Polymeric Routes to Silicon Carbide and Silicon ... - ACS Publications

32 Polymeric Routes to Silicon Carbide and Silicon Nitride Fibers Downloaded by TUFTS UNIV on June 20, 2017 | http://pubs.acs.org Publication Date: May 5, 1989 | doi: 10.1021/ba-1990-0224.ch032

William H. Atwell Dow Corning Corporation, Midland, MI 48686-0995

The chapter gives a brief review of the preparation and use of silicon -containing preceramic polymers to prepare ceramic fibers. Several key issues are discussed that must be addressed if this area of technology is to continue to advance.

TTHE USE OF CHEMC IAL APPROACHES

to improve the processing, properties, and performance of advanced ceramic materials is a rapidly growing area of research and development. One approach involves the preparation of organometallic polymer precursors and their controlled pyrolysis to ceramic materials. This chapter will review the preparation and application of silicon-, carbon-, and nitrogen-containing polymer systems. However, the discussion is not exhaustive; the focus is on systems with historical significance or that demonstrate key technological advances. Silicon-containing preceramic polymers are useful precursors for the preparation of ceramic powders and fibers and for ceramic binder applications (I). Ceramic fibers are increasingly important for the reinforcement of ceramic, plastic, and metal matrix composites (2, 3). This chapter will emphasize those polymer systems that have been used to prepare^ceramic fibers. An overview of polymer and fiber processing, as well as polymer and fiber characterization, will be described to illustrate the current status of this field. Finally, some key issues will be presented that must be addressed if this area is to continue to advance.

0065-2393/90/0224-0593$06.00/0 © 1990 American Chemical Society

Zeigler and Fearon; Silicon-Based Polymer Science Advances in Chemistry; American Chemical Society: Washington, DC, 1989.

594

SILICON-BASED POLYMER SCIENCE: A COMPREHENSIVE RESOURCE

Routes to Silicon Carbide Fibers The earliest work on silicon-earbide-related fibers was by Verbeek and Win ter (4). Using the principles developed earlier by Fritz and co-workers (5), Verbeek and Winter (4) reported that the high-temperature pyrolysis of tetramethylsilane or methylchlorosilanes gives branched polycarbosilane (PCS) polymers containing a structure with alternating silicon and carbon atoms (equation 1). (CH3)4Si — ^ polycarbosilane (PCS)

(1)

ψ spin aid ψ pyrolysis

Downloaded by TUFTS UNIV on June 20, 2017 | http://pubs.acs.org Publication Date: May 5, 1989 | doi: 10.1021/ba-1990-0224.ch032

SiC fibers Spinning of PCS, often with the use of an organic-polymer spinning aid such as poly(ethylene oxide), followed by curing and high-temperature pyrolysis gives black silicon-carbide-like fibers. However, interest in this technology was ignited by the extensive and pioneering work of Yajima and co-workers (6, 7). In an early work, Yajima and Hayashi (6) applied the known Kumada rearrangement to Burkhard's poly(dimethylsilane) polymer and obtained melt-spinnable polycarbosilane polymers (equation 2). [(CH3)2SiL

PCS ^*$>

eure and p y r o l y ^

^

&

^

(

£

)

φ Ti(OR) 4

poly(titanocarbosilane) —» SiTiC fibers (3) The spun fibers were cross-linked (cured) by air oxidation and pyrolyzed to give silicon-carbide-type fibers. Yajima et al. (7) reported further that heteroatoms such as titanium could be incorporated into the polymers and ceramic fibers to enhance their stability (equation 3). Following the pioneering works of Verbeek and Yajima, numerous in vestigators began to explore the scope of polymer systems that would provide useful ceramic compositions. West and co-workers (8) prepared polysila styrene (PSS) polymers by the sodium coupling of phenylmethyldichloro silane and dimethyldichlorosilane (equation 4). (C 6 H 5 )(CH 3 )SiCl 2 + ( C H 3 ) 2 S i C l 2 ^ polysilastyrene (PSS) φΔ

SiC powders

(4)

One aspect of their elegant and extensive studies dealt with the pyrolysis of these polymers to ceramic powders, fibers, and film. Schilling and co-workers (9) used the known disilylation reaction of


32.

ATWELL

Polymeric Routes to Silicon Carbide and Silicon Nitride Fibers 595

chlorosilanes to produce polycarbosilane polymers with controlled levels of cross-linking (equation 5). (CH 3 ) 2 SiCl 2 + (CH 2 = C H ) ( C H 3 ) S i C l 2 - ^ polycarbosilane φ

Δ

SiC powders

(5)


Baney and co-workers (10) used the known Si-Si and Si-Cl bond redistri bution of methylchlorodisilanes to prepare poly(methylchlorosilane)s. Py rolysis of these branched polysilane polymers gives nearly stoichiometric amounts of silicon carbide (equation 6). methylchlorodisilane fraction

R4PC1 catalyst

> p0ly(methylchlorosilane) ψ

Δ

SiC powders and fibers (6)

Routes to Silicon Nitride Fibers The earliest report of fibers containing silicon and nitrogen is by Verbeek (11). Chlorosilanes and amines are used to prepare polycarbosilazane poly mers, which are converted into amorphous SiCN-containing ceramic fibers (equation 7). CH 3 SiCl 3 + C H 3 N H 2

(CH 3 SiNCH 3 ) x (polycarbosilazane) φ spin aid φ

Δ

SiCN amorphous fiber

(7)

Seyferth and co-workers (12) used the ammonolysis of dichlorosilane to pre pare carbon-free polysilazanes that could be converted into silicon nitride (equation 8). H 2 SiCl 2 + N H 3

(H 2 SiNH) x -^-> Si 3 N 4 powder

(8)

Gaul and co-workers (13,14) prepared a variety of polysilazanes by using chlorosilanes and hexamethyldisilazane to control polymer molecular weight, rheology, and spinnability (equation 9). Fibers were prepared from the poly(methyldisilylazane) (MPDZ) system (equation 10). CH 3 SiCl 3 + [(CH3)3Si]2NH -> (CH 3 SiNH),

SiCN powder

(9)

methyldichlorosilane + [(CH3)3Si]2NH -> (CH 3 SiSiCH 3 NH) x (MPDZ) Φ

SiCN amorphous powder Zeigler and Fearon; Silicon-Based Polymer Science Advances in Chemistry; American Chemical Society: Washington, DC, 1989.

(10)

596


Some key observations can be drawn from these early pioneering efforts. • Pyrolysis of organosilicon polymers in nonoxidizing atmo spheres provides a route to nonoxide ceramic compositions. • The yield of inorganic char (residue) increases with increasing polymer cross-linking.


• Both amorphous and crystalline materials are obtained, de pending on the polymer compositions and temperatures used. Recent workers have extended this field in several key areas. Cannady (15) and LeGrow and co-workers (16) extended the use of hexamethyldisilazane and prepared poly(hydridosilazane) (HPZ) polymers (equation 11). HSiCl 3 + [(CH3)3Si]2NH

(HSiNH) 3 / 2 (HPZ) ψ cure φ Δ

S13N4 fiber and powder

(11)

Gas-phase curing of polymer fibers with trichlorosilane followed by pyrolysis gave high-nitrogen, low-carbon, silicon-nitride-type fibers. Burns and co-workers (17) prepared a series of alkyl-, aryl-, and arylalkylsubstituted polysilazane polymers (equation 12), and a mechanistic study of pyrolysis was carried out to determine the effect of substituents on char yield, char composition, and stability of the resulting ceramic powders. RSiCl 3 + N H 3 -> (RSiNH) 3/2

SiCN powders

(12)

Okamura and co-workers (18) have taken air-cured PCS polymer and, through pyrolysis in the presence of ammonia, prepared essentially carbonfree silicon oxynitride fibers (equation 13). However, if the PCS polymer fiber is cured by electron beam radiation (to prevent oxygen addition), the same ammonia pyrolysis conditions provide nearly stoichiometric quantities of silicon nitride fibers (equation 14). PCS PCS

cured PCS - ^ Λ Si 2 N 2 0 fiber e

' e c t r o n b e a m ) cured PCS - ^ Λ Si 3 N 4

(13) fiber

(14)

Quite recently, Arai and co-workers (19) prepared higher molecular weight poly(dihydridosilazane) polymers through amine complexation (equa tion 15). H 2 SiCl 2 + N H 3 -> (H 2 SiNH) x - * Si 3 N 4 powders and

fibers


(15)

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Dry spinning in the presence of spinning aids followed by pyrolysis gives low-carbon silicon-nitride-type fibers. From this recent work, several additional key observations can be drawn: • Char compositions can be controlled quite precisely through the proper choice of polymer composition, pyrolysis temper ature, and atmosphere, as well as through an improved un derstanding of the mechanisms of polymer pyrolysis.


• Many of the inorganic chars are thermally unstable at elevated temperatures, generally as a result of poor control of stoichiometry, that is, high oxygen content and excessively high car bon/ silicon ratios. • Past workers have placed a high emphasis on polymer synthe sis, although recent work is more balanced, with an increasing emphasis on detailed characterization of the char.

Fibers The following section will provide an overview of fiber processing, as well as descriptions of the polymer-derived fibers that have been prepared. A general processing scheme is depicted in Figure 1. Typically, spinning is carried out by melt-spinning extrusion through a spinneret to obtain polymer fibers with a diameter of 10-20 μπι. However, solvent spinning has been reported (19). The fragile polymer fibers need to be cured (cross-linked) below their melting point to prevent coalescence

C Η A R A C Τ Ε R I Ζ A Τ I

ORGANO SILANES POLYMERIZATION J

.

PRECERAMIC POLYMER

I

FIBER

±

CURE

POLYMER FIBER

CROSSLINKED FIBER PYROLYSIS

0 Ν

SPINNING

CERAMIC FIBER -•CHARACTERIZATION

Figure 1. Process scheme for fiber production.



598


during pyrolysis and to increase char yields. Proper handling must be ex ercised to prevent damage to the fiber surface. Pyrolysis must be carried out with careful control of atmospheric con ditions. The gases formed during pyrolysis cause the fibers to undergo a weight loss of 20-40% (60-80% volume decrease), and the fibers achieve a final diameter of 10-15 μπι. This diameter is desirable for many ceramic and plastic composite applications. A more complete discussion of this fiber processing has been presented recently (20). A summary of some of the polymer-derived SiCN fibers is presented in Table I. This table includes the polymer precursor systems for these fibers, the compositional data, and the developing companies. Reliable elemental analyses of these fibers are often difficult to obtain because of incomplete sample combustion. In addition, many of these fiber compositions are being changed as their development proceeds. Table I. Polymer-Derived SiCN Fibers Elemental Fiber Polymer Si Ν Nicalon PCS 58.4 0.1 Tyranno" PCS-Ti 49.3 0.3 Si 3 N 4 (H 2 SiNH)„ 62.5 34.3 SiCN HP2 59.4 28.9 b SiC PMS 69.1 "Thisfibercontains 1.2 wt % titanium. h — means not determined.

Composition (wt%) Total C 0 99.7 31.2 10.1 27.9 20.2 98.9 0.5 3.1 100.4 10.1 3.1 101.5 29.5 1.0 99.6

Company Nippon Carbon Company Ube Industries Toa Nenryo Kogyo K.K. Dow Corning Corporation Dow Corning Corporation

Characterization. Characterization work at all stages of fiber proc essing, as illustrated in Figure 1, is important. The discussion in this chapter will be limited to certain aspects of polymer and ceramic fiber characteriza tion, which will be illustrated with poly(methylsilane) (MPS) and HPZ poly mers and the ceramic fibers derived from these polymers. The characterization of polymer-derived ceramic fibers is an important but somewhat neglected area, although several publications have now dealt with the extensive characterization of preceramic polymers for ceramic fiber spinning and pyrolysis. A wide variety of characterization techniques have been used; in addition to elemental analyses and IR and NMR spectroscopy, gel permeation chromatography (GPC), thermal mechanical analysis (TMA), thermogravimetric analysis (TGA), and temperature-viscosity studies are valuable techniques. The TMA and temperature-viscosity characterization data are particularly useful for predicting polymer melt spinning conditions. As a result of these characterizations, a picture of the fiber structures is now emerging (16,18,21,22). For a more detailed discussion of these techniques, the reader is referred to a recent article (16) on HPZ polymer. The results of X-ray diffraction (XRD), transmission electron microscopy


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(TEM), high-resolution electron microscopy (HREM), and electron spec troscopy (XPS) have shown that the continuous phase in these fibers is amorphous. For some fibers (21), the evidence supports the presence of a microcrystalline β-silicon carbide, dispersed in the amorphous phase, with an average crystallite size of 2-4 nm (Figure 2). The presence of a microcrystalline carbon phase similar to graphitic carbon is supported (21) by elemental analyses and Raman, magic-angle-spinning NMR ( 1 3 C MAS NMR), and ESR (electron spin resonance) spectroscopy.

10e

20e

30°

40e

50°

60°

2Θ Figure 2. X-ray diffractograms of powdered ceramic fibers (Cu K ). CGN is ceramic-grade Nicalon, and SGN is standard-grade Nicalon. MPDZ-PhVi is poly(methyldisilylazane)-derived fiber. See reference 21 for complete details. 2

The chemical bonding in these ceramic fibers appears to approach a structure in which individual silicon atoms are simultaneously bonded to C, N, and Ο in a random distribution. This structure is most dramati cally demonstrated by 2 9 Si MAS NMR (21, 23) spectroscopy. In Si-C-O fi bers, five tetrahedral silicon bond arrangements are possible: SiC 4 , SiC 3 0, S i C 2 0 2 , SiC0 3 , and Si0 4 . Figure 3 shows that all five arrangements are indicated in the spectra of standard-grade Nicalon (SGN) fiber. In contrast, the 2 9 Si MAS NMR spectrum of MPS-derived silicon carbide fibers (Figure



600

-11 Model Compounds: SiC 0 = Me SIX (x = OSiMe OMe) 3

3

3>

SiC = Me Si 4

4

S i C 0 = Me SiX (x = OSiMe ) 2

2

2

2

3

SiC0 = MeSiX (x = OSiMe ) 3

3

Si0 = SIX (x = OSiMe )

Intensity

4

4

3

D P h 2 = (Ph SiO) 2

T Downloaded by TUFTS UNIV on June 20, 2017 | http://pubs.acs.org Publication Date: May 5, 1989 | doi: 10.1021/ba-1990-0224.ch032

3

P h

-34

KSU

= (PhSI0 ) 3/2

-74

(SIC2| ° 2 > \ | D 2 β-S C \ Ph

(SiCC

?l a

pri 0

50

x

Τ

x

-111

\jIV

-50

-100

.. -150

Chemical Shift, δ

-200 PPM

Figure 3. Si MAS NMR spectrum of powdered SGN fibers. 29

4) is rich in SiC 4 structure, with smaller amounts of S1C3O and S i C 2 0 2 . Additional data are available for other fibers (16, 21, 23). Physical Properties. Polymer-derivedfibersshow low densities (de termined by density gradient column technique) compared with that re ported for crystalline silicon carbide and silicon nitride (—3.20 g/cm3). Table II gives the reported values of tensile strength, elastic modulus, and density for several fibers. Numerous nitrogen and krypton BET (BrunauerEmmett-Teller) surface area measurements have been carried out on these fibers. Essentially, the surface area is the geometric area. In addition, mer cury porosimetry studies show the absence of significant surface-intercon nected porosity (24). Present studies, including small-angle X-ray-scattering studies, suggest the existence of considerable closed porosity on the 2-4nm level as an explanation for the low densities and low surface areas. Finally, the surface and bulk compositions of these fibers can vary dra matically. Figures 5 and 6 show the scanning Auger microscopy (SAM) profiles of SGN and HPZ ceramic fibers, respectively. Because the fi ber-matrix interfaces are critical in many applications (2), this type of char acterization of fiber surface composition and chemistry must be carried out.


32.

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-15

|SiC202

ι

I

ι ι ι ι

100

I

ι ι ι ι

ι ri π ι

50

0

-50

I I I I [ ι ι ι ι

-100

ι

ι ι ι ι

-150

ι

I I I

-200 PPM

Chemical Shift, δ Figure 4. Si MAS NMR spectrum of powdered MPS-derived ceramic fiber. 29

The preceding brief discussion, shows that a wide variety of character ization techniques have been used to characterize preceramic polymers and the derived ceramic fibers. Although some structural details remain elusive, the structural understanding of thesefibershas advanced dramatically during the past 5 years.

Key Issues The remaining discussion will focus on four key issues that must be addressed and understood to advance the technology of polymer-derived ceramic fibers. Table II. Properties of SiCN Fibers Fiber S13N4 S13N4

Nicalon Tyranno SiC

Tensile Strength

Elastic Modulus

Density

2.45 2.80 2.98 2.80 1.75

196 217 196 196 210

2500 2620 2555 2400 2750

(GPa)

(GPa)


fe/m3)


602


60

120

180 240 Depth (nm)

300

Figure 5. SAM profile of SGN fiber.


360

32.

ATWELL


Polymer Properties. Thefirstissue is that fiber processing is driven by polymer properties and chemistry. Some key polymer requirements re lated to the process scheme presented in Figure 1 are the following: • Rheology—stable under melt-spinning conditions • Chemical reactivity—uniform, rapid, and controllable cure in solid state • Composition—uniform and highly pure


• Pyrolysis—high char yield of controllable chemical composition Control of all these requirements is essential if suitable fiber processing and properties are to be developed. Control of polymer purity and curability are the critical areas for future work. Some purity requirements will become quantified as this discussion proceeds. With regard to fiber cure, gas-phase chemical cure and energetic methods have been most effective to date (16,18). Fiber processing involving the dry spinning of thermally unstable polymers (4,11, 19) may be of future importance, because the need for a separate cure step is eliminated. Fiber Strength. The second key issue is that fiber tensile strengths must be improved. This property in brittle materials is controlled by flaw size and distribution. The tensile strength of current fibers is limited by the purity of the polymer precursor (20, 22). An analysis of tensile strength versus flaw size suggests that flaws with diameters of —1 μπι limit fiber tensile strength to —1.75 GPa. Reduction offlawdiameter to 0.1 μπι gives a fiber strength of -2.86 GPa (20). Flaw frequency and flaw size can be limiting factors in determining the strength of these ceramic fibers. Tensile strength measurements are typically made with a single fiber having a gauge length of 2.54 cm. A frequency of two flaws per 2.54 cm would be expected if only 1 ppm of Ι-μπι-diameter spherical flaws is present in a ΙΟ-μπι-diameter fiber. This flaw population would limit the attainable tensile strength to —1.75 GPa. Fractographic analysis of ceramic fiber fracture surface by scanning electron microscopy (SEM) is an invaluable tool for the study of strength-limiting flaws (20, 22). Obviously, fiber processing also provides opportunities to introduce fiberweakening flaws, and the attention required in each process step cannot be overemphasized. Elastic Modulus. The third key issue relates to improvements re quired in the elastic modulus of these ceramic fibers. Studies indicate that modulus improvements will depend on careful control of composition and increases in fiber density (24). Earlier discussions in this chapter summarized the low densities of these ceramic fibers relative to their crystalline coun terparts, as well as the existence of considerable pore volume.


604



In a recent study (24), significant increases in the density of HPZ-derived fiber was achieved without changes in weight or composition. The resultant modulus changes are shown in Figure 7. Extrapolation suggests that the predicted modulus for amorphous, fully dense HPZ-derived fiber corresponds to a value of —360 GPa.

Thermal Stability. The last key issue relates to improvements required in the high-temperature stability of ceramic fibers. Thermal stability may refer to the percent retention of initial tensile strength either at temperature or subsequent to heat aging at temperature. The fibers known at present show a significant reduction in tensile strength after heat aging at 1200-1400 ° C (25-27). The results can vary widely not only with fiber type but also with fiber batch and processing conditions, such as atmospheric conditions and sample geometry. Careful definition and control of protocols are required for reproducible results. One detailed study (28) suggests that strength loss is a result of crystal growth and the formation of fine holes or pores around grain boundaries. Other studies (25) suggest that strength loss of Nicalon during nitrogen aging is due to growth of existing internal or surface flaws. However, when aged in argon, Nicalon fiber strength reduction appears to be due to grain growth and pores between grains. Although the mechanisms of strength reduction are not understood, this factor must be controlled if these fibers are to attain their full potential use in reinforced structures.


32.

ATWELL


Summary The polymer-derived approach is a viable method for the preparation of ceramic fibers for reinforcement applications. During the past decade, sig nificant advances have been made in materials, processes, and character ization capabilities. This approach holds great promise for future advances; however, focus must be placed on addressing the key issues that have been identified.


Acknowledgments I acknowledge the support of a Defense Advanced Research Projects Agency (DARPA)-funded Air Force contract (F33615-83-C-5006) administered by the Air Force Wright Aeronautical Laboratories/Materials Laboratory and the Dow Corning Corporation. I thank A. Zangvil, University of Illinois, for the SAM profiles.

References 1. For a recent review of this subject, see Baney, R. H.; Chandra, G. In Encyclo

pedia of Polymer Science and Engineering; Wiley: New York, 1988; Vol. 13, p

312. 2. Mah, T.; Mendiratta, M . G.; Katz, A. P.; Mazdiyasni, K. S. Ceram. Bull. 1987, 66, 304-317. 3. Petrisko, R. Α.; Stark, G. L. SAMPE Proc. 1988, 33, 1015. 4. Verbeek, W. ; Winter, G. Ger. Offen. 2 236 078, 1974. 5. Fritz, G.; Matern, E. Carbosilanes Synthesis and Reactions; Springer: New York, 1986. 6. Yajima, S.; Hayashi, J.; Omori, M . Chem. Lett. 1975, 931. 7. Yajima, S.; Iwai, T.; Yamamura, T.; Okamura, K.; Hasegawa, Y. J. Mater. Sci. 1981, 16, 1349. 8. Ultrastructure Processing of Ceramics, Glasses and Composites; West, R. C.;

Hench, L. L.; Ulrich, D. R., Eds.; Wiley: New York, 1984; p 235. 9. Schilling, C. L.; Wesson, J. P.; Williams, T: C. Am. Ceram. Soc. Bull. 1983, 62, 912. 10. Baney, R. H.; Gaul, J. H.; Hilty, T. K. Organometallics 1983, 2, 859. 11. Verbeek, W. U.S. Patent 3 853 567, 1974. 12. Seyferth, D.; Wiseman, G. H . ; Prud'Homme, C. J. Am. Ceram. Soc. 1983, 66, C-13. 13. Gaul, J. H . U.S. Patent 4 312 970, 1982. 14. Emergent Process Methods for High Technology Ceramics; Davis, R. F.; Palmer,

Η., III; Porter, R. L., Eds.; Plenum: New York, 1984; Vol. 17,p235. 15. Cannady, J. P. U.S. Patent 4 535 007, 1985. 16. LeGrow, G. E.; Lim, T. F.; Lipowitz, J.; Reaoch, R. S. Am. Ceram. Soc. Bull. 1987, 66, 363. 17. Burns, G. T.; Angelotti, T. P.; Hanneman, L. F.; Chandra, G.; Moore, J. A. J. Mater. Sci. 1987, 22, 2609. 18. Okamura, K.; Sato, M.; Hasegawa, Y. Ceram. Int. 1987, 13, 55.


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19. Arai, M.; Hayato, N . ; Osamu, F.; Takeski, I. Eur. Patent Appl. 86308752.4, 1986, 20. Salinger, R. M.; Barnard, T. D.; Li, C. T.; Mahone, L. G. SAMPE Q. 1988, 19, 27. 21. Lipowitz, J.; Freeman, Η. Α.; Chen, R. T.; Praek, E. R. Adv. Ceram. Mater. 1987, 2, 121. 22. Sawyer, L. C.; Jamieson, M.; Brikowski, D.; Haider, I.; Chen, R. T. J. Am.

Ceram. Soc. 1987, 70, 798. 23. Lipowitz, J.; Turner, G. L. Polym. Prepr. (Am. Chem. Soc., Div. Polym. Chem.)

1988, 29, 1. 24. Rabe, J. Α.; Lipowitz, J.; Frevel, L.; Sander, W. Proceedings of the 12th Annual


Conference on Composites and Advanced Ceramics; Cocoa Beach, FL, 1988, in

press. 25. Lipowitz, J., Dow Corning Corporation, Midland, MI, unpublished results. 26. Clark, T. J.; Jaffe, M.; Rabe, J.; Langley, N . R. Ceram. Eng. Sci. Proc. 1986, 7, 901. 27. Sawyer, L. C.; Chen, R. T.; Haimbuch, F., IV; Hagert, P. J.; Prack, E. R.; laffe, M. Ceram. Eng. Sci. Proc. 1986, 7, 914.

28. Langley, N . R.; Filsinger, D. H.; Rabe, J. Α.; laffe, M.; Clark, T. J. Proceedings of a Joint NASA/Department

of Defense Conference; N A S A Conf. Publ. 2445;

Government Printing Office: Washington, DC, 1986; ρ 7.

29. Proceedings of a Symposium on High Temperature Chemistry; Ishikawa, H.;

Ichikawa, H.; Teranishi, H.; Munir, Ζ. Α., Eds.; Electrochemical Society: Pen nington, NJ, 1988; Vol. 4, ρ 205.

RECEIVED for review December 20, 1988. ACCEPTED revised manuscript March 27, 1989.


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