Synthesis, microstructure and properties of SiCN ceramics prepared from tailored polymers

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  • Synthesis, microstructure and properties of SiCNceramics prepared from tailored polymers

    G. Ziegler, H.-J. Kleebe*, G. Motz, H. Muller, S. Tral, W. WeibelzahlInstitute for Materials Research (IMA), University of Bayreuth, D-95440 Bayreuth, Germany

    Dedicated to Prof. Dr. S. Somiya on the occasion of his 70th birthday

    Abstract

    Different liquid polymers in the system SiCN with tailored structures were prepared by ammonolysis from functionalized chlorosilanes.

    Crosslinking to an unmeltable polymer with initiators at low temperatures and subsequent ceramization were studied applying 29Si solid-

    state nuclear magnetic resonance (NMR) spectroscopy in combination with Fourier transformed infrared (FTIR) spectroscopy and

    thermoanalytical techniques.

    Microstructure development, in particular, the devitrification of the corresponding bulk polymer-derived SiCN glasses was investigated

    by scanning electron microscopy (SEM) and transmission electron microscopy (TEM). Preparation of monolithic samples was performed

    by mixing liquid polysilazane with SiCN-powder particles, derived from the same precursors by heat treatment at 3008C, and subsequentannealing at temperatures exceeding 10008C to initiate crystallization. Depending on the functionalities of the SiCN-precursor and theprocessing conditions, different microstructures were obtained.

    The material prepared from the HVNG precursor revealed a homogeneous amorphous micro structure with only a small fraction of

    crystallized spherical inclusions after exposure at 15408C for 6 h in nitrogen atmosphere. In contrast, investigating ceramic monolithsderived from another SiCN precursor, a different crystallization sequence was observed. The material derived from the HPS precursor

    showed crystallization of large a-Si3N4 grains within the bulk. As will be discussed in detail, devitrification of these polymer-derivedglasses is promoted by local rearrangements and possible phase separations within the amorphous bulk. Moreover, local decomposition and

    residual porosity can affect the crystallization behavior, which strongly differs depending on the polymer employed.

    In addition to the crystallization phenomena observed, different oxidation response was monitored for the two SiCN ceramics discussed

    here. Moreover, fracture strength and hardness data were recorded, which, however, did not substantially differ between the polymer-

    derived ceramics investigated. # 1999 Elsevier Science S.A. All rights reserved.

    Keywords: Synthesis; Microstructure; Properties; SiCN ceramics; Tailored polymers

    1. Introduction

    Organometallic compounds (precursors) have attracted

    considerable interest in recent years, owing to their promis-

    ing potential for the formation of high-purity non-oxide

    ceramics, amorphous fibers and surface coatings [13].

    Since the pioneering work of Verbeek and Winter [4] in

    addition to Yajima [5] in the mid 1970s, a wide variety of

    precursors have been developed for the preparation of

    different non-oxide ceramics [68]. The major advantages

    of such polymer-based materials is their intrinsic homoge-

    neity on an atomic level, low processing temperatures, since

    the precursors can be transformed into amorphous covalent

    ceramics at temperatures between 80010008C, and the

    applicability of established polymer processing techniques.

    In general, processing of ceramic materials via organome-

    tallic compounds involves the synthesis of the precursor

    from monomer units followed by crosslinking into an

    unmeltable, preceramic network and finally the pyrolysis

    at elevated temperatures. The latter heat treatment initiates

    the organicinorganic transition and results in an amor-

    phous, non-oxide covalent glass. Post-annealing of such

    amorphous non-oxide ceramics at temperatures exceeding

    10008C yields a partially or completely crystallized ceramic.A number of studies reported in literature address the

    pyrolysis behavior of the polymeric precursors at tempera-

    tures around 10008C, whereas less work has been focused onthe crystallization behavior and the thermal stability of these

    precursor-derived amorphous structures. TEM investiga-

    tions by Monthioux and Delverdier [9,10] as well as Kleebe

    et al. [11,12] focused on the crystallization phenomena

    Materials Chemistry and Physics 61 (1999) 5563

    *Corresponding author.

    E-mail address: achim.kleebe@uni-bayreuth.de (H.-J. Kleebe)

    0254-0584/99/$ see front matter # 1999 Elsevier Science S.A. All rights reserved.PII: S 0 2 5 4 - 0 5 8 4 ( 9 9 ) 0 0 1 1 4 - 5

  • observed in SiCN-based glasses, while the work reported by

    Bill and Aldinger [13] described the microstructure devel-

    opment of monolithic SiBCN and SiPCN ceramics. Up to

    now, little understanding has been developed concerning the

    problem of thermal degradation of the amorphous SiCN-

    ceramic materials. Various aspects may be important, start-

    ing from the polymer architecture, the chemical composi-

    tion, the residual porosity within the amorphous structure

    (open/closed system), local kinetics and thermodynamics as

    well as the ambient atmosphere. It is also thought that the

    aforementioned parameters affect the resulting material

    properties such as fracture strength, fracture toughness

    and oxidation resistance [1416].

    Here we report on the study of two different SiCN

    ceramics, derived from tailored precursors, starting from

    polymer synthesis followed by detailed microstructure

    characterization of bulk ceramics in addition to the acquisi-

    tion of their corresponding properties such as oxidation

    behavior and mechanical response. This general approach

    synthesis-characterization-microstructure reflects the con-

    cept followed at the Institute for Materials Research in

    Bayreuth.

    2. Experimental procedures

    2.1. Polymer synthesis and characterization

    All preparation steps were carried out in an inert gas

    atmosphere due to air and moisture sensitivity of both educts

    and products. Synthesis followed standard procedures [7],

    i.e., dissolving of different di- and trifunctionalized chlor-

    osilanes in toluene and passing ammonia through the solu-

    tion. When the reaction has ended, it is necessary to purge

    with argon to eliminate excess ammonia. Subsequent filtra-

    tion of the ammonium chloride from the reaction mixture

    and distillation of the solvent leads to colorless or pale

    yellow silazane precursors. Rheological measurements

    were performed on a cone-plate-viscosimeter (Rheolab

    MG 10, Physica Metechnik, Germany). Molecular weights

    were determined cryoscopically in cyclohexane or p-xylene.

    The precursors were crosslinked by using dicumylper-

    oxide as a radicalic initiator and subsequent thermal treat-

    ment at 3008C for 5 h in N2-atmosphere. For all theexperiments, powder samples were used, which were

    obtained from the as-received unmeltable solids by ball

    milling with zirconia milling media. The powders were

    sieved and the fraction

  • perforation and subsequent light carbon coating to minimize

    electrostatic charging under the electron beam.

    2.4. Properties

    Oxidation stability was examined employing thermogra-

    vimetry (STA 409, Netzsch) at temperatures ranging from

    11008C to 14008C for 72 h in flowing air (150 ccm/min).Before testing, the polished specimens were tempered at

    14508C in N2-atmosphere to exclude any possible mass lossdue to the escape of hydrogen. The hardness values of the

    specimens were determined using a Vickers indenter (98 N),

    with an average of 10 indentations for each sample. Fracture

    strength was measured by four-point bending tests of five

    specimens each using a 40/20 mm support span and a

    crosshead speed of 0.5 mm/min. Youngs modulus was

    determined from the load/displacement curves by following

    Eq. (1)[25]:

    E Fl20l1

    16Jy0; (1)

    where F represents the load applied, l0 is the distance

    between the inner load points, l1 gives the distance between

    the inner and outer supports, y0 the deflection of the center of

    the specimen relative to the position of inner supports, and J

    is the moment of inertia, J bh312

    , where b is the width of the

    specimen and h represents its height in the direction of the

    deflection.

    3. Results

    3.1. Synthesis and characterization of polymers

    The polymer synthesis is mainly based on two concepts.

    First, various reactive functional groups at the silicon atoms

    were introduced to control further branching reactions via

    hydrosilylation (>SiHH2C=CHSi>) and/or polymeriza-tion of the vinyl substituents. These reactions can be

    induced by heating to 3008C or preferably at a lowertemperature of about 1308C by adding a radicalic initiator.Second, modification of the molecular weight and viscosity

    is achieved by either using di- and trifunctional chlorosi-

    lanes or by bonding sterically different substituents to the

    silicon atoms. As a result of both concepts, the polysilazanes

    HVNG and HPS were synthesized and can be described by

    the structural units given in Fig. 1. The polysilazane HVNG

    consist of mixed di- and trifunctional units, i.e., every

    second silicon atom is bridged by two and the other half

    by three nitrogen atoms to other silicon atoms. In contrast,

    the HPS precursor only consist of twofold bridged silazane

    units.

    With the additional possibility of crosslinking, the mole-

    cular weight was raised from 440 g/mol (HPS) to 620 g/mol

    (HVNG). The viscosity of the silazanes also strongly

    depends on their intrinsic structure. Therefore, a viscosity

    increase from a highly liquid (HPS, 0.05 Pas) to a honey like

    precursor (HVNG, 29 Pas) was recorded (compare Table 1).

    In all silazane systems, first a mass change was observed

    during pyrolysis between 1508C and 3508C. In general, atlower degrees of branching (HPS) the mass loss is about 15

    30 wt%, owing to the release of gaseous oligomers besides

    ammonia, as identified by coupled TG-FTIR measurements

    (Fig. 2). When, however, increasing the degree of branching

    (HVNG), the mass loss is reduced to about 5 wt%. A second

    stage of mass loss was observed between 3508C and 7508C,where methane is released which leads to the degradation of

    the organic substituents (methylene groups, ethylene

    bridges). Above 8008C, no significant mass changes wereobserved. After heating to 10008C, the ceramic yield for theHPS in comparison to the HVNG precursor is markedly

    lower (73 versus 82 wt%), due to the escape of more volatile

    Fig. 1. Structural units of the precursors HVNG and HPS.

    Table 1

    Properties of the precursors and the resulting polymer-derived ceramics.

    Educts Precursor Molecular

    Weight (g/mol)

    Synthesis

    Yield (%)

    Viscosity at

    208C (Pas)Ceramic yield

    at 10008CElementary composition at 10008C(mass %)

    ViSiCl3 HVNG 620 92 29 82 Si50.3 O0.85Me(H)SiCl2 C20.6 H

  • oligomers and a higher amount of methyl groups within the

    HPS. In addition, the elementary composition shows a

    higher carbon concentration in the resulting amorphous

    ceramic (Table 1). Changes in the structure of the solid

    intermediates during thermal treatment were investigated by

    29Si solid state NMR spectroscopy. Fig. 3(a) (HVNG) and

    Fig. 3(b) (HPS) show the 29Si-spectra of both polysilazanes

    as a function of pyrolysis temperature up to 15008C. Solidstate NMR characterization of silazanes is typically diffi-

    cult, since rather broad signals appear in the NMR spectra of

    the crosslinked polymers as well as the amorphous materi-

    als. In the present study we used the peak assignment

    reported in literature [1722]. The 29Si-spectrum of the

    crosslinked HVNG polymer cured at 3008C shows threesharp peaks (Fig. 3(a)). In contrast to the corresponding

    solution spectrum of this precursor, a new signal at

    2.5 ppm appears, whereas the peak at 33.5 ppm andthe high intensity of the signal at 20 ppm refer tounreacted vinyl and SiH groups, respectively. Resonances

    having a 29Si-chemical shift in the range of 2.5 ppmcorrespond to silicon atoms on (N)2Si(C

    sp3)2 sites and

    denote crosslinking via hydrosilylation [22]. An exact

    assignment of the signals appearing after crosslinking can

    only be made via comparison with other SiCN precursors

    like HPS. This spectrum shows only two peaks (Fig. 3(b)).

    The chemical shift with high intensity at 3.5 ppm isrelated to the hydrosilylation reaction, while the higher

    intensity compared to the signal at 22 ppm (unreactedSiH groups) and the fact that a peak for Si-vinyl groups in

    the range of about 15 ppm cannot be observed, indicatethat all vinyl groups reacted via hydrosilylation and/or

    polymerization. In the 29Si spectrum of HVNG at 5008C,

    Fig. 2. FTIR spectra (coupled with TG) of gaseous species which escaped

    during pyrolysis (HVNG).

    Fig. 3. 29Si NMR spectra of (a) HVNG-and (b) HPS-derived powder samples heat treated at various temperatures.

    58 G. Ziegler et al. / Materials Chemistry and Physics 61 (1999) 5563

  • the peak corresponding to unreacted vinyl groups has dis-

    appeared, indicating that crosslinking is completed. More-

    over, these signals are broadened (higher degree of

    crosslinking) and shifted to higher fields, owing to the

    degradation of organic groups and the enrichment of SiN

    surroundings. The latter is in agreement with TG-FTIR

    measurements (Fig. 2), where the evolution of methane

    was observed. The same effects were detected in the

    5008C 29Si NMR spectrum of the HPS precursor.Based on TG analysis, a second mass loss of about 3 wt%

    was detected in the temperature range between 8008C and14008C in conjunction with a density increase from 2.3 to2.6 g/cm3. The accompanying 29Si-NMR spectrum indi-

    cates rearrangements in the amorphous state, whereas at

    10008C only one broad peak was monitored, which corre-sponds to a homogeneous amorphous SiCN matrix. Anneal-

    ing at 15008C, however, leads to a heterogeneous SiCNmaterial. The broad peak finally separates into the three

    signals for SiC4, SiN3C and SiN4 [18,2024]. Longer

    annealing times (48 h) at 15008C cause the formation ofthe thermodynamically stable crystalline phases SiC

    (16 ppm) and Si3N4 (48 ppm). At 16008C, no Si3N4but only a SiC signal is detected by 29Si NMR measure-

    ments, which narrows at 17008C indicating crystal growthof SiC. Upon crystallization, the density increases from

    about 2.6 to 3.25 g/cm3 (corresponding to SiC) with a

    substantial mass loss of 26 wt% due to the decomposition

    of amorphous SiCN and Si3N4 by nitrogen evaporation.

    3.2. Microstructure

    The materials investigated exhibited a residual open

    porosity of about l5 vol% after high-temperature annealing

    at 15408C for 6 h in N2-atmosphere and can, therefore, beconsidered as open systems that allow for the escape of

    gaseous species formed during pyrolysis. This open porosity

    in turn affects the high-temperature stability of these poly-

    mer-derived glasses, as will be discussed in the following.

    On the other hand, since the materials revealed a high degree

    of coalescence between the powder particles and the binder

    phase upon pyrolysis, as can be seen at the fracture surfaces

    of the HVNG- (Fig. 4(a)) and the HPS-derived (Fig. 4(b))

    bulk glasses after annealing at 15408C, the materials locallycontain regions without residual porosity which is consid-

    ered here as the corresponding closed systems. Apart

    from the porosity present, the matrix of the polymer-derived

    materials revealed a homogeneous glass-like fracture sur-

    face, as shown in the SEM micrographs of Fig. 4. Distinc-

    tion between former powder particles and binder phase is

    not possible. This coalescence between powder particles,

    pre-heat treated below 6008C, and the binder phase uponpyrolysis indicates the possibility of structural rearrange-

    ments within these polymer-derived compounds, due to the

    presence of various functionalities.

    The pore sizes of the HPS-derived material are in the

    range of 13 mm in diameter, whereas the pore diameters ofthe HVNG-derived material are much larger, up to 10 mm indiameter. A second major difference, besides the pore size,

    was the occurrence of crystallized spherical inclusions

    commonly observed in the HVNG-derived glass, as

    depicted in the TEM micrograph of Fig. 5(a). These sphe-

    rical inclusions, only observed in the HVNG material,

    contained the thermodynamically stable crystalline phases

    Si3N4, SiC and graphite (compare the HRTEM image of

    Fig. 6(a)). It should be emphasized that in order to ratio-

    nalize the observed phase assemblage, a nitrogen over-

    pressure within these globules of about four atmospheres

    is required. Except of these spherical inclusions, the bulk

    material of the HVNG material remained completely amor-

    phous.

    In contrast, using the HPS-polymer for preparation of the

    monolithic SiCN-glass sample, no spherical inclusions

    could be found. The material appeared completely homo-

    geneous and amorphous, when employing SEM as the

    characterization tool (Fig. 4(b)). Additional TEM investi-

    gations, however, revealed large a-Si3N4 crystallites within

    Fig. 4. SEM micrographs of fracture surfaces of (a) HVNG- and (b) HPS-derived ceramic monoliths annealed at 15408C, 6 h, N2 atmosphere, containing3008C polymer powder particles. A distinction between former powder particles and void filling binder phase is not feasible. Note the different pore size ofthe two materials which, however, contain the same overall porosity.

    G. Ziegler et al. / Materials Chemistry and Physics 61 (1999) 5563 59

  • the glass after exposure to 15408C (see Fig. 5(b)). More-over, employing HRTEM imaging, it could be revealed that

    the bulk of the HPS-derived material was in fact not

    completely amorphous as suggested by SEM, but showed

    in some areas the formation of SiC nuclei, shown in the

    HRTEM image of Fig. 6(b).

    The characterization of the two different precursor mate-

    rials clearly revealed different crystallization phenomena. In

    the one case, spherical precipitates, filled with SiC, Si3N4and graphite, were observed while the HPS-derived sample

    revealed large a-Si3N4 crystallites besides a small numberof globules which only contained SiC. However, the actual

    reason for this marked difference in high-temperature

    response, i.e., the respective crystallization behavior, is

    not yet unequivocally known and a generalized discussion

    proved to be rather complex, as will be shown in Section 4.

    3.3. Properties

    Density and open porosity of the infiltrated and pyrolyzed

    HVNG bodies changed from 1.68 to 2.05 g/cm3 and from

    25% to 8%, respectively, after four infiltration cycles. The

    effort to decrease the residual porosity below 8%, using a

    higher number of infiltration cycles was not successful,

    because all the accessible pore channels were already closed

    after four infiltration cycles. This leads to small porosity

    gradient within the sample with a rather dense outer rim and

    a porous inner core structure.

    The Vickers hardness strongly depends on the annealing

    temperature of the material (Fig. 7), whereas different pre-

    cursors show only a small variance in hardness. The Vickers

    hardness increases between 10008C and 12008C from 7.9 to12.8 GPa and from there on remains constant up to 15008C.Annealing the specimen at 15508C, however, leads to apronounced decrease of the hardness to about 5.5 GPa, since

    crystallization occurs which creates additional porosity due

    to the strong density change.

    Fracture strength shown in Fig. 8 depends on the overall

    microstructure and on the silazane used [25]. Due to process

    optimization by die pressing at 1408C, large structuraldefects could mainly be eliminated which resulted in higher

    strength values. Therefore, the fracture strength improved

    from an average value of about 104 MPa (pyrolysis tem-

    perature of 14008C) to about 130 MPa. Monolithic samples

    Fig. 5. TEM micrographs of (a) one spherical inclusion observed in the HVNG- derived SiCN-material, (b) a-Si3N4-crystallites within the matrix of theHPS-derived material after annealing at 15408C, 6 h, N2 atmosphere.

    Fig. 6. HRTEM micrographs of (a) one spherical inclusion in HVNG revealing the crystalline phases a-Si3N4, SiC and C, and (b) the matrix of the HPS glassafter annealing at 15408C, 6 h, N2-atmosphere. Note that the formation of small SiC nuclei was observed in some regions within the glass structure.

    60 G. Ziegler et al. / Materials Chemistry and Physics 61 (1999) 5563

  • prepared from HPS showed the highest strength values with

    a maximum of 235 MPa. The Youngs modulus of the

    HVNG material also increased from 109 to 118 GPa by

    employing the warm die-pressing technique.

    The oxidation resistance of the monolithic SiCN ceramic

    was tested by isothermal oxidation in air. In general, the

    SiCN materials are stable due to the formation of a SiO2-

    protection layer. Commonly, porous non-oxide ceramics

    oxidize by internal and external oxidation (e.g., RBSN),

    whereas internal oxidation dominates at lower temperatures

    and larger channel radii. The complete mass gain of the

    pyrolyzed HVNG ceramic was not larger than 1% at 14008Cafter 72 h oxidation. Increasing the oxidation temperature

    leads to an increase in mass gain, as given in Fig. 9. The

    oxidation behavior of bulk material derived from the HPS-

    precursor differs strongly from the HVNG material. In this

    case, the total mass gain after isothermal treatment for 72 h

    at 14008C was 0.07% and, hence, about two orders ofmagnitude lower as compared to the HVNG material.

    4. Discussion

    The general idea of using polymer powders (crosslinked

    at 3008C) derived from polymers with different basic struc-tural units was based on the assumption that the micro-

    structure development and the respective thermal stability of

    these polymer-derived materials is directly influenced by

    the polymer architecture. It was suggested that the devi-

    trification of polymer-derived SiCN glasses can be

    described by a stepwise change of the microstructure,

    initiated by a rearrangement of the polymer network upon

    heat treatment, which yields phase separation within the

    amorphous state. The phase separation and, consequently,

    the thermal stability of the glass structure can therefore be

    controlled by the architecture of the starting polymer.

    However, the resulting pore structure (open/closed system)

    can also strongly affect the stability of the amorphous

    ceramic. Both precursors studied here yielded a homoge-

    neous microstructure after pyrolysis, where coalescence

    between powder particles and polymer binder had occurred.

    However, employing the HPS-precursor for preparation of

    the SiCN glass, a microstructure with a high amount of

    small pores was observed, whereas the HVNG-precursor

    resulted in a microstructure with only a small amount of

    much larger pores. The crystallization phenomena, in par-

    ticular, the occurrence of all the stable phases of the SiCN

    system within one spherical inclusion in the HVNG-derived

    material, supports the assumption that this reflects the

    crystallization behavior of a closed system. In contrast,

    the HPS system can be considered as an open system, since

    the regions without any porosity between the pore network

    are very small. It becomes evident, that no crystallization

    areas, containing all the thermodynamically stable phases

    SiC, Si3N4 and graphite, can be found within the HPS-

    derived material since the open system allows for the escape

    of nitrogen during rearrangement and decomposition of the

    amorphous SiCN network prior to crystallization. This

    results in SiC enriched areas and the formation of SiC

    Fig. 7. Vickers hardness (HV10) of different precursors (HVNG, HPS)

    annealed at various pyrolysis temperatures (6 h, N2).

    Fig. 8. Four-point fracture strength of monolithic HVNG samples

    prepared at different forming and pyrolysis temperatures. Note that one

    fracture strength data point of the HPS material is also shown (triangle),

    giving the highest strength value.

    Fig. 9. Mass change of two SiCN ceramics (HVNG, HPS) due to the

    formation of a protective silica layer by isothermal oxidation in air at

    different temperatures.

    G. Ziegler et al. / Materials Chemistry and Physics 61 (1999) 5563 61

  • nuclei within the matrix. In addition, it is assumed that the

    large a-Si3N4 crystallites, observed in the HPS material,were formed in proximity to closed pores, where the gen-

    eration of a sufficiently high nitrogen partial pressure, which

    allows for the formation and stabilization of Si3N4, was

    enabled. It is important to note, that, apart from the polymer

    architecture, the residual porosity plays a dominant role

    with respect to crystallization of the bulk polymer-derived

    ceramics.

    The resistance of SiCN ceramics against oxidation is

    affected by both, temperature and precursor type. Since

    processing of the investigated silazanes resulted in different

    microstructures of the pyrolyzed monoliths, giving different

    pore structures, it is assumed that the detected higher

    oxidation rates of the HVNG specimens are in fact a result

    of the much wider pore channels. The isothermal oxidation

    experiments imply that above 12008C and after an initialoxidation stage, the pore channels are mostly closed by a

    protective silica scale which prevents further oxidation.

    Oxidation of the HPS-precursor material, however, shows

    nearly no detectable mass change due to the smaller pore

    size of the material. The above-mentioned results again

    emphasize the effect of the residual porosity and the diffi-

    culty to distinguish between the influence of polymer

    architecture itself and the given pore structure. The latter

    is also thought to affect fracture strength obtained as well as

    KIc and hardness.

    5. Conclusions

    One of the key topics when employing newly developed

    polymer-derived glasses, is their stability at high tempera-

    tures, in particular, the stability of the amorphous state.

    Crystallization of bulk SiCN glasses is controlled by a

    stepwise change in microstructure, which is thought to

    yield phase separation within the amorphous phase,

    structural rearrangement as well as chemical degradation.

    The crystallization strongly depends on the polymer

    architecture and on the residual porosity of the system.

    One question that remains to be solved in this context is,

    if the process of phase separation is required for crystal-

    lization to occur and, therefore, would be responsible for

    the observed degradation in thermal stability. Or, on the

    other hand, if the residual porosity, i.e., the pore size and

    the ratio between open and closed porosity, in fact over-

    rules the influence of the polymer architecture and

    chemistry.

    Acknowledgements

    The authors would like to thank the Volkswagenstiftung

    Hannover and the Deutsche Forschungsgemeinschaft

    (DFG) Bonn for financial support throughout the work.

    References

    [1] T. Vaahs, M. Bruck, W.D.G. Wocker, Polymer-derived silicon

    nitride and silicon carbonitride fibres, Adv. Mater. 4(3) (1992)

    224226.

    [2] M.R. Mucalo, N.B. Milestone, I.G. Vickridge, M.V. Swain,

    Preparation of ceramic coatings from pre-ceramic precursors, J.

    Mater. Sci. 29 (1994) 44874499.

    [3] J. Lucke, M. Keuthen, G. Ziegler, Infiltration and pyrolisis behaviour

    of polymer derived UD reinforced CMCs, in: A. Bellosi (Ed.),

    Proceedings of the Fourth Euro-Geramics, vol. 4, Faenza Editrice,

    Italy, 1995, pp. 187192.

    [4] W. Verbeek, G. Winter, German Patent 2 236 078 (1974).

    [5] S. Yajima, K. Okamura, J. Hayashi, M. Omori, Synthesis of

    continuous SiC fibres with high tensile strength, J. Am. Ceram.

    Soc. 59(7)(8) (1976) 324327.

    [6] M. Birot, J.-P. Pillot, J. Dunogue`s, Comprehensive chemistry of

    polycarbosilanes, polysilazanes, and polycarbosilazanes as precur-

    sors of ceramics, Chem. Rev. 95(5) (1995) 14431478.

    [7] J. Lucke, J. Hacker, D. Suttor, G. Ziegler, Synthesis and

    characterization of silazane-based polymers as precursors for

    ceramic-matrix composites, Appl. Organomet. Chem. 11 (1997)

    181194.

    [8] R. Riedel, J. Bill, A. Kienzle, Boron modified inorganic polymers

    precursors for the synthesis of multicomponent ceramics, Appl.

    Organomet. Chem. 10 (1996) 241.

    [9] O. Delverdier, M. Monthioux, A. Oberlin, A. Lavedrine, D. Bahloul,

    P. Goursat, Thermal behaviour of polymer-derived ceramics II: Si

    CN system from a new PVSZ precursor, High Temp. Chem.

    Processes 1 (1992) 139149.

    [10] M. Monthioux, O. Delverdier, Thermal behavior of (organosilicon)

    polymer-derived ceramics. V: Main facts and trends, J. Europ.

    Ceram. 16 (1996) 721737.

    [11] H.-J. Kleebe, Microstructure and stability of polymer-derived

    ceramics; the SiCN system, Phys. Stat. Sol. (a) 166 (297) (1998).

    [12] H.-J. Kleebe, D. Suttor, H. Muller, G. Ziegler, Decomposition/

    crystallization of polymer-derived SiCN ceramics studied by

    scanning- and transmission electron microscopy, J. Am. Ceram. Soc.

    81 (1998) 29712977.

    [13] J. Bill, F. Aldinger, Precursor-derived covalent ceramics, Adv. Mater.

    7(9) (1995) 775787.

    [14] L. An, R. Riedel, C. Konetschny, H.-J. Kleebe, R. Raj, Newtonian

    viscosity of amorphous silicon carbonitride at high temperature, J.

    Am. Ceram. Soc. 81 (1998) 13491352.

    [15] R. Riedel, M. Seher, J. Mayer, D.V. Szabo, Polymer-derived Si-based

    bulk ceramics, Part I: Preparation processing and properties, J. Eur.

    Ceram. Soc. 15 (1995) 703715.

    [16] R. Riedel, H.-J. Kleebe, H. Schonfelder, F. Aldinger, A covalent

    micro/nano-composite resistant to high-temperature oxidation,

    Nature 374 (1995) 526528.

    [17] C. Gerardin, F. Taulelle, J. Livage, Pyrolysis of a polyvinylsilazane,

    polymeric precursor for silicon carbonitride: structural investigation

    by 1H;13 C;29 Si;15 N and 14N nuclear magnetic resonance, J. Chim.

    Phys. 89 (1992) 461467.

    [18] R.H. Lewis, G.E. Maciel, Magnetic resonance characterization of

    solid-state intermediates in the generation of ceramics by pyrolysis

    of hydridopolysilazane, J. Mater. Sci. 30 (1995) 50205030.

    [19] N. Brodie, J.-P. Majoral, J.-P. Disson, An NMR study of the step by

    step pyrolysis of a polysilazane precursor of silicon nitride, Inorg.

    Chem. 32 (1993) 46464649.

    [20] G.R. Hatfield, K.R. Garduner, Review solid state NMR: applica-

    tions in high performance ceramics, J. Mater. Sci. 24 (1989) 4209

    4219.

    [21] D.C. Apperley, R.K. Harris, G.L. Marshall, D.P. Thompson, Nuclear

    magnetic resonance studies of silicon carbide polytypes, J. Am.

    Ceram. Soc. 74(4) (1991) 777782.

    62 G. Ziegler et al. / Materials Chemistry and Physics 61 (1999) 5563

  • [22] J. Seitz, J. Bill, N. Egger, F. Aldinger, Structural investigations of Si/

    C/N-ceramics from polysilazane precursors by nuclear magnetic

    resonance, J. Eur. Ceram. Soc. 16 (1996) 885981.

    [23] K.R. Carduner, C.S. Blackwell, W.B. Hammond, F. Reidinger, G.R.

    Hatfield, 29Si NMR characterization of a-and b-silicon nitride, J.Am. Ceram. Soc. 112 (1990) 46764679.

    [24] J.S. Hartman, M.F. Richardson, B.L. Sheriff, B.G. Winsborrow,

    Magic angle spinning NMR studies of silicon carbide: polytypes,

    impurities, and highly inefficient spinlattice relaxation, J. Am.

    Chem. Soc. 109 (1987) 60596067.

    [25] R.J. Roark, Formulas for Stress and Strain, McGraw-Hill, New York,

    1965.

    G. Ziegler et al. / Materials Chemistry and Physics 61 (1999) 5563 63

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