situ scan the temperature

F. Tiptz. et al. I Solid Store tonics

100 (IW7)

291

289-29.5

was kept constant for 3 min. Platinum reflections of the supporting heat-tape were used as the internal standard. The displacement error due to the thermal expansion of the tape during heating was corrected for precise lattice parameter determination by applying the Rietveld retincment technique. Raman experiments were carried out at room temperature in air using standard equipment consisting of an argon ion laser (Spectra Physics, Mod. 164), a double-monochromator (Jobin Yvon, Mod. Ramanor HG2-S), equipped with holographic gratings (2000 grooves mm-‘) and a cooled photon counting system, as described in [IS]. Impedance spectra were recorded between 300 and 1000°C in the frequency range of 10 ‘- IO6 Hz using a Schlumberger Solatron 1260 frequency rcspouse analyzer. Platinum thin films were sputtered onto both pellet faces as electrodes. The data obtained were fitted with an equivalent circuit program package [19]. 48

50

52

54

201” 3. Results and discussion 3.1. XRD In dependence of the sintering temperature a cubic phase (c-ZrO,) or a cubic phase coexisting with a rhombohedral phase (c-ZrOz + fi-phase [ IS]) has been observed in XRD measurements at room temperature. Fig. 2 illustrates this situation for the example of the Bragg angle window 28 = 47-54”, where the cubic 220 and the rhombohedral 220 and 220 reflections [15] are found. After sintering at about 1000°C and at 1600°C only the 220 reflection of the cubic lattice was observed. At intermediate sintering temperatures the reflections of both C-ZIG, and P-phase occur with different relative intensity. Sintering at low temperatures, i.e. 900-l 100°C. leads to broad peaks due to small particle size and/or highly disordered crystallites. The deviation from the cubic lattice seems to be small with respect to the rhombohedral peak splitting. After heat treatments between 1200 and 14OO”C, both c-ZrO, and p-phase are formed in comparable amounts during cooling. Sintering at about 1400°C seems to give a maximum amount of P-phase. The occurrence of two phases after cooling down to room

Fig. 2. Part of the XRD pattern of S&L between 47 and 54” (2@) with cubic 220 reflection at 50.7” and 220, 220 reflectiom of the rhombohedral p-phase [15] at 50.1” and 51.3”. respectively.

temperature is accompanied by a decrease of the cubic lattice parameter a (Fig. 3). This might be an indication of a diffusional redistribution in the crystalline parts as well as of an ongoing crystallization. The similarity of the lattice parameters of the spray-dried and co-precipitated powder with the same composition indicates that the decrease of CI with increasing sintcring temperature is not a property of the spray-dried powder alone. The situation is different again at higher sintering temperatures. Even at 1500°C only a small amount of P-phase is present and after sintering at 1600°C the sample consists of a single cubic phase. Its lattice parameter u =508. I l(9) pm is in good agrcemcnt with literature data 15,161. A temporary swelling was observed for spraydried powders of this composition and of yttriacontaining zirconias during the recording of sintering curves in a dilatomctcr 1201. Therefore, high-temperature XRD measurements were performed using the raw powder after calcination to study possible phase transformations in-situ. Here attention was

F. Tiet? et 01. I Solid Slate Ionirs IO (1997) 289-295

due to the thermal expansion of the lattice without any indication of superposed structural reorganization or transformation. Slight fluctuations above 1400°C arc caused by sample height changes due to Pt tape expansion and plastic deformation. The linear thermal expansion coefficient was calculated by linear regression as u = 11.2X 10-e K ’ between 25 and 1000°C and (Y= 12.9 X 10h K ’ between 900 and 1500°C. The former value is in excellent agreement with the value obtained by dilatometry ((~~~_,,,~~~c= 10.9x 10” K--l 120,211). I

3.2. Raman spectroscopy

I

I

1600

1200

800

sintering temperature / “C Fig. 3. Lattice parameter a at room temperature of spray-dried cubic ScSZ (W) vs. sintcring temperature. Additionally, the data of a co-precipitated powder of the same composition arc shown (0). The large error bars arc due to broad or overlapping peaks because of inhomogeneities and the occurrence of the P-phase in the sintering tcmperamre regimes of KM-1000°C and 12001400°C.

Raman spectra were recorded on samples sintered at 1000, 1200 and 1600°C (Fig. 5). These samples arc denoted in the following as ScSZlO, ScSZ12, and ScSZ16, respectively. Spectra of the fresh surface as well as the fractured surface were recorded for all specimens. No difference was observed in the two cases. apart from a small relative intensity change of the band at about 260 cm-’ for ScSZ12 (see Fig. shows

paid to the XRD peak shapes and the temperature dependence of lattice parameters during heating. In contrast to the room temperature XRD diagrams the in-situ experiments from 90%1500°C show the cubic phase only. The temperature dependence of the cubic lattice parameter LI from 25 to 1500°C is given in Fig. 4. From 900 to 1500°C it increases smoothly

a complex

0

I&o

A0

temperature

250

500

(a:

,\-750

I’

wavenumber / cm-’

II

I “C

Fig. 4. Temperature dependence of the cubic lattice parameter determined from high temperature in-situ measurements.

with

fractured surface

,

40

structure

surface

-

;

band

a

Fig. 5. Raman spectra of (ZrO,),,,,,(Sc,O,),, ,,,,(N20,),,, recorded from the fresh surface (dotted line) and fractured surface (continuous line) of specimens sintered at I(XX)“C (a), 1200°C (b), and I6WC (c), respectively.

I:. Tietz et al. I Solid State IonicsIcx)(1997) 289-295

several features: the bands at 580-630 cm --’ arc composed of at least three components with an intensity maximum at 605 cm-‘. Apart from the slightly differently shaped modes in this part of the spectrum and the very sharp peak at 150 cm-’ the spectrum looks like that of 8ScSZ reported by Yamamoto 1131. Also the peak at 260 cm ‘, usually referred to as a fingerprint of t-ZrO, [13,22], is present with remarkable intensity. Therefore the Raman measurements indicate a composite character of the powder treated at 1000°C which consists at least of t-ZrO, and c-ZrO,. However, due to the broad XRD peaks as well as the complex Raman band structure the occurrence of the P-phase with small lattice distortions (with respect to the cubic lattice) cannot be excluded. After sintering at 12OO”C, the shoulder at 630 cm-’ disappears and the modes around 600 cm-’ arc shifted to lower wavcnumbers peaking at 580 cm ’ . Also the band at 260 cm-’ becomes much less intense and the mode at 1.50 cm-’ is broadened. This spectrum appears very similar to that of 1 IScSZ [ 131. It therefore seems that the triplet structure observed in our ScSZl2 sample in the region of 600 cm ’ is characteristic of the P-phase. ScSZl6 shows a Raman spectrum in which a single unstructured band appears at 615 cm -‘. This broad peak is a characteristic feature of stabilized c-ZrO, and turns out to be comparable with published spectra 113,221. In addition, the bands centered at about 155, 260, and 480 cm ’ indicate the presence of a tetragonal phase component. Since no peak splitting in the XRD pattern of this sample was observed, these Raman peaks are attributed to the so-called t”-phase [23]. This assumption seems reasonable with respect to the phase diagram prcsented in 1241, because the chosen composition is very close to the t-P phase boundary. 3.3. Impedance

spectroscopy

Significant changes of the electrical properties were also observed for differently sintered samples as can be seen in Fig. 6. Since no attempt was made to obtain high-density samples during this study, the data were corrected for porosity 1251. Although after this normalization the data are in good agreement with earlier investigations [6,7,16], the temperature

293

T/T 1000

800

400

600

O-l-

% s1

-2-

'g

-3-

be 2 -4-5-

SASZ, 1000

0

SASZ,l200"C SASZ.16OO"C SASZ (gh),1600°C

0 -6-

“C

A

. 08

1.0

1.2

b

~

1.4

1.6

1.8

lo3 T-’ / K-’ Fig. 6. Ionic conductivity temperatures

as indicated

of ScSZ samples sintered at different in the legend,

(gb=grain

boundary

conductivity).

dependence of the ionic conductivity is more interesting here. The datapoints in Fig. 6 interconnected by lines indicate temperature intervals from which activation energies of the ionic conductivity were calculated according to the equation log (T,,,,,, = log A - iE,I(RT). After sintering at lOOO”C, the conductivity data follow a straight line in a (T vs. 1 lT plot from which an activation energy of 1.17 eV can be derived. The conductivity of the other two samples, ScSZl2 and ScSZ16, resulted in curved lines with different shapes as a function of temperature. Whereas ScSZ12 exhibits two distinct temperature regions with a wide cubic # rhombohedral transition region [6] and an activation energy of 0.83 eV above 7OO”C, the other sample has a continuously changing activation energy, typical of cubic ScSZ 14-71. The activation energies are 1.35 eV and 0.72 CV for the temperature regions 300-500°C and 700- 1OOO”C, respectively. For ScSZ16 the conductivity of grain boundaries has an activation energy (1.28 eV> similar to the total conductivity (see Fig. 6). Since below 600°C the is strongly conductivity influenced by grain boundaries it can be concluded that all the specimens investigated have grain boundaries differing in con-

204

F. Ti~fz et al. / Solid Smte Ionics 100 (1997) 289-295

stitution and composition. Whereas ScSZlO has a conductivity linearly increasing with reciprocal temperature over the whole temperature region investigated, the other samples show a predominant bulk response at high temperatures. This difference may be explained by the fact that the major part of ScSZlO is highly disordered or amorphous and forms continuous paths within the whole sample as found in a transmission electron microscopy (TEM) study (Fig. 7). Due to the ionic redistribution and crystallization proccsscs at higher sintcring temperatures. the P-phase is predominantly formed in ScSZ12 at the expense of the intergranular phase and is well pronounced as shown by the steep conductivity change between 450 and 600 “C. Above this temperature region the activation energy is still 15% higher than the value for pure cubic ScSZ. This is attributed to the further co-existence of amorphous, i.e. low-conducting, fractions within the crystalline bulk 121) and it should be emphasized here that the shape of the graph of ScSZ12 is not completely

comparable with that of pure P-phase [6,14,15], because the cubic # rhombohedral transition appears in a wider temperature region. In the case of S&Z16 it can be assumed that the grain boundary phase consists of a glassy phase [26,27] because of comparable activation energies [4,26,28,29] and the fact that SiO, impurities have been segregated on the surface of pores detected by clement mapping in scanning electron microscopy (SEM) investigations. Although a cationic redistribution with increasing sintering temperature is evident from all the cxpcrimental techniques applied, the appearance of the different phases at room temperature seems to arise mainly from slightly different SC contents in the crystalline phases. which is revealed especially by the crystallographic phases detected with XRD and Raman spectroscopy. Unfortunately, the investigations so far give no definite information about the role of the added alumina or other impurities. First analytical measurements by SEM/EDX have shown reasonably that the Si distribution becomes inhomogeneous with increasing sintering tcmpcrature. Especially after high annealing temperatures (> 14OO”C), Si-rich layers on pore surfaces are formed. In the case of Al no such systematic segregation at microdomains was observed. Although some Al-rich pm-scale regions were also detected here, they seem to be distributed randomly in the sample. Most of the AI,O:, doping is dissolved in the crystal structure of ScSZ. Further analytical work is necessary to elucidate the effect of additives and impurities.

4. Conclusions

Fig. 7. TEM image of a spray-dried ScSZ grain after sintering at 1000°C. The dark particles are small crystallites surrounded by an amorphous matix (gray).

ScSZ powder with I mol% Al,O, was synthesized homogeneously by spray drying. Although the alumina doping should stabilize the cubic zirconia phase down to room temperature, this stabilization is effective only for specimens sintered at the highest temperatures ( > 1500°C). At lower sintering temperatures the rhombohedral P-phase also appears in significant amounts. This may be related to diffusional cationic redistributions during annealing. Additionally, the segregation of Si-rich layers on grain surfaces was observed occurring at similar sintering temperatures as are necessary [or a successful cubic

F. Tietz et al. I Solid State Ionics 100 (1997) 289-295

phase stabilization. Although the experimental results are still scarce at this stage of the investigations, nevertheless this Si-segregation may have an influence on phase formation and stabilization at high temperatures.

Acknowledgements The authors thank Prof. T. Schober for TEM investigations, Dr. E. Wallura for SEM/EDX analysis, W. Jungen for powder preparation, P. Lersch for XRD measurements, and N. Steffens for experimental assistance.

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