Stabilized bismuth oxide

Abstract

Disclosed is a stable bismuth oxide composition having excellent oxygen ion conductivity comprising: from 50 to 90 mole % Bi.sub.2 O.sub.3, from 10 to 40 mole % of a rare earth oxide, such as yttria; and from 0.1 to 10 mole % of an oxide compound, such as ZrO.sub.2 or ThO.sub.2. The composition retains its oxygen ion transport capabilities even after prolonged annealing.

Description

BACKGROUND OF THE INVENTION

FIELD

This invention relates to ceramics generally and to stable bismuth oxide ceramics specifically.

STATE OF THE ART

Ceramic electrolytes, such as zirconia, have been used in automotive sensors and other applications wherein the oxygen ion transport capabilities of ceramic electrolytes have been utilized. Cubic bismuth oxide (Bi.sub.2 O.sub.3) is much more effective than zirconia as a conductor of oxygen ions. However, bismuth oxide tends to be unstable and its conductive characteristics are not retained over longer periods of time (e.g. 50 percent diminution over 100 hours) and also its strength tends to diminish. A tendency exists for the cubic phase of bismuth oxide to convert to monoclinic at temperatures below about 730.degree. centigrade (C.).

Some work in the area of stabilization of bismuth oxide has occurred. Similar to stabilization of zirconia, yttria has also been used with some enhancement of stability in bismuth oxide. However, even yttria-bismuth oxide compositions do not have the long-term stability required to be an effective electrolyte for sensors and other oxygen transport devices in industry.

Bismuth oxide is known to exhibit two polymorphs; namely, cubic above 730.degree. and monoclinic below 730.degree. C. The cubic phase, which is stable between about 730.degree. C. and the melting point Bi.sub.2 O.sub.3 of 825.degree. C. is of the CaF.sub.2 type. Its chemical formula suggests that in order to crystallize in the CaF.sub.2 -type structure, there must be 25 mole % anion vacancies in the structure. The very high concentration of oxygen vacancies is believed to be the primary reason for the exceptional ionic conductivity of Bi.sub.2 O.sub.3 in the cubic form. However, pure Bi.sub.2 O.sub.3 cannot be thermally cycled through the transformation temperature since the volume change associated with the cubic to monoclinic transformation leads to disintegration of the material. It has, however, been shown that numerous rare earth and alkaline earth oxides can be added to lower the transformation temperature and thus enhance the stability range of the cubic phase. The published phase diagrams indicate that in most cases the oxide additive extends the stability range of the Bi.sub.2 O.sub.3 -oxide solid solution, which decomposes eutectoidally below about 700.degree. C. In the case of yttria (Y.sub.2 O.sub.3) as the additive, the work of Datta and Meehan shows that the cubic phase can be stabilized to temperatures at least as low as 500.degree. C. and possibly lower.

It has also been documented that oxygen ion conductivity of yttria-stabilized bismuth oxide is at least two orders of magnitude higher than that of stabilized zirconia. Thus, it would appear that yttria-stabilized Bi.sub.2 O.sub.3 would be an ideal candidate as a solid electrolyte in devices which require high ionic conductivity at moderate temperatures. Recent work has shown, however, that long-term annealing treatment at approximately 600.degree. C. leads to the decomposition of Y.sub.2 O.sub.3 -Bi.sub.2 O.sub.3 solid solutions suggesting that the phase diagram given by Datta and Meehan is incorrect. It is to be noted that Datta and Meehan, who used materials of very high purity (typically in excess of 99.99% or greater), experienced difficulty in achieving equilibrium in several of the compositions examined in their work. The significance of the purity of the starting materials will be discussed later. Recent work by inventors also shows that compositions of other rare earth oxide-bismuth oxide systems such as Er.sub.2 O.sub.3 -Bi.sub.2 O.sub.3, Sm.sub.2 O.sub.3 -Bi.sub.2 O.sub.3, and Gd.sub.2 O.sub.3 -Bi.sub.2 O.sub.3 also decompose when annealed at approximately 600.degree. C.

Stability of the solid electrolyte is the principal requirement of any realistic device based on Bi.sub.2 O.sub.3. Thus, it is imperative that the solid electrolyte remain stable over the design life of the device. If this cannot be guaranteed, Bi.sub.2 O.sub.3 -based solid electrolytes will be of little practical value. The fact that it is possible to retain Bi.sub.2 O.sub.3 stabilized by Y.sub.2 O.sub.3 and other rare earth oxides at lower temperatures long enough to make conductivity measurements suggests that the decomposition upon annealing must somehow depend on the kinetics of mass transport. The same is true of other systems such as Er.sub.2 O.sub.3 -Bi.sub.2 O.sub.3, Sm.sub.2 O.sub.3 -Bi.sub.2 O.sub.3 and Gd.sub.2 O.sub.3 -Bi.sub.2 O.sub.3 3 and the like. If so, the factors which suppress the kinetics of mass transport are expected to slow down the kinetics of destabilization. Assuming that the products of the decomposition are formed by a diffusional process, it would appear that factors which suppress the pertinent diffusion coefficient would impart kinetic stability. Conversely, factors which enhance the diffusion coefficient would lead to rapid deterioration of the solid electrolyte.

SUMMARY OF THE INVENTION

It has recently been discovered that the inclusion of a finer amount of zirconia or hafnia or thoria in a yttria-bismuth oxide ceramic body greatly enhances the stability of the bismuth oxide, enabling it to remain in the cubic phase over very long periods of time.

Typical bismuth oxide-yttria compositions are 75% bismuth oxide and 25% yttria. Inclusion of zirconia, hafnia, or thoria in amounts up to about 10 mole % enhances the stability of the cubic phase of bismuth oxide by several orders of magnitude.

Stable bismuth oxide ceramic compositions according to the present invention include a) bismuth oxide substantially (i.e. >60%) present in the cubic phase as the predominant component; b) a rare earth oxide present in a quantity of about 10 to about 40 mole percent; and c) an oxide compound having a cationic element with a valence of four or more present in quantities of about 0.i to about 10 mole percent. Such a ceramic composition, present substantially as a high-temperature phase, has excellent oxygen ionic conductivity even after days of use at high temperatures.

In its broader aspect the invention involves the stabilization of bismuth oxide by inclusion of a rare earth oxide such as oxides of yttrium, ytterbium, neodymium, gadolinium, samarium, lanthanum, terbium, erbium, and the like by the incorporation of an oxide of a cation having a valence of four or greater, such as zirconia, hafnia, thoria, titania, stannic oxide, tantalum oxide, and niobium oxide and the like. The rare earth oxide may be included in amounts up to about 40 mole % while the polyvalent cationic oxide may be included in amounts up to about 10% mole. Generally amounts as small as about 0.5 and particularly at levels of about 2.0 to 5.0 mole % as minimum quantities have been found very effective.

BRIEF DESCRIPTION OF THE FIGURES

FIG. 1 is an optical microphotograph of the as-sintered sample of III(a);

FIG. 2 is the X-ray diffraction peaks of the as-sintered samples of Y.sub.2 O.sub.3 -Bi.sub.2 O.sub.3 with CaO, ZrO.sub.2 and ThO.sub.2 as dopants as well as undoped samples;

FIG. 3 is the X-ray diffraction traces of doped and undoped samples after annealing; FIG. 3(A) at 2% CaO; FIG. 3(B) at 1% CaO; FIG. 3(C) undoped; FIG. 3(D) at 2% ZrO.sub.2 ; FIG. 3(E) at 2% ThO.sub.2 ;

FIG. 4 shows the optical microphotographs of the undoped samples annealed at 600.degree. for various times; 4(A) at 0 hrs.; FIG. 4(B) at 170 hrs.; FIG. 4(C) at 210 hrs.; FIG. 4(D( at 250 hrs; FIG. 4(E) at 290 hrs;

FIG. 5 shows the microstructures of the as-sintered and annealed samples doped with 5.0 mole % ZrO.sub.2 ;

FIG. 6 plots the fraction transformed X(t), determined using integrated peak intensities from X-ray diffraction traces v. ln(t);

FIG. 7 replots the fraction transformed in FIG. 6 as a function of annealing time;

FIG. 8 is a graph plotting transformed fraction versus time at a temperature of 600.degree. for an undoped composition of Bi.sub.2 O.sub.3 -Er.sub.2 O.sub.3 (80/20 mol. %) and the same composition doped with 5 mol. % ZrO.sub.2 ;

FIG. 9 is a graph plotting conductivity versus time for conditions and composition of FIG. 8; (FIG. 10 is a graph similar to FIG. 8 for a Bi.sub.2 O.sub.3 -Gd.sub.2 O.sub.3 composition;

FIG. 11 is a graph similar to FIG. 9 for the Bi.sub.2 O.sub.3 -Gd.sub.2 O.sub.3 composition;

FIG. 12 is a graph similar to FIG. 8 for a Bi.sub.2 O.sub.3 -Sm.sub.2 O.sub.3 composition;

FIG. 13 is a graph similar to FIG. 9 for the Sm.sub.2 O.sub.3 -Bi.sub.2 O.sub.3 composition.

DETAILED DESCRIPTION OF THE PREFERRED EMBODIMENTS

The compositions of the present invention are believed to suppress cation transport and thereby enhance the kinetic stability of Y.sub.2 O.sub.3 -Bi.sub.2 O.sub.3, Er.sub.2 O.sub.3 -Bi.sub.2 O.sub.3, Gd.sub.2 O.sub.3 -Bi.sub.2 O.sub.3, and Sm.sub.2 O.sub.3 -Bi.sub.2 O.sub.3 solid solutions. Since diffusion of Y (or Er, Gd, Sm) and Bi on the cation sublattice is expected to occur by the movement of point defects, namely vacancies and interstitials, the defect mobilities and defect concentrations are expected to determine the interdiffusion coefficient. For example, if the interstitial mobility is lower than the vacancy mobility, enhancement of the interstitial concentration (to a point) at the expense of the vacancy concentration should lead to lower the interdiffusion coefficient. The opposite would be the case if the interstitial mobility is greater than the vacancy mobility.

Incorporation of a dopant of a valence different than that of the host ion seems to alter the defect concentration. For example, a tetravalent cationic dopant (e.g. Zr or Th) should enhance cation vacancy concentration while a divalent cationic dopant (Ca) should enhance interstitial concentration.

Bismuth oxide for use in the present invention will typically have a substantial quantity (>90%) of the cubic (high temperature) phase, and negligible or a very small quantity of the rhombohedral (low temperature) phase. Bismuth oxide also has a rhombohedral high temperature phase which may be used in place of the cubic high temperature phase, specifically in alkaline earth-Bi.sub.2 O.sub.3 oxide systems. The bismuth oxide will be present in amounts from about 50 to about 90 mole %, with a preferred amount being about 75 mole percent.

Dopants useful in the present invention will generally be present in the ceramic composition in an amount not exceeding about 10 mole percent. Dopants for use in the invention will typically have a cationic portion wherein the cation will have a valence of 4.sup.+ or more. Examples of useful dopants include oxides of zirconium, titanium, thorium, hafnium, tantalum, niobium, tungsten, molybdenum, and vanadium. These dopants will preferably be present in an amount of 2 mole %.

Stabilizers for use in the invention are typically rare earth oxides present in an amount of about 10 to about 40 mole percent. Oxides of the entire lanthanide series are believed to work in this invention, although the preferred rare earth oxides are yttria, erbia, gadolinia and samaria.

EXAMPLES

In the present work, baseline material containing 1) 25 mole % Y.sub.2 O.sub.3 and 75 mole % Bi.sub.2 O.sub.3 with CaO, ZrO.sub.2 and ThO.sub.2 as dopants and, 2) 20 mole % Er.sub.2 O.sub.3, Gd.sub.2 O.sub.3, and Sm.sub.2 O.sub.3 and 80 mole % Bi.sub.2 O.sub.3 with 5 mole % ZrO.sub.2 were annealed at 600.degree. C. and 650.degree. C. for up to several days. The samples were subsequently characterized using X-ray diffraction, optical microscopy and electron microprobe analysis. The experimental procedure, results and a discussion of the results are presented in the following. The experimental procedures for the Y.sub.2 O.sub.3 -Bi.sub.2 O.sub.3 system, described in the parent application and those for the Er.sub.2 O.sub.3 -Bi.sub.2 O.sub.3, Gd.sub.2 O.sub.3 -Bi.sub.2 O.sub.3, and Sm.sub.2 O.sub.3 -Bi.sub.2 O.sub.3 are described separately.

EXPERIMENTAL PROCEDURE

(a): Fabrication of Samples Y.sub. 2 O.sub.3 -Bi.sub.2 O.sub.3 System

Y.sub.2 O.sub.3 and Bi.sub.2 O.sub.3 from readily available commercial sources were used for the present work. CaCO.sub.3 was used as the source of CaO while zirconia (ZrO.sub.2) and ThO.sub.2 powders were used directly as dopants. Chemical analyses of the as-received powders, supplied by the vendors, are given in Table I.

TABLE I______________________________________Chemical Analysis of the as-received PowdersMa- Impuritiesterial (concentration)______________________________________Bi.sub.2 O.sub.3 Fe Mg 20 ppm 8 ppmY.sub.2 O.sub.3 La Nd Gd Zr Na Ca Fe 75 ppm 50 ppm 15 ppm 20 ppm 15 ppm 6 ppm 5 ppm______________________________________

Y.sub.2 O.sub.3 and Bi.sub.2 O.sub.3 powders were mixed in the molar ratio of 1:3. The powder mixtures were milled using either zirconia or alumina milling media. The dopants were added to the baseline mixture. The maximum amount of Ca0 added as 2 mole% while that of ZrO.sub.2 and ThO.sub.2 was 10 mole%. Much of the work reported here is on CaO -doped, undoped and ZrO.sub.2 doped (up to 2 mole %) samples. The powders were wet milled in acetone. The powder mixtures were subsequently dried in an oven and pellets were green-formed by die-pressing followed by isostatic pressing. The samples were then sintered in air at 1000.degree. C. for 24 hours. The as-sintered samples were examined using X-ray diffraction with CuK.alpha. radiation. The density of the samples was determined by the conventional fluid immersion technique.

(b): Thermal Treatments

The as-sintered samples were annealed in air at 600.degree. C. and 650.degree. C. for up to several days. The samples were periodically removed from the furnace, characterized and reinserted into the furnace for further annealing.

(c) X-ray Diffraction

After annealing, the samples were ground and then examined by X-ray diffraction. X-ray diffraction traces were obtained using CuK.sub.60 radiation. The objective was to determine the phases present, their structures and the respective lattice parameters and estimate the extent of phase transformation on the basis of integrated peak intensities. The integrated peak intensities were determined using the peak of the CaF.sub.2 structure and the (on the hexagonal basis) peak of the rhombohedral phase which formed upon annealing in the undoped and CaO-doped samples.

(d) Optical Microscopy and Chemical Analysis

Samples were polished and thermally etched at 900.degree. C. for 60 minutes in order to reveal the grain structure. For the samples that had undergone phase transformation, no thermal etching was done so as to prevent reconversion of the transformed phase(s) into the original cubic phase. The relative difference between the hardness of the two phases led to relief polishing which was adequate to distinguish the morphology of the phase(s). Chemical analysis was conducted on Cameca SX50 electron microprobe. The objective was to determine the compositions of the original solid solutions as well as the decomposed samples.

(e): Sample Preparation: Er.sub.2 O.sub.3 -Bi.sub.2 O.sub.3, Gd.sub.2 O.sub.3 -Bi.sub.2 O.sub.3, and Sm.sub.2 O.sub.3 -Bi.sub.2 O.sub.3 Systems

Bi.sub.2 O.sub.3 and Re.sub.2 O.sub.3 (where Re=Er, Gd or Sm) were weighted in proportion such that Bi.sub.2 O.sub.3 =80 mol. % and RE.sub.2 O.sub.3 =20 mol. %. The powders were mixed and ball-milled in acetone for 72 hours. Subsequently, 5 mol. % ZrO.sub.2 was added to the powder mixtures. The mixtures were mixed for additional 8 hours. The powder mixtures were dried in vacuum and agglomerates were sieved through 200 mesh nylon cloth. Disks of dimensions approximately 16 mm in diameter and 3 mm in thickness were green formed by die pressing followed by isostatic pressing under a pressure of about 200 MPa. The disk compacts were first bisque-fired in air at 450.degree. C. for 12 hours to remove any volatiles. Subsequently, the disks were sintered in air at 940.degree. C. for 24 hours. In order to metastably retain the high temperature phase, which is the high conductivity cubic phase, samples were rapidly cooled to room temperature in 30 minutes. The samples will be referred to as as-sintered samples.

(f) Thermal or Heat Treatment

The as-sintered samples were annealed in air at 600.degree. C. for various periods of time. The samples were periodically removed from the furnace, characterized with regards to the phase present by X-ray diffraction and re-inserted into the furnace for further annealing. These samples will be referred to as as-annealed samples. The samples were annealed for over 160 hours.

(g) X-Ray Diffraction Analysis

The as-sintered and as-annealed samples were surface ground to remove surface layer and then examined by X-ray diffraction with CuK.alpha. radiation. The objective was to determine the phases present and the extent of phase transformation; the lower the extent of transformation, the higher is the phase stability.

(h) Ionic Conductivity Measurement

The ionic conductivity was measured using Hewlett Packard 4192A Impedance Analyzer. Porous gold electrodes were deposited both surfaces of the disk-shaped samples and the conductivity was measured by a.c. two-probe complex admittance method. The frequency range was between 5 hz and 100 khz. All of the conductivity measurements were conducted at 600.degree. C. as a function of time. The objective was to explore the effectiveness of the dopants in enhancing the phase stability (cubic phase) and preventing or minimizing the degradation in conductivity.

RESULTS

III. Y.sub.2 O.sub.3 -Bi.sub.2 O.sub.3 System

III (a): As-Sintered Samples

The as-sintered, undoped samples typically had a density in excess of 7.85 g/ml. Samples doped with ZrO.sub.2 were lower in density with density decreasing with increasing ZrO.sub.2 content. Samples containing 2 mole % and 10 mole% ZrO.sub.2 had densities of 7.67 and 7.33 gms/ml. CaO-doped samples also exhibited somewhat lower density which was attributed to the residual CaCO.sub.3 that may have been present after the densification had begun. The as-sintered samples were orange in color. An optical photomicrograph of the as-sintered sample is shown in FIG. 1. The same was thermally etched at 900.degree. C. for 60 minutes. Typical grain size of the as-sintered samples was about 23.mu.m. As seen in the micrograph, the porosity can be seen within the grains as well as at the grain boundaries. X-ray diffraction of the as-sintered samples with and without any dopants indicated that all of them were single phase with CaF.sub.2 as the only crystal type present. FIG. 2 shows X-ray diffraction peaks of the as-sintered samples of the following compositions: # 2(a)--2mole% CaO, #2(b)--1 mole % CaO, #2(c)--undoped, #2(d)--2 mole % ZrO.sub.2, and #2(e)--2 mole % ThO.sub.2. As seen in the diffraction trances, the as-sintered samples are all single phase cubic with CaF.sub.2 structure. Full scans over angle range between 20.degree. and 60.degree. showed that the dopants were dissolved in the material For ZrO.sub.2 concentration beyond 7.5 mole %, additional peaks corresponding to the dopant were observed suggesting that the dopant solubility is lower than 7.5 mole %. Within the accuracy of the measurements, the lattaice parameters were found to be independent of the dopant type and concentration. Table II gives the measured lattice parameters as a function of the dopant type and concentration.

III.(b): Thermal Treatments

FIGS. #3(a), #3(b), #3(c), #3(d), and #3(e) show X-ray diffraction traces respectively of 2 mole% CaO-doped, 1 mole % CaO-doped, undoped, 2 mole % ZrO.sub.2 -doped, and 2 mole % ThO.sub.2 -doped samples annealed at 600.degree. C. and various periods of time. As shown in the FIG. , the undoped and CaO-doped samples contain two phases, namely cubic and rhombohedral. However, the ZrO.sub.2 - and ThO.sub.2 -doped samples remained single phase cubic. In the undoped and CaO-doped samples, within the limits of the X-ray analysis, no peaks corresponding to phases other than the cubic and rhombohedral phases were observed. As seen in these figures, the extent of phase transformation is the greatest in the CaO-doped samples with no phase transformation occurring in the ZrO.sub.2 - and ThO.sub.2 -doped samples. FIGS. #4(a), #4(b), #4(c),

TABLE II______________________________________Lattice parameters of the cubic phaseof undoped, ZrO.sub.2 -doped andCaO-doped samples7.5 mole % ZrO.sub.2 Lattice parameter (A)______________________________________7.5 mole % ZrO.sub.2 5.49755 mole % ZrO.sub.2 5.50593 mole % ZrO.sub.2 5.50532 mole % ZrO.sub.2 5.5087Undoped 5.50751 mole % CaO 5.51302 mole % CaO 5.5123______________________________________ #4(d), and #4(e) show optical micrographs of the undoped samples annealed at 600.degree. C. for 0, 170, 210, 250, and 290 hours, respectively. The plate-shaped particles are of rhombohedral phase. As seen in the micrographs, the volume fraction of the rhombohedral phase increases with increasing annealing time.

In the subsequent work, undoped, 1.0 mole% CaO-doped and ZrO.sub.2 -doped samples with 2.0 , 3.0, 5.0, 7.5 and 10.0 mole % ZrO.sub.2 were annealed at 600.degree. C. for up to 1020 hours. After annealing, the samples were ground and polished, examined by X-diffraction, thermally etched (only the ZrO.sub.2 -doped), and examined under an optical microscope. It was observed that all of the ZrO.sub.2 -doped samples remained cubic. The grain size of the samples was also unaltered after the thermal treatment. Microstructures of the as-sintered and annealed samples doped with 5.0 mole % ZrO.sub.2 are shown in FIGS. 5(a) and 5(b), respectively, contrast the undoped samples and samples doped with 1.0 and 2.0 mole % CaO exhibited a substantial amount of transformation within the first 100 hours. In samples doped with 2.0 mole % CaO, the transformation was nearly complete in about 200 hours. Again, within the limits of the accuracy of X-ray analysis, no peaks corresponding to any other phases were observed. Fraction transformed, X(t), determined using integrated peak intensities from X-ray diffraction traces is plotted vs. ln(t) in FIG. 6. The powder mixtures of the samples used in this study were milled using alumina grinding media. Similar experiments done using samples made from powers milled with zirconia media showed that the kinetics of phase transformation were slower. The significance of this will be discussed later.

FIG. 6 suggests that the fraction transformed may be described by an Avrami equation as follows: ##EQU1## In the preceding equation, .tau. is the pertinent relaxation time. The fraction transformed shown in FIG. 6 is replotted as a function of annealing time as ln ln (1/(1-X(t))]vs. ln(5) in FIG. 7. The parameters m and -mln (.tau.) are identified with the slope and intercept respectively in FIG. 7. The experimentally determined values of m are between .about.2 for samples doped with 2 mole % CaO to 3.2 for undoped samples. The relaxation times, .tau., for the undoped, 1 mole % CaO-doped and 2 mole %CaO-doped were determined to be .sup..about. 211, .sup..about. 182 and .sup..about. 101 hours, respectively.

III.(c): Chemical Analysis

Chemical analysis of the undoped, as-sintered samples, conducted using Cameca SX50 electron microprobe, showed that the Y.sub.2 O.sub.3 and the Bi.sub.2 O.sub.3 concentrations were respectively 23.3 mole % and 76.7 mole %. The as-sintered samples were single phase cubic as mentioned earlier. The undoped samples, after annealing at 600.degree. C. for 170 hours, were also chemically analyzed. The samples were two phased, containing both the cubic and the rhombohedral phases. The concentrations of Y.sub.2 O.sub.3 and Bi.sub.2 O.sub.3 in the cubic and the rhombohedral phases were determined to be 23.53 mole % Y.sub.2 O.sub.3, 76.47 mole % Bi.sub.2 O.sub.3 and 23.03 mole % Y.sub.2 O.sub.3, 76.97 mole % Bi.sub.2 O.sub.3, (1% of the absolute value), the compositions of the two phases appear to be nearly identical. In the samples annealed for 210 hours, which contained greater amount of the rhombohedral phase, the compositions of the cubic and the rhombohedral phases were determined to be 23.64 mole % Y.sub.2 O.sub.3, 76.36 mole % Bi.sub.2 O.sub.3 and 23.1 mole % Y.sub.2 O.sub.3, 76.9 mole % Bi.sub.2 O.sub.3, respectively. This suggests that the transformation from the cubic to the rhombohedral phase must be a composition invariant transformation.

III.(d) Er.sub.2 O.sub.3 -Bi.sub.2 O.sub.3 System

FIG. 8 shows a comparison between the ZrO.sub.2 -doped and undoped samples as far as the extent of phase transformation is concerned. After about 400 hours at 600.degree. C., the undoped samples transformed to a about 80% (dark squares) while no transformation was evident in the ZrO.sub.2 -doped samples (open squares). FIG. 9 shows the corresponding ionic conductivity. Note that the ionic conductivity remained unaltered and high (about 0.1 (1/ohm.cm)) for the ZrO.sub.2 -doped samples (open squares) even after about 190 hours at 600.degree. C. By contrast, the undoped samples (dark squares) rapidly degraded to less than 0.001 (1/ohm.cm). This shows the profound effect of ZrO.sub.2 suppressing the phase transformation, enhancing the phase stability and maintaining high conductivity. This effectively enhances the utility of the material as a solid electrolyte.

III.(e) Gd.sub.2 O.sub.3 -Bi.sub.2 O.sub.3 System

FIG. 10 shows a comparison of fraction transformed between the undoped and the ZrO.sub.2 -doped samples. The phase transformation (destabilization) in this system is much more rapid than that in the Er.sub.2 O.sub.3 -Bi.sub.2 O.sub.3 or the Y.sub.2 O.sub.3 -Bi.sub.2 O.sub.3 systems. Consequently, note that essentially 100% of the transformed (low temperature, undesirable) phase is present right at the beginning of annealing treatment (dark circles). By contrast, the ZrO.sub.2 -doped samples are initially fully cubic (high conductivity phase) and transform to less than 10% even after 180 hours (open circles). The conductivity showed similar trends as shown in FIG. 11. Note that the conductivity of the undoped samples is rather low and remains as such with continued annealing (dark circles). The ZrO.sub.2 -doped samples (open circles), on the other hand, initially exhibit much higher conductivity which subsequently decreases due to the formation of the low conductivity phase. Even so, the role of ZrO.sub.2 in stabilizing the high conductivity, cubic phase is readily apparent.

III.(f) Sm.sub.2 O.sub.3 -Bi.sub.2 O.sub.3 System

FIG. 12 shows a comparison between the undoped and ZrO.sub.2 -doped samples with regards to phase stability: undoped (dark triangles) and ZrO.sub.2 -doped (open triangles). Similar trends as in the other two systems are obvious. FIG. 13 shows a comparison between the respective conductivities. Here again, the beneficial effect of ZrO.sub.2 additions is clearly seen; undoped (dark triangles) and ZrO.sub.2 -doped (open triangles).

The results demonstrate the generality of the approach suggested in the parent application and set forth hereinabove. It may thus be concluded that the addition of ZrO.sub.2 (or ThO.sub.2) stabilizes high conductivity cubic phase in the RE.sub.2 O.sub.3 -Bi.sub.2 O.sub.3 system over a wide range of compositions. The RE.sub.2 O.sub.3 concentration may be between 5 to 40% while the dopant (ZrO.sub.2 or ThO.sub.2 or the like tetravalent dopant) may be between 1% and 10%.

Of the compositions represented by FIGS. 8 through 13, those stabilized with erbia provided the best results, approaching yttria in effectiveness. For example, in FIG. 8, virtually no phase transformation was detectable after about 400 hours of annealing at 600.degree. C. (The annealing temperature 600.degree. C. was chosen for tilitarian reasons; this is approximately the operating temperature of most solid oxide fuel cells in which a stable Bi.sub.2 O.sub.3 electrolyte offers advantages because of generally higher conductivity than zirconia, the conventional SOFC electrolyte of choice.)

Higher quantities of stabilizer, e.g. erbia, gadolinea or samaria, approaching 40 mole % generally suppressed phase transformation to a greater degree than lower quantities. The presence of larger quantities of ZrO.sub.2, e.g. higher than 5% also improves phase stabilization. Some compromise in composition may be required inasmuch as the Bi.sub.2 O.sub.3 molar composition is decreased the base conductivity may decrease; however, the loss of conductivity with time at annealing temperatures may be reduced, i.e. the conductivity remains more constant for Bi.sub.2 O.sub.3 compositions containing high quantities of stabilizer and dopant even though the initial conductivity may not be as good as Bi.sub.2 O.sub.3 compositions containing smaller amounts of stabilizer and dopant.

DISCUSSION

While not intending to be bound by one theory of the invention, the following may help to explain the extraordinary results obtained herein. The transport of oxygen is not necessary for the occurrence of phase transformation (although not precluded from occurring) since it is common to both end members in the same stoichiometry. Thus, the factors which increase or decrease the interdiffusion coefficient on the cation sublattice enhance or suppress the phase transformation kinetics respectively. The interdiffusion coefficient, D, on the cation sublattice in the Y.sub.2 O.sub.3 -Bi.sub.2 O.sub.3 system is given by ##EQU2## where X.sub.Y and X.sub.Bi are the mole fractions of and D.sub.Y and D.sub.Bi are the diffusion coefficients of Y and Bi, respectively. Similar diffusion coefficinets can be given in the other system by replacing D.sub.Y and X.sub.Y by D.sub.Er, D.sub.Gd and X.sub.Gd, and D.sub.Sm and X.sub.Sm, respectively. The cation diffusion is expected to occur via point defects, e.g. vacancies and interstitials. Thus, both D.sub.Y and D.sub.Bi are expected to depend upon the concentration of cation vacancies and interstitials and on vacancy and interstitial mobilities. If the vacancy concentration far exceeds the concentration of interstitials, the interdiffusion coefficient varies linearly with the vacancy concentration. Alternatively, if the interstitial concentration far exceeds the vacancy concentration, the interdeffusion coefficient will vary linearly with the interstitial concentration. For a "pure" system, i.e. "intrinsic" system, the vacancy and the interstitial concentrations will be the same and equal to D.sub.F where K.sub.F is the Frenkel product on the cation sublattice. Now, in general the vacancy and the interstitial mobilities will be different. If the vacancy mobility is greater/lower than the interstitial mobility, an increase in the vacancy concentration for the intrinsic value will lead to the enhancement/suppression of the interdiffusion coefficient. Alternatively, a decrease in the vacancy concentration will suppress/enhance the interdiffusion coefficient. The net effect is that at some unique concentrations of the point defects, the interdiffusion coefficient will exhibit a minimum. Indeed a minimum in cation diffusion has been observed by Petuskey and Dieckmann and coworkers in iron aluminate spinal and magnetite, respectively. In these works, the pertinent defect concentrations were altered by varying the partial pressure of oxygen. Attainment of a minimum in the interdiffusion coefficient is consistent with minimizing the kinetics of processes which are controlled by the interdiffusion coefficient.

Deviation of defect concentrations from the intrinsic values can be effected by either changing the partial pressures of the constituents or by the incorporation of aliovalent dopants, the latter approach probably being the more convenient one for the present work. Addition of ZrO.sub.2 or ThO.sub.2 as dopants is expected to enhance the cation vacancy concentration above the intrinsic value. A plausible reaction may be given by ##EQU3## A similar equation can be written for ThO.sub.2 dopant. The preceding equation shows that incorporation of a tetravalent cation leads to an enhancement in the cation vacancy concentration. Concurrently, the interstitial concentration will be suppressed through the Frenkel product. Addition of ZrO.sub.2 or ThO.sub.2 can also lead to an increase in the oxygen interstitial concentration. In principles, both of the defect reactions are correct and thus both, the cation vacancy concentration and the oxygen interstitial concentration, will increase. A change in the defect concentration on the anion sublattice, however, is not expected to affect the kinetics of phase separation and will not be discussed further. It should be mentioned, however, that the relationship between the cation vacancy concentration and the dopant concentration would depend upon the pertinent electroneutrality condition which in turn will depend upon the majority defects. In such a case, the knowledge of anion defect concentration may be necessary for determination of the relationship between the dopant concentration and the cationic point defects.

Incorporation of a cationic dopant of a valence less than 3 should enhance the cation interstitial and oxygen vacancy concentrations. Again, oxygen defect concentration is not relevant as far as this discussion is concerned. Dissolution of CaO may occur by the following reaction: ##EQU4## where "I" denotes an interstitial.

The preceding discussion suggests that ZrO.sub.2 and CaO as dopants are expected to affect the interdiffusion coefficient, D, in opposite ways. Indeed, it was observed that the kinetics of phase transformation were enhanced by CaO and were suppressed by ZrO.sub.2. The effect observed is so profound that the addition of 2 mole % ZrO.sub.2 stabilizes the cubic phase at 600.degree. C. for at least 1020 hours, the maximum duration of time that the samples were annealed. By contrast, phase transformation was essentially 100% complete in .sup..about. 290 hours in samples doped with 2 mol % CaO. These results indicate that the cation vacancy mobility is lower than the cation interstitial mobility in Bi.sub.2 O.sub.3 -Y.sub.2 O.sub.3 solid solutions. The fact that the samples milled with zirconia media exhibited slower kinetics is in accord with expectations since some amount of media wear must have occurred thereby inadvertently doping the samples with zirconia. The present work also shows that Bi.sub.2 O.sub.3 -Y.sub.2 O.sub.3 cubic solid solutions can be kinetically stabilized.

The examples set forth hereinabove generally dealt with yttria as the preferred stabilizer for Bi.sub.2 O.sub.3. While excellent results were obtained with a Bi.sub.2 O.sub.3 -Y.sub.2 O.sub.3 system doped with zirconia, as other tetravalent metal oxide, other rare earth elements, as set forth hereinabove, are useful stabilizers in place of yttria.

Although the invention has been described with a certain degree of particularity in both composition and theory, it is understood that the present disclosure has been made only by way of example and that numerous changes in details of composition may be resorted to without departing from the scope of the following claims.

Claims

1. A stable bismuth oxide ceramic composition consisting essentially of:

bismuth oxide substantially present in a high temperature phase as the predominant component;

a rare earth oxide present in a quantity of about 10 mole percent to about 40 mole percent; and

an oxide compound having a cationic element having a valence of four or more present in minor quantities of about 0.1 mole percent to about 10 mole percent.

2. The stable bismuth oxide ceramic composition of claim 1 wherein the rare earth element of said rare earth oxide is selected from the group consisting of yttrium, lanthanum, cerium, neodymium, samarium, europium, gadolinium, terbium, erbium, thulium and ytterbium.

3. The stable bismuth oxide ceramic composition of claim 2 wherein said cationic element of valence four or more is selected from the group consisting of zirconium, hafnium, thorium, tantalum, niobium, tungsten, molybdenum, titanium and vanadium.

4. The stable bismuth oxide ceramic composition of claim 3 wherein said oxide compound is present in an amount of about 5 mole percent.

5. The stable bismuth oxide of claim 4 wherein said cationic element is thorium and said rare ear oxide is yttria.