Precursor-Derived Ceramics∗

Chapter 11.1.10

Precursor-Derived Ceramics
Markus Weinmann,y Emanuel Ionescu,

Ralf Riedel

and Fritz Aldingerz y

H.C. Starck GmbH, Im Schleeke 78-91, Goslar, Germany,

Technische Universita¨t Darmstadt, Institut fu¨r Materialwissenschaft, Petersenstrasse 23, Darmstadt, Germany, z Max-Planck-Institut fu¨r Metallforschung, Heisenbergstrasse 3, Stuttgart, Germany

Chapter Outline 1. 2. 3. 4.

Introduction Precursor Synthesis Polymer-to-Ceramic Transformation High Temperature Properties

1025 1026 1050 1065

1. INTRODUCTION The condensation of organometallic compounds into merely inorganic materials by a proper thermal treatment under controlled atmosphere is a unique and fairly simple process of producing new types of ceramics, which ared due to their origindreferred to as precursor- or polymerderived ceramics (PDCs). This procedure allows for an easy control of elemental composition, chemical homogeneity, and material architecture on an atomic scale. It provides an inimitable access for controlling and adjusting the design and the microstructure of ceramic materials that cannot be achieved using conventional processing techniques such as melting or sintering. In recent years, this process scheme has therefore created substantial interest, both scientifically and for practical purposes. The synthesis and processing of PDCs usually do not require any additives. PDCs thus display the intrinsic chemical and physicalechemical properties of the constituting pure phases, for example, an exceptional oxidation and corrosion stability as well as crystallization and creep resistance up to very high temperatures. Multinary PDCs frequently possess even superior properties compared with their constituting binary or ternary subsystems. Several review articles [1e5] and books [6,7] have been published on this topic, as well as special issues in the Journal of the American Ceramic Society




Dedicated to Prof. Dr. Klaus Mu¨ller, who passed away far too young.

Handbook of Advanced Ceramics. http://dx.doi.org/10.1016/B978-0-12-385469-8.00056-3 Copyright Ó 2013 Elsevier Inc. All rights reserved.

5. Applications 6. Conclusions References

1080 1092 1092

(“Ultrahigh-Temperature Polymer Derived Ceramics”) [8], Journal of the European Ceramic Society (“Polymerderived Ceramics”) [9], and Soft Materials (“Preceramic Polymers”) [10], which focused on the synthesis and properties of PDCs. The overall process for the formation of PDCs is outlined in Figure 1. This consists of three major steps. At first, suitable monomers are transformed into polymeric macromolecules referred to as precursors, which are subsequently crosslinked at moderate temperatures into preceramic networks

FIGURE 1 Flow diagram for the preparation of PDCs.

1025

1026

with defined rheological behavior to provide proper processing capabilities. The networks are finally transformed into inorganic materials by a thermal heat treatment at temperatures usually ranging from 1000 to 1400  C. The general idea behind the process concept is to generate preferred structural features in the organometallic precursors (this can mostly be realized by comparably simple preparative organometallic chemistry) and to subsequently transform the precursors into ceramics under retention of the especially designed building blocks. Thermolysis must thereby include a series of controllable condensation steps, which are aimed for the specific architecture of the desired ceramic material. As a consequence of being mostly reaction controlled, thermolysis is especially valuable for the densification of covalently bonded inorganics. Conventionally, low diffusion capability of nonoxide ceramics even at very high temperatures is compensated by the use of sintering aids. Such additives are mostly of the oxide type and therefore degrade the otherwise unique properties of nonoxide ceramics such as thermal, chemical, and mechanical stability over wide temperature ranges. In contrast, thermolysis of ceramic precursors provides a means to realize the production of nonoxide ceramic thereby keeping up their excellent properties without compromises. One remarkable consequence of the low atomic mobility is that as-thermolyzed PDCs can be formed which remain amorphous up to rather high temperatures. The reason for this is that substantial thermal activation is required for nucleation, grain growth, and the formation of crystalline phases. Crystallization thereby very much depends on the chemical/elemental composition, the atomic structure and the microstructure of the materials, and the atmosphere applied. Under certain conditions, crystallization can be retarded even up to 1800  C providing novel glass-like structural materials with superb thermomechanical properties. However, also crystalline materials are of great interest because their microstructure formation can be controlled during devitrification to a large extent providing, for example, a means for stabilizing nanosized morphologies. This chapter intends to review the fundamentals of synthetic approaches and the processing of silicon-based PDCs, to reveal the possibilities of microstructure development and to give hints on their potential application. It has been shown that the basis for the microstructure design of PDCs is the type of macromolecule, that is, its chemical composition and molecular structure. Therefore, at first synthetic procedures, which deliver different types of organometallic, silicon-based ceramic precursors are evaluated. In the following, thermolysis reactions will be considered, because the structure of the ceramic materials is also widely determined by the structure of the molecular

Handbook of Advanced Ceramics

precursors. Further on, phase reactions controlling the high temperature stability of these materials will be discussed and finally recent trends in technologically relevant applications be reported.

2. PRECURSOR SYNTHESIS 2.1. General Comments The most intensively investigated preceramic polymers for the preparation of PDCs are polysilanes, polycarbosilanes (PCSs), polysiloxanes, as well as polysilazanes and polysilylcarbodiimides [11e13]. Since the Yajima process to synthesize silicon carbide fibers using PCSs, significant development in the synthesis and processing of PDCs has been achieved. Thus, using silicon-based polymers, technologically important ceramic components such as complex-shaped monoliths, fibers, coatings or infiltrated porous media, and powders can be prepared. The molecular structure and the type of preceramic polymer influence not only the chemical composition but also the microstructure and the phase composition of the final ceramic [14]. In this way, the chemical and physical properties of PDCs can be varied and adjusted to a great extent by the design of the molecular precursor. Thus, the synthesis of preceramic polymers is one of the key issues in the PDC field.

2.2. Precursors to Silicon Carbide Silicon carbide is the only (solid) binary compound of silicon and carbon. It is used for multifold purposes. The most important feature is its outstanding mechanical behavior. Silicon carbide is one of the most frequently used ceramic hard materials [15]. Silicon carbide has an extremely low self-diffusion coefficient, even at a very high temperature. Processing, that is, compaction of silicon carbide powders thus requires sintering aids, which maydto a certain extentd unfavorably determine the properties of the processed ceramic materials [16e19]. In contrast, precursor thermolysis of suitable organosilicon compounds is a method, which delivers sinter aidfree silicon carbide. As will be shown in the next sections, suitable precursors for this purpose are, for example, polysilanes, PCSs, polyalkenylsilanes, polyalkynylsilanes, or their mixtures (copolymers or polymer blends). In the following, the most important of the abovementioned organosilicon polymers will be reviewed chronologically, starting with polysilanes, which are composed of alternating [SiR2] units, followed by polysilaethylenes, which are composed of [SiR2eCR2] building blocks. Finally, PCSs with alternating [SieCeC] units will be discussed.

Chapter | 11.1.10

Precursor-Derived Ceramics

1027

2.2.1. Polysilanes Polysilanes are well-studied polymers [20e22]. However, although many research groups all over the world have extensively investigated synthetic approaches, especially with the goal to obtaining polysilanes as precursors to phase-pure silicon carbide, there are only a very limited number of procedures described in the literature that deliver near-stoichiometric silicon carbide. Polysilanes, also referred to as poly(silylene)s, usually have a linear polymer backbone consisting entirely of silicon atoms to which two organic substituents R are bonded. When sufficiently purified, poly(diorganosilane)s, [SiR1R2]n (R1, R2 ¼ singly bonded organic group) are stable in oxidative and hydrolytic environments and thus can be handled easily without requiring inert gas techniques. Polysilanes are often insoluble, not meltable, and intractable, which impedes their applicability in many practical processes. Chemical and physicochemical properties are strongly influenced by the molecular weight and the nature of the silicon-bonded substituents. Polysilanes were first synthesized in the early 1920s by Kipping and Sands [20] by a WurtzeFittig type dehalogenation reaction of dichlorodiphenylsilane with molten potassium in xylene delivering poly(diphenylsilane). Ph Cl

Si

Cl

Ph

Si

(1)

Ph

n

Ph = C 6H5

In 1949, Burkhardt published the first clear description of the synthesis of a polysilane with a potential for technological application [23]. Dichlorodimethylsilane was reacted with sodium in a WurtzeFittig type reaction in benzene in analogy to the abovementioned procedure. CH 3 Si

CH 3 Cl

Na, benzene

Si

- NaCl CH 3

R

R H

Si

H

Catalyst - H2

Si H

H

n

(3)

Catalysts: η 2-Alkynyl-Titanocenes, η2-Alkynyl-Zirconocenes

Ph K, xylene - KCl

Cl

processing (purification) including hydrolysis to remove excess alkali metal. A further essential drawback of the Wurtz coupling is the limited functional group tolerance, which allows only for the synthesis of polysilanes with more or less “inert” side groups such as alkyl or aryl groups. The mechanism of the Wurtz-type coupling of chlorosilanes seems to be very complex. Alkali metal surface reactions are probable, but dehalogenation reactions in solution have also to be considered. Moreover, the mechanisms involving silyl radicals are discussed as well as mechanisms based on silyl anions, which appear to be key species in the chain growth of polysilanes. An alternative procedure for the synthesis of polysilanes, which delivers products in high purity and higher yields, is the dehydrocoupling of silanes in the presence of suitable catalysts such as h2-alkynyl titanocene or -zirconocene as described by Chang and Corey in 1989 [24].

(2)

CH 3 n

Poly(dimethysilane) is a white powder, which is poorly soluble in organic solvents and which does not melt without decomposition. It will be highlighted below that it represents the first precursor for the production of refractory silicon carbide fibers. In the meantime, a large number of polysilanes with different silicon-bonded substituents and molecular weight have been synthesized. Product yields are frequently unsatisfactory because of (i) the low solubility of the polymers in organic solvents and (ii) the expandable

In contrast to the WurtzeFittig reaction, no solid byproducts are formed. The only byproduct is hydrogen. As a consequence, purification steps such as hydrolysis and filtration are not required. The authors also suggested a conclusive mechanism of the silane polymerization by dehydrocoupling. By elimination of the side-on coordinated alkynyl ligand, a highly reactive 14VE (VE ¼ valence electron) metallocene species forms, which inserts into a SieH bond of the silane, thus forming a silylmetallocene. This species most likely eliminates hydrogen thereby forming a metallasilanediyl intermediate and subsequently adding a silane molecule by hydrosilylation of the Si]M bond. The disilylated metallocene rearranges by a hydrogen shift to the transition metal center and reductive elimination of the oligomer/polymer (Scheme 1). Woo et al. who investigated dehydrocoupling reactions of silanes using zirconocene and hafnocene hydride complexes as catalysts postulated alternative mechanisms [25]. In contrast to the above, they suggested a mechanism that is based on two s-bond metathesis reactions that pass through four-center transition states (Scheme 2).

2.2.2. Polycarbosilanes In contrast to polysilanes, polymer backbones in PCSs are not exclusively composed of silicon atoms. Here, aliphatic

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Handbook of Advanced Ceramics

SiMe3

R

R

Cp2M

H

SiMe3

H

Si

H

H

n+1

[M]

H

Si

H

H

n

H3SiR - Me3SiC CSiMe3

R Si

R

H

n

H

Cp2M - Cp2M

Si

R

H H H Si Si R R H

H(SiHR)n-1

H

Si

H(SiHR)n-1

Si(R)H2 H

R

R

H(SiHR)n-1 R H

H

Si

H

H3SiR

Si

Si

H

SCHEME 1 Suggested mechanism for the formation of polysilanes by metallocene-catalyzed dehydrogenative coupling of alkylsilanes, RSiH3, via metalla-silanediyl intermediates [24].

units link the silicon atoms. Frequently, also unsaturated units as well as aromatic ring systems serve as building blocks for the formation of PCSs. PCSs of the general structure [R2SieCH2]n have been synthesized by a variety of reaction pathways. The most important and intensely investigated access to PCSs certainly is the Yajima process [26]. This process gained much attractiveness as it thermally converts insoluble and not meltable poly(dimethylsilane) into soluble and processable poly(methylsilyleneemethylene), [HSi(CH3)eCH2]n. This precursor can be either cast or spun. Silicon carbide fibers obtained from this precursor are commercially available under the trade names NicalonÔ or High NicalonÔ, depending on the processing applied (cf. Section 5.5)

[M] = CpCp*(Cl)Hf Cp = 5-C5H5, Cp* =

5

-C5Me5

the polymer backbone into the polymer chain. The process is radical induced. The mechanism including the single key steps in the Kumada rearrangement of poly(dimethylsilane) is shown in Scheme 3. Initially, SieSi bonds in the polymer are homolytically cleaved thus forming silyl radicals, followed by H-radical migration from a methyl group to the silyl radical (initiation). The methylene radical that is thereby generated inserts into an SieSi bond whereby a new silyl radical is formed (rearrangement). Propagation proceeds by a repeated H-radical migration to the silicon radical with the formation of a methylene radical. Termination most probably takes place by radical recombination. PCSs with C1 building blocks can also be obtained by Grignard reactions of chloromethyl trichlorosilane Cl3SieCH2Cl in diethyl ether solution [27]. Subsequent

CH3

CH3

CH3

CH3

- H2

SCHEME 2 Proposed mechanism for the formation of polysilanes by hafnocene-catalyzed dehydrogenative coupling of alkylsilanes, RSiH3, via metalla-silacyclobutane transition state [25].

-C5H5, M = Ti, Zr, R = CH3, C6H5

Cl

R

H

H

5

Si

Si [M]

"Hydrosilylation"

H

Cl

H

[M]

H

H Metalla-silanediyl

R Cp2M

Si

H2 Cp2M Si

H

H

R

[M]

H

Cp2M

Cp =

H

Na -NaCl

400ºC

Si

Si CH2 H

CH3 n

n

(4)

KumadaRearrangement

The key step in the Yajima process is a rearrangement of the polymer backbone involving pendent methyl groups, referred to as Kumada rearrangement or Kumada reaction, which occurs at around 400  C. It is a methylene migration from one of the silicon-bonded methyl groups attached to

treatment of the as-obtained poly(dichlorosilylene-methylene) with LiAlH4 delivers branched PCS with an idealized structure [H2SieCH2]n. Due to its compositional analogy to polyethylene, it is also referred to as polysilaethylene. Thermolysis of polysilaethylene delivers near-stoichiometric SiC.

Chapter | 11.1.10

Precursor-Derived Ceramics

1029

Initiation Me Me Si

Me 400ºC

Si

Si

2

Me

Me Me Me Me Si

Si

+

Me Me

Me

Me CH 2

Si

Si

Me

Me Me

Si

Me Si

+

H

Me

Rearrangement Me CH 2

Me

Si

Si CH 2 Si

Si

Me

Me Me Propagation Me Si CH 2 Si Me

+

Me

Me

Me Me

Me

Si

Si

Si

Me Me

Me CH 2

H CH 2 Si

Me

Si

+

Si

Me Me

Me Monomer Unit

SCHEME 3 Key steps in the Kumada rearrangement of poly(dimethylsilane) [26].

Cl

Cl Cl

Si

Mg, Et2O

CH2Cl

H

Si CH2

- MgCl

Cl

Cl

R1 Si

CH2Cl

R2

2

R

- MgCl2

1

R

R Si

Si

Pt-catalyst

Si CH2

1

R

R2

n

2

R , R = H, CH3, Cl

ROH

1. Mg 2. CH3COCl / FeCl 3 OR

Cl

n

R1

1

b)

(5)

acid as catalysts [28]. Wu and Interrante developed a modified multistep polysilaethylene synthesis [29] starting from methyltrichlorosilane, Cl3SieCH3, which is a cheap, commercially available compound. Halogenation using chlorine in the presence of ultraviolet light delivers chloromethyl trichlorosilane, Cl3SieCH2Cl (this reaction is not shown in Scheme 4). Cl3SieCH2Cl could not directly be converted into 1,1,3,3-tetrachloro-1,3-disilacyclobutane. Rather, two of the three silicon-bonded chlorine atoms were initially substituted by ethoxy (eOEt) groups, which

2

a) Mg

Si CH2 H

n

The disadvantages of this reaction pathway are long reaction times that are required for obtaining sufficient polymer yields. Changing the solvent toward a more polar one such as tetrahydrofuran strongly influences the selectivity of the reaction. If starting from chloromethyl dialkylor diarylchlorosilanes, according to reaction pathway (a), 1,3-disilacyclobutane derivatives are obtained. The transformation into linear polymers can be achieved by ring-opening polymerization (ROP) reactions using Pt-group metal complexes, such as chloroplatinic

Cl

LiAlH4

Si

CH2Cl

OR a) R1, R2 = H, alkyl, aryl b) R1, R2 = Cl

c) R1, R2 = Cl

LiAlH4

H Si CH2 H

n

SCHEME 4 Synthesis of PCSs by Pt-catalyzed ring-opening polymerization (ROP) of 1,3-disilacyclobutane derivatives [28,29].

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Handbook of Advanced Ceramics

could be selectively realized in alcoholysis reactions before the reductive coupling. The ethoxy groups in 1,1,3,3-tetraethoxy-1,3-disilacyclobutane were subsequently replaced with chlorine atoms by a treatment with acetic chloride acid in the presence of FeCl3. Pt-catalyzed ROP and subsequent reduction with LiAlH4 then delivered linear polysilaethylene in an approximately 60% yield. Procopio and Berry published a procedure for the synthesis of oligosilaethylene by dehydrogenative coupling reactions [30] (Scheme 5). Oligomeric carbosilanes were obtained from trimethylsilane under mild conditions in the presence of H3Ru(SiMe3)(PMe3)3, which was obtained in situ from H2Ru(PMe3)4 and HSiMe3. The proposed mechanism involves the formation of a disilyl ruthenium complex, which upon a b-hydride shift transforms into a h2-silylene species (in analogy to ethylene, this is the monomer unit of poly(silylene-methylene)). Insertion of the silylene into the remaining RueSi bond results in the formation of a singly silylated ruthenium complex and chain growth by one SiMe2CH2 unit. PCSs with aliphatic C2 building blocks in the polymer backbone are precursors to SiC, which have reached remarkable attractiveness. Most of the publications in this field trace back to work of Corriu et al. who have applied hydrosilylation of vinylhydridosilanes, (R1)(R2)SiH(HC] CH2) (R1, R2 ¼ H, Cl, alkyl, aryl, NR2), for the preparation of linear PCSs of the general type [(R1)(R2)SieC2H4]n [31,32]. Even though this type of hydrosilylation does not

occur regioselectively, formation of the b-addition product (C2H4]CH2CH2) is strongly preferred. 2

R H 1

R

HSiMe3

- H2

LnRu SiMe3 H LnRu

SiMe2 CH2

H

Si

R2 C C R3

H2PtCl6

R

n

H2 Me2 Si CH2 LnRu H SiMe3

SiMe3

1

2

HSiMe3

(6)

Si n

Polymerization by hydrosilylation of vinylsilanes requires a catalyst such as H2PtCl6 dissolved in 2-propanol, also known as Speyers’ reagent [33]. In 2-propanol solu2 tion, PtCl2 6 is reduced to PtCl4 , which is the catalytically active species. The mechanism which involves 16VE and 18VE species is shown in Scheme 6. Initially, one of the chloride ligands in the 16VE fragment PtCl2 4 is replaced with a vinylhydridosilane ligand, thus forming a h2-olefin PtCl 3 complex. In a second step, a vinylhydridosilane molecule adds oxidatively (insertion of [Pt] into SieH) whereby an octahedral 18 VE complex forms, followed by an insertion of the h2-coordinated olefin into a platinumehydrogen sbond. The vacancy, which forms upon this migration, is occupied by an external chloride ligand that adds to the complex. Finally, the carbosilane is eliminated reductively with regeneration of the catalytic species PtCl2 4 . Similarly, PtCl2 4 -catalyzed hydrosilylation of alkynylhydridosilanes of the general type (R1)(R2) SiH(C^CR3) is a synthetic pathway for obtaining linear PCSs in which alternating silicon atoms and olefinic C2 units constitute the polymer backbone [34,35].

Si CH2 Me

H

R

R1

Me H3Ru(SiMe3)(PMe3)3

H2PtCl6 2

R

H2Ru(PMe3)4 + HSiMe3 - PMe3

Si

H C

Si

C

R1

R3 n

(7)

R1, R2 = H

Formation of poly(alkenylsilanes) of this type is only observed if R1, R2 s H. The use of dihydridoalkynylsilanes (R)SiH2(C^CH) as starting compounds, in contrast, delivers highly crosslinked, glass-like PCSs. The reason is that the olefin units, which are generated by the hydrosilylation of the alkynyl units, may react further and transform into aliphatic C2-building blocks.

H LnRu

SiMe2

-hydride shift

CH2

Me3Si 2

-silene

HH

R

Ln = (PMe3)3H2

SCHEME 5 Proposed mechanism for the formation of oligosilaethylene by H2Ru(PMe3)4-catalyzed dehydrogenative coupling of trimethylsilane, Me3SiH via a h2-silene ruthenium complex [30].

H

Si H

R C C H

H2PtCl6

H C

Si

C

(8) n

further cross-linking

Chapter | 11.1.10

Precursor-Derived Ceramics

1031

Mechanism: H2PtCl6 1

2

R

R

i

Si

Si

H

PrOH H

2

1

R

R

[PtCl4]

21

Si

2

R

R

16VE Red. Elim.

-

Cl 2-

Cl Cl Pt

H 1

Cl

Cl

Si

R

2

R

Si

Cl

1

R

2

1

R

R

Cl Pt

H Si

Cl 2

R

H Cl

-

Cl

1

Cl 1

R

2

R

Cl

Ox. Add.

Pt Si

Si

H

1

R

H

R

Si

2

R

2

R

18VE SCHEME 6 Proposed mechanism for the formation of PCSs by PtCl2 4 catalyzed hydrosilylation of vinylsilanes [33].

PCSs of the general type [(R1)(R2)SieC^C]n, in which silicon atoms are linked via C^C building blocks, are not accessible by hydrosilylation reactions. In contrast, metathesis reactions of dichlorosilanes with dialkalimetal acetylides or the respective Grignard reagents deliver this type of polymers in good yields [34].

Cross-linked Products

R1 n Cl

Si

referred to as poly[(silylene)diacetylenes] or poly[(silylene) diethynylenes]) obtained from 1,4-dilithio-1,3-butadiyne proved to be precursors with interpenetrating networks to silicon carbide/metal carbide (M ¼ Ti, Nb, Ta) nanocomposites [36]. Preparation of branched polymers by regiospecific hydrosilylation of poly[(silylene)diethynylenes] was also reported [37].

R1 Cl

MC CM - MCl

R2

Si

C

R2 M = Li, Na, BrMg

The alkynyl unit represents a highly reactive site within the polymer backbone. This is an important issue with respect to obtaining highly crosslinked precursors for ceramics to prevent polymer degradation and thus volatilization of low molecular weight species during thermolysis. Moreover, the alkynyl group can be functionalized by, for example, hydroboration, hydrosilylation, or attachment of transition metal complex fragments by coordination, providing an access for the synthesis of multicomponent ceramics. In addition, poly(butadiynylsilanes) [(R1)(R2)SieC^CeC^C]n (also

C n

Functionalization by - Hydrosilylation - Hydroboration - Complex formation

(9)

2.3. Precursors to Silicon Nitride For the preparation of silicon nitride-based PDCs oligo- or polysilazanes, [(R1)(R2)SieN(R3)]n (R ¼ H, alkyl, aryl, etc.) or poly(silylcarbodiimide)s, [(R1)(R2)SieN]C]N]n have been studied intensively during the last 30 years, resulting in an immense number of structurally different precursors that were published. In general, such precursors deliver silicon nitride/carbon or silicon nitride/silicon carbide/carbon composites if R1, R2, and/or R3 contain

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Handbook of Advanced Ceramics

carbon atoms. The formation of phase-pure silicon nitride can only be realized by thermolysis of perhydridopolysilazane (PHPS) in which R1, R2, and R3 consist of hydrogen and/or nitrogen or by performing the polymerto-ceramic conversion in an ammonia atmosphere. The latter, however, is limited to powders or parts with comparably small dimensions and/or cross sections.

2.3.1. Carbon-free Polysilazanes

Cl Si

Cl + 6n NH 3

Cl

Si H

NH Cl + 3 NH 3

- NH 4 Cl

NH

SiH 2

SiH 2 NH

(schematical representation)

[cat.] , - H2 Further crosslinked products

Narsavage et al. published the polymer-to-ceramic conversion of tetrakis(ethylamino) silane in an inert gas atmosphere (Ar), from which they expected to obtain phase-pure silicon nitride [38]. The idea behind this concept was that silicon is exclusively bonded to nitrogen. It was expected that Si3N4 formation by volatilization of the nitrogen-bonded ethyl groups takes place during thermolysis. However, carbon was not completely degassed, and the resultant material was an amorphous silicon carbonitride. Consequently, the only remaining possibility for obtaining carbon-free Si3N4 is thermolysis of polysilazanes, containing only Si, N, and H. The first attempts for the preparation of such polymers were reported already in 1885 by Schutzenberger and Colson [39], by ammonolysis of tetrachlorosilane according to Eqn (10). Polymer formation was accompanied by precipitation of ammonium chloride as a coproduct, which was removed by filtration. Glemser and Naumann performed more detailed investigations on this system [40]. n Cl

SiH 2

H

[Si(NH) 2]n + 4n NH 4Cl

Cl

(10) Silicon diimide, Si(NH)2, is a colorless solid, which is insoluble and infusible. With increasing temperature, it decomposes with the elimination of ammonia to form amorphous Si3N4 at 1000  C. Further heat treatment to 1400e1500  C results in the crystallization of a-Si3N4. A processable Si3N4 precursor, PHPS, can be obtained by the ammonolysis of dichlorosilane in polar solvents such as diethyl ether or dichloromethane as shown in Scheme 7 [41]. It is composed of linear and cyclic motifs but rapidly ages by loosing hydrogen and further crosslinking, thus increasing its viscosity gradually from becoming oily over waxy to glassy. Thermolysis to 1050  C releases a ceramic material in 70% yield, which according to X-ray investigations consists of a-Si3N4, b-Si3N4, and elemental silicon. Excess Si could be removed almost completely by performing thermolysis in an ammonia atmosphere. However, a critical issue with respect to transferring this process into a technical scale is the use of H2SiCl2, which is a highly

SCHEME 7 Synthesis of processable precursors to phase-pure silicon nitride by ammonolysis of dichlorosilane and subsequent crosslinking [41].

flammable gas, which can disproportionate with the formation of SiH4 and SiCl4. A modified process, allowing for a safer handling of the highly reactive starting compounds was developed at the Tonen Company. Before ammonolysis, H2SiCl2 was modified by reaction with pyridine, which resulted in the formation of H2SiCl2
(NC5H5)2 [42]. Blanchard and Schwab modified the process developed by Seyferth et al. They avoided the formation of free silicon by synthesizing a PHPS, which delivers stoichiometric silicon nitride even after thermolysis in an inert gas atmosphere, using coammonolysis of dichlorosilane and trichlorosilane [43].

2.4. Precursors to Ternary SieCeN Ceramics 2.4.1. Carbon-containing Polysilazanes While the synthesis of PHPS is limited to a small number of suitable starting compounds such as H2SiCl2, HSiCl3, SiCl4, and NH3 or H2NeNH2 as silicon and nitrogen sources, respectively, the chemistry of carbon-containing polysilazanes of the general type [(R1)(R2)SieN(R3)]n (R1, R2, R3 ¼ H, alkyl, aryl etc.) is a very multifaceted field. The first publications in this regard were published by Rochow appeared already in 1966 [44]. In contrast to H2SiCl2 and HSiCl3, functionalized chlorosilanes of the general type (R1)(R2)SiCl2 (R1, R2 ¼ H, alkyl, aryl, halide etc.), which are in general used in the synthesis of SieCeN polymers, are mostly cheap and commercially available. Therefore, much effort has been made within the last 30 years toward specifically designing polysilazanes for numerous applications. In contrast to PHPS that deliver amorphous SieN ceramics, carbon-containing polysilazanes release amorphous SieCeN ceramics after thermolysis to approximately 1100  C. Interestingly, these materials possess improved high temperature mechanical properties compared to silicon nitride derived from PHPS. A particularly critical issue is “free,” that is, excess silicon. The reason is that silicon segregation is known to lower the crystallization temperature of silicon-nitride-based materials drastically [45]. In carbon-containing SieCeN materials,

Chapter | 11.1.10

Precursor-Derived Ceramics

1033

which usually contain “free” carbon, silicon segregation is suppressed because of the formation of silicon carbide. There have been a remarkable number of synthetic pathways described in the literature, which deliver polysilazanes. The desired application, the choice of the chemical composition and molecular structure of the monomers, reaction conditions and eventually further crosslinking finally determine the synthetic approach to the precursors. The procedure used most frequently for the synthesis of polysilazanes is ammonolysis or aminolysis of chlorosilanes. Ammonolysis of chlorosilanes in general proceeds according to the following: R4-xSiClx +

3 x NH3 2

1

n

[R4-xSi(NH)x/2]n + x NH4Cl

x = 1-3

(11) The single steps are (i) initial substitution of Cl atoms with NH2 units followed by (ii) a condensation step upon which SieNHeSi linkages are built. Depending on the number of chlorine atoms per silicon atom and the nature of the silicon-bonded groups R, monomeric, dimeric, cyclic, short chain linear, or highly crosslinked polysilazanes form. Ammonolysis of monochlorosilanes R3SiCl results in monomeric species. In the case of “small” silicon-bonded substituents R (R ¼ H, CH3), disilazanes R3SieNHe SiR3 are formed. Certainly, if silicon-bonded substituents are sufficiently small and reaction temperatures are high, these species condense further, thereby transforming into tertiary amines bearing Si3N motifs [46]. Against this, bulky substituents R, such as isopropyl, isobutyl, or mesityl inhibit the abovementioned condensation step due to steric effects, thus causing the formation of monomer silylamines R3SiNH2 [47,48]. Hexamethyldisilazane (HMDS), (H3C)3SieNHeSi (CH3)3, synthesized by the ammonolysis of trimethylchlorosilane on a large scale, is industrially by far the most important monomeric silazane. Because it is a volatile liquid (boiling point: 125  C), it cannot be directly used for the solid-state thermolysis process. However, it is an important single-source precursor for the preparation of SieCeN coatings by the chemical vapor deposition (CVD) processes [49] and a very important source for the synthesis of polysilazanes by trans-amination (redistribution, metathesis) reactions [50]. Ammonolysis of dichlorosilanes R2SiCl2 delivers oligo- or polysilazanes of the general type [R2SieNH]n, presuming sufficiently high reaction temperatures and sterically less demanding silicon-bonded substituents R such as hydrogen, methyl, ethyl, or phenyl. Reaction of ammonia with dichlorosilanes having sterically demanding

substituents R such as i-propyl, t-butyl, or mesityl delivers silyldiamines, R2Si(NH2)2 according to the abovedescribed ammonolysis of monochlorosilanes R3SiCl. Usually, product mixtures are obtained, which consist of 6 and 8-membered rings (trimers and tetramers, respectively) besides low molecular (M < 2000) chain molecules. The low molecular weights of the tri- and tetrameric species result in low boiling temperatures and cause their volatilization during thermolysis, whereas the chain molecules frequently degrade during the heat treatment with volatilization of low molecular gas species [51]. As a consequence, ceramic yields of polysilazanes are frequently unsatisfactory. Seyferth and Wiseman published a method for increasing the ceramic yields of cyclic silazanes bearing SieH and NeH units considerably by crosslinking with catalytic amounts of KH [52,53]. In the case of cyclotetra(methylsilazane), the ceramic yield could be increased from <30% to approximately 84%. The authors proposed a mechanism, which involves silyleneeimine motifs as key species [53] (Scheme 8). It is supposed that the SieH units of a second Si4N4 ring add intermolecularly in a hydrosilylation type reaction to the highly reactive silyleneeimine intermediate, which could not be isolated. The as-formed species can again be deprotonated and react further. However, there are also mechanisms desirable excluding silyleneeimine formation, which were postulated for the dehydrocoupling of boron-modified silanes using n-butyl lithium as a catalyst [54]. Here, n-butyl lithium initially deprotonates the silazane N-atoms resulting in highly nucleophilic amides. Subsequent intermolecular substitution of a silicon-bonded hydride by an amide results in the formation of a new SieN bond. The hydride released deprotonates an NeH unit, thereby closing the catalytic cycle. A possibility for the synthesis of highly branched high ceramic yield polysilazanes is the attachment of reactive sites, for example, vinyl groups to the polymer backbone. These allow for further crosslinking reactions during thermolysis and thus very efficiently avoid polymer backbone degradation. Vinyl-substituted oligo- or polysilazanes can be obtained by ammonolysis of vinyldichlorosilanes (H2C] CH)Si(R)Cl2 (R ¼ H, CH3) according to the following. H C Cl

Si

CH 2

H C 1 n

Cl + 3 NH3

R

CH 2

Si

N

R

H

+ 2 NH4Cl n

R = H, CH 3

(12) Depending on the silicon-bonded unit R and reaction conditions applied, [(H2C]CH)Si(R)eNH]n can crosslink

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H3C H N Si

H

H

Si

H3C

H

NH

Si

NH

H

[KH] CH 3

Si

HN

H3C H N Si

H

Si

H3C

- H2

N Si

HN Si

H

CH 3

NH

H CH 3

CH 3

H3 C H H N Si Si NH H3 C Si CH 3 HN Si NH H H CH 3 H

H3C H N Si

H

Further crosslinked product

H3C

Si

H

H3C

SiH Si

NH

H

H

Si

N

HN

H3C H N Si

NH

Si

H

CH 3

Si

HN

NH

H

CH 3

H

CH 3

SCHEME 8 Crosslinking of cyclotetra(methylsilazane) with catalytic amounts of KH. By increasing the crosslinking density, ceramic yields could be increased from <30% to >80% [52,53].

by three different reaction pathways (cf. Scheme 9): (i) olefin polymerization, (ii) hydrosilylation, and (iii) dehydrocoupling (the latter reactions only take place if R ¼ H). It should be noted that both olefin polymerization and hydrosilylation are addition reactions, which take place without the formation of byproducts. This is an important issue that makes vinylsubstituted polysilazanes candidates for the production of polymer-derived ceramic matrix composites [55].

Nuclear magnetic resonance (NMR) spectroscopy and chemical analysis suggest that hydrosilylation is the main reaction in the thermally induced crosslinking of [(H2C] CH)SiHeNH]n [56]. Addition of catalytic amounts of Speyers’ reagent can accelerate it. Thermally induced trans-amination (not shown in Scheme 9) in contrast, which would be accompanied with ammonia evaporation, occurs to a minor extent. By adding catalytic amounts of

Olefin Polymerization

CH

CH2

Hydrosilylation

CH H C

Si

N

R

H

n

R = H, CH3

Si

CH2

Si

N

R

H

ion addit

C Si

n

R=H n

addit

ion

R=H - H2

H

N H

H2C Si

Dehydrocoupling

CH3

CH2 N n

SCHEME 9 Possible crosslinking reactions of PVSs.

CH2 N H

n

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1035

potassium hydride, dehydrocoupling reactions become favored. Crosslinking in this case proceeds by SieN bond formation with the elimination of hydrogen. Possible mechanisms for catalysis, hydrosilylation, and dehydrocoupling were already mentioned above and are not discussed here in detail. In contrast, potassium hydride promoted anionic polyaddition reaction of the vinyl groups (anionic olefin polymerization) was not observed. Due to the absence of SieH motifs in [(H2C]CH) Si(CH3)eNH]n, neither hydrosilylation nor dehydrocoupling can be considered for thermal crosslinking. However, based on the structural characterization by solid-state 13Cand 29Si-magic angle spinning Nuclear Magnetic Resonance (MAS NMR), Fourier Transform Infrared (FTIR), and elemental analysis, Bill et al. concluded that between 250  C and 350  C, which is approximately 120  C higher than the temperature required for hydrosilylation reactions of [(H2C]CH)SiHeNH]n, the vinyl groups are transformed into aliphatic hydrocarbons [57]. More detailed information on thermolysis reactions of polysilazanes are provided in Section 3.2. Very important silazane-based precursors are copolysilazanes obtained by coammonolysis or eaminolysis of chlorosilane mixtures. Polymerization of such mixtures allows for the synthesis of modified polymers with tunable properties such as solubility, rheology, and (latent) reactivity, which cannot be achieved using single-source precursors. Originally, copolysilazanes were developed to improve the processing of the precursors for silicon nitride fibers, that is, to adjust proper rheology and latent reactivity. Moreover, the importance of chemical constitution, that is, phase composition was recognized [58]. The main subject in this regard was to decrease the amount of free carbon in SieCeN ceramics to shift decomposition temperatures toward higher values. A commercially available precursor in this series is NCP200, a “polyhydridomethylsilazane,” which is obtained by the coammonolysis of dichlorodimethylsilane and H 3 Cl

Si H

dichloromethylsilane [59]. Because of commercial availability, low cost, good processability, and high ceramic yield of approximately 80% (this value, however, depends on the processing), this precursor could successfully be applied for the preparation of dense silicon carbonitride ceramics [60]. Structurally comparable is NN710 a “polyperhydridomethylsilazane,” which can be synthesized by coammonolysis of dichloromethylsilane and dichlorosilane [61]. Recently, a structurally similar precursor, which is derived from this system and which delivers stoichiometric silicon carbide/silicon nitride, was published [62]. The polymer was obtained according to Scheme 10 by coammonolysis of H2SiCl2 and H3CSiHCl2 in a 3:1 molar ratio and subsequent base-catalyzed crosslinking using nBuLi as a catalyst. While the as-obtained polymer is a viscous oil, which rapidly ages, the crosslinked precursor is a hard glass-like infusible solid, which is insoluble in all common organic solvents. In mass-spectra-coupled simultaneous thermogravimetric analysis (TGA-MS) investigations, it was shown that this type of polymer could be pyrolyzed without the formation of gaseous byproducts other than hydrogen, thus delivering amorphous SiC/Si3N4 with a ceramic yield of 94.5% (c.f. Figure 10). According to the procedures described above, ammonolysis of trichlorosilanes RSiCl3 usually results in the formation of [(R)Si(NH)1.5]n. Such poly(silsesquiazane)s (these are polymeric silazanes in which three siliconbonded substituents are nitrogen atoms and in which N:Si ¼ 1.5) were first published in 1967 by Andrianov and Kotrelev [63] who performed ammonolysis of methyltrichlorosilane. Cl Cl

Si

Cl

- 8 NH4Cl

+ 3 NH4Cl

1

/n

Si

N

H

H

3

H

1.5

n

1

/n

Si

Si

N

H

H

n

-H2

H

Thermolysis

94.5% ceramic yield

R

CH3

n-BuLi

1200°C/Ar

N

(13)

H

a-SiC/Si3N4

Si

R = CH3, C2H5

H 12 NH3

1 n

Cl + 4.5 NH3

R

CH3 Cl + Cl

Si

CH3 Si

N H

3

H

N n

(schematically) SCHEME 10 Synthesis of a precursor for stoichiometric SiC/Si3N4 [62]. A key step is subsequent crosslinking of the initially obtained precursor in order to avoid thermal degradation with volatilization of low-molecular weight species other than hydrogen.

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In contrast to poly(silsesquioxane)s, (RSiO1.5)n, which can be obtained by hydrolysis of trichloro- or trialcoxysilanes under water-starved conditions and which are structurally well investigated [64e66], there is not much known about the molecular structure of poly(silsesquiazane)s. However, similar to poly(silsesquioxane)s, they are obtained as a mixture of molecules with cage structures as well as oligomeric and highly branched polymeric products. The structure of a hexameric silsesquiazane (H3C)6Si6(NH)9, which was obtained from sodium amide and (H3C)SiCl3 in n-hexane solution at 78  C, was reported by Ra¨ke et al. Sodium amide was used instead of ammonia and thus formation of ammonium chloride precipitate avoided. Single crystal X-ray analysis displayed a cage-type structure of (H3C)6Si6(NH)9 in which two sixmembered ring systems [(H3C)Si(NH)]3 with chair conformation are linked via the silicon centers with each one NH unit [67]. Polysilazanes are also accessible by transition metalcatalyzed dehydrocoupling of silanes and ammonia or amines. In contrast to the above-described ammonolysis and aminolysis of chlorosilanes, no solid byproducts are formed in the dehydrocoupling reactions; the only byproduct is hydrogen, which evaporates during the reaction. In this regard, it has to be mentioned that dehydrocoupling of silanes and ammonia allows for the preparation of highly crosslinked polysilazanes that due to their comparably low solubility are difficult to be obtained using the conventional salt elimination process. Laine and Blum found that ruthenium catalysts such as Ru3(CO)12 are most suitable. They reported the first catalytic synthesis of polysilazanes in 1986 [68e71]. R1 H

Si R2

R1 H + NH3

[Ru3(CO)12] - H2

Si NH R2

n

(14)

An alternative method for the synthesis of polysilazanes, which also avoids the formation of solid byproducts during the polymerization step, relies on the redistribution reactions of chlorosilanes with HMDS [50]. H

R1 Cl

Si

Cl +

N Me3Si

1

SiMe3

Si NH

/n

2

R2

R

n

+ 2 Me3SiCl

(15) The reaction can be performed in solution or without solvent using excess HMDS. Both HMDS and the couple product trimethylchlorosilane are volatile liquids, which can easily be removed from the crude reaction mixture by a distillation step under reduced pressure. As pointed out above, the molecular weight of polysilazanes can be increased by a crosslinking step, which occurs either by trans-amination or dehydrocoupling reactions. In both cases, polymers are obtained with retention of the silazane ring structures. The conversion of low molecular weight cyclic silazanes, for example, cyclodi-, tri-, or tetrasilazanes, by ROP in contrast results in the loss of cyclic structural motifs. The idea behind this process is to obtain linear high molecular weight polysilazanes with proper rheological properties, as they are required for polymer fiber spinning. Ring opening of cyclosilazanes can be induced by treatment with catalytic amounts of a strong base such as MeLi, n-BuLi, or t-BuLi [73], strongly acidic compounds such as triflic acid methylester, F3CSO3CH3 [74e77], or by transition metal catalysts such as Ru3(CO)12 [68]. For example, Suom and coworkers investigated in detail F3CSO3CH3e and organo alkali salt-induced ROP of cyclodisilazanes covering small substituents in dichloromethane solution [78].

R1, R2 = alkyl

H3C

CH 3

CH 3

Si

In addition, the authors demonstrated that the transition metal complex noted above is also useful for crosslinking low-molecular-weight oligosilazanes such as trimethylcyclotrisilazane, [(H3C)SiHeNH]3, or tetramethylcyclotetrasilazane [(H3C)SiHeNH]4. Liu and Harrod reported on the dehydrocoupling of ammonia and less reactive tertiary silanes such as diphenylmethylsilane using dimethyltitanocene as a catalyst [72]. Primary silanes (PhSiH3) surprisingly reacted retarded even though the opposite was expected. As an explanation for this phenomenon the authors referred to the possibility of homocoupling reaction of the silanes (SieSi linking), resulting in disilanes or trisilanes, which possess lower reactivity toward crossdehydrocoupling with ammonia than secondary or tertiary silanes R2SiH2 or R3SiH, respectively.

R1

R

N

N Si

H3C

R

[F3CSO3CH 3] 30ºC, 30 min

Si

N

CH 3 R

CH 3

n

R = CH3, C2H5....

(16) They observed that four-membered ring systems, which have higher tension, react faster and more readily than cyclotrisilazanes, independent of basic or acidic conditions that were applied. Remarkably, polysilazanes with molecular weights of up to 18,000, using acidic conditions and up to 100,000 using basic conditions, could be obtained by this method. Another remarkable issue using ROP is the possibility for synthesizing copolymers when starting from cyclodisilazanes with differently substituted silicon atoms.

Chapter | 11.1.10

Precursor-Derived Ceramics

1037

R1 H3C

C2H3 Si

H3C

N

N

CH 3

Si

Si H3C

Cl

C2H3

CH 3 ROP

Si

N

(17)

R2

Disilylated carbodiimides are known since the early 1960s and mainly Ebsworth and Mays, as well as Pump and Wannagat investigated the synthesis and reaction behavior of these monomeric compounds [79e84]. However, because of their volatility disilylated carbodiimides cannot be used directly as precursors for ceramics. In 1964, Pump and Rochow reported for the first time on poly(silylcarbodiimide)s, which were obtained by the reaction of dichlorosilanes with disilver cyanamide [85].

Si 2

R

R1 Cl + Ag2N-CN

1

- AgCl

/n

Si

N C N

2

n

R

(18) However, for a transformation into a technical scale and the application of the polymers as a preceramic material, this approach is by far too expensive. A patent of Klebe and Murray in 1968 described a modified procedure for the synthesis of poly(silylcarbodiimide)s [86]. In a trans-silylation reaction starting from bis(triorganosilyl)carbodiimides (usually bis (tri-methylsilyl)carbodiimide) and di- or trichlorosilanes, the precursors could be obtained in simple procedures and high yields. n MeSiCl3 + 3n Me3Si-N=C=N-SiMe3 -3n Me3SiCl

n MeSi[(-N=C=N-)1.5]n

R = alkyl N C

N n

n MeSiCl3 + 3n H-O-H

Substitution

n MeSi(-N=C=N-SiMe3)3 -1.5 Me3Si-N=C=N-SiMe3

/n

Besides the simplified processingdno solid byproducts formed in this reactiondtremendously reduced costs for the precursor synthesis were also advantageous. In the beginning of the 1990s (>20 years after the appearance of the patent of Klebe and Murray!), poly (silylcarbodiimide)s were rediscovered and their applicability as preceramics studied in much detail. The synthetic approach was improved significantly by developing a nonoxide solegel process, and the molecular structure of cyclic tetra(dimethylsilylcarbodiimide) using single crystal X-ray diffraction (XRD) determined [12,87e95]. Recently, air-stable linear carbon-rich polysilylcarbodiimides [R1R2SieNCN]n with R1 ¼ phenyl, R2 ¼ Ph, Me, H, and vinyl were reported, which were synthesized via the liquid phase reaction of R1R2SiCl2 with bis(trimethylsilyl)carbodiimide. The increased carbon content of these precursors induces an increase in the carbon content in the final ceramics, resulting in an improved thermal stability of the derived SieCeN ceramics and higher crystallization temperatures. Remarkably, Si3N4 crystallization in the resulting ceramics upon annealing at high temperatures is inhibited [96]. Due to the similar electronegativity of the carbodiimide group and oxygen, polysilylcarbodiimides can be regarded as pseudo chalcogenides. Accordingly, polysilylcarbodiimide gels are non-oxide analogous to classical oxide gels, for example, to tetraethoxysilane- (TEOS) Si(OC2H5)4 or tetramethoxysilane-based (TMOS) (Si(OCH3)4) systems, with respect to structural, spectroscopic, and rheological properties [89,91,97] (Scheme 11).

2.4.2. Poly(silylcarbodiimide)s

Cl

1

- R3SiCl

(19)

R1 Si

R1

Cl + R3Si-N=C=N-SiR 3

R2

N

CH 3 CH 3 CH 3 CH 3 n

CH 3

Si

-3n HCl

n MeSi(-O-H)3 Condensation

Gel

-1.5 H-O-H

n MeSi[(-O-)1.5]n

Pyrolysis SixCyNz

SixCyOz

SCHEME 11 Nonoxide solegel process for the reaction of trichloromethylsilane with bis(trimethylsilyl)carbodiimide (left) and its analog oxide solegel type (right) [91]. For color version of this figure, the reader is referred to the online of this book.

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Handbook of Advanced Ceramics

[B(CH3)eN(C6H5)]3, were mixed and reacted at an elevated temperature. In a similar manner, the authors reacted blockcopolymers carrying Si(CH3)2 and Si(C6H5)2 units in the polymer backbone with the above-noted borazine derivative [99,100]. Such polymers were spun and rendered infusible by either subsequent curing in oxidative atmosphere or by irradiation with actinic rays. Investigations to determine the molecular structure of the precursors were not performed.

N

A

C3N4

C

B

Si3N4

SiC

2.5.1. Borazine-based SieBeCeN Precursors

Si

FIGURE 2 Ternary SieCeN phase diagram (T < 1440  C, 1 atm N2). A and B: compositions SiC2N4 and Si2CN4, respectively, the first crystalline ternary phases in the SieCeN system obtained from poly (silylcarbodiimide) [Si(NCN)2]n [95].

Thus, highly crosslinked poly(silylcarbodiimide) can be obtained via a nonoxide Sto¨ber process by reacting a diluted solution of tetrachlorosilane and bis(trimethylsilyl)carbodiimide in the presence of pyridine (pyridine was added to accelerate the trans-amination reaction) according to Eqn (20) [98]. By using this method, the formation of the first crystalline phases in the SieCeN system upon thermal treatment was observed (Figure 2) [95]. Cl

The chemistry of borazine and its derivatives is an important field in inorganic chemistry. It is thus natural that scientists deal with potential precursors to SieBeCeN ceramics long before this topic became relevant for materials science. Such an example is work published by Noeth in the early 1960s [101]. He investigated the reaction of trichloroborane or alkyldichloroborane with (trialkylsilyl) 0 dialkylamines, R3SiNR2, and HMDS. It turned out that 0 chloroboranes readily react with R3SiNR2 under SieN bond cleavage and BeN bond formation. For example, (trimethylsilyl)dimethylamine, Me3SieNMe2 reacted with trichloroborane to form dimethylaminodichloroborane and trimethylchlorosilane.

Tol, [Py] Cl

Si

Cl + 2 Me3Si-N=C=N-SiMe3

Cl 1

/n

Si N C N

- Me3SiCl

(20)

2 n

b-SiC2N4 (space group Pn3m) crystallizes >400  C. It possesses an anticuprite structure, which can be described as two interpenetrating crystobalite lattices. Above 920  C, it decomposes with the loss of cyanogene and nitrogen to form Si2CN4 (space group Aba2), which is similar to the crystalline structure of sinoite [95].

2.5. Precursors to Quaternary SieBeCeN Ceramics Organometallic polymers containing the elements silicon, boron, carbon, nitrogen, and hydrogen are a comparably young class of preceramic polymers. However, the exceptional high temperature stability of their derived ceramics, which is superior to those noted in the previous sections, coupled with a high resistance toward oxidation, that is, self-healing capabilities under oxidative conditions, has prompted considerable interest in these new materials. In 1986, Takamizawa et al. first described a method for the preparation of an inorganic fiber containing silicon, carbon, boron, and nitrogen. For this purpose, polydimethylsilane, [(H3C)2Si]n and B-trimethyl-N-triphenylborazine,

Me3Si-NMe 2 + BCl 3

Me2N-BCl 2 + Me 3SiCl

(21) Using HMDS, (Me3Si)2NH, instead of Me3SieNMe2 and alkyldichloroborane RBCl2, resulted in the formation of monomeric B-alkylated borazine derivatives. H N 3 Me3Si-NH-SiMe3 + 3 RBCl2

RB

BR

HN

NH

+ 6 Me 3SiCl

(22)

B R

Highly crosslinked borazine-based precursors to SieBeCeN ceramics with a comparably low silicon content were synthesized by Narula et al. starting from B-chloroborazine derivatives and HMDS [102]. H N

H N ClB HN B Cl

BCl

(Me 3Si)2NH

NH

-Me 3SiCl

B

H N B

HN

NMe B n

(23) Even though not evident from the molecular structure depicted above, silicon and carbon are present in the oligomer-capping groups.

Chapter | 11.1.10

Precursor-Derived Ceramics

1039

Seyferth and Plenio published further borazine-based polymers in 1990. They reported on the synthesis of “borasilazanes” as polymeric precursors for borosilicon nitride [103,104]. Initially, borane undergoes addition to the cyclic silazane with the formation of a silazaneeborane adduct (not shown) and elimination of dimethyl sulfide. Subsequent dehydrocoupling (intermediate I) and trimerization result in the formation of a borazane (proposed intermediate II in Scheme 12) in which silazane rings are maintained. A shift of a boron-bonded hydride to silicon, accompanied by an ROP reaction of the silazane motif finally releases the precursor. Blum and Laine independently published a similar procedure starting from poly-N-methylsilazane, [SiH2e N(CH3)]n, and ammonia or trimethylamine borane complexes BH3
NH3 or BH3
N(CH3)3, respectively, in the presence of catalytic amounts of Ru3(CO)12 [105,106]. They obtained viscous products, which were not characterized by spectroscopic methods in more detail. Ceramic yields were in the range of 50%. A series of SieBeCeN precursors, which were obtained by dehydrogenative, and/or dehydrosilylative coupling reactions of monomeric or oligomeric silazanes with borazine as the boron source, were published by Sneddon et al. [107e113]. For example, hydridopolysilazane (HPZ) was reported to react smoothly with liquid borazine in different stoichiometry at temperatures between 50  C and 90  C to give

borazine-substituted polysilazanes in high yields. Spectroscopic investigations indicated that the polysilazane used retained its backbone but was substituted with pendent borazine substituents by means of a borazineeboron to polymerenitrogen linking, resulting from both hydrogen and trimethylsilane elimination [107,108]. H N

Si(CH3)3 Si H

N

N

N

HN

NH

H

B H

H N HB

50-70ºC - HSi(CH3)3 - H2

BH

HN

(24)

NH B

Si

N

N

H B HN

NH

HB

BH

H

N H

However, molecular weight and distribution studies suggested an inhomogeneous distribution of borazine throughout the obtained polymers. The evidence for dehydrosilylative coupling reactions was confirmed by reaction of tris(trimethylsilylamino)silane with borazine [109,110].

H

N

[Si]

[Si]

3 H

BH

+

H

H [Si]

HB

[Si]

+ 3 H3B * S(CH3)2

N H

N

[Si]

3 N

-H2, - S(CH3)2

H

N

[Si]

BH2

proposed intermediate (I) [Si] = SiHMe Trimerization H H B

[Si] H–([Si]NH)

N

N

2

HB

[Si]

[Si] NH[Si]

BH

H

N -

H -Shift, Si-N bond clevage

N [Si] NH[Si]

2

2

H

N

H [Si]

HH

[Si]

N

[Si]

B [Si]

N

N

[Si]

N H

B B H H H [Si] N [Si] H

H

N H

[Si]

N H

proposed intermediate (II) SCHEME 12 Synthesis of borazine-based polyborosilazanes via dehydrogenative coupling of borane dimethyl sulfide and cyclic silazanes [103,104].

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Handbook of Advanced Ceramics

H Si

NH

HB

BH

HN

NH B H

SiMe3

SiMe3

50-70ºC

+

NH

NH Me3Si

- HSiMe 3

H B Si

BH

B N H

N

NH Me3Si

(25) NH

HN

H NH

A crucial issue in the synthesis of borazine-based polymers is the lack of selectivity in the polymerization reactions. This fact causes difficulties in controlling the structure and properties of the precursors and the ceramic materials. Especially if a precursor is designed for drawing fibers, latent reactivity mustdthough it is absolutely required for obtaining high ceramic yields and rendering green fibers infusibledbe adjusted very carefully. The pendent borazine units in the above-noted precursors possess both BeH and NeH functions, which upon a thermal treatment may react further and increase molecular weight and viscosity of the polymeric precursors. Consequently, they are not suitable for processes that require stable melt viscosities, as, for example, in the generation of polymer fibers by melt spinning [113]. As a possible approach out of this dilemma, Sneddon et al. designed monofunctional and difunctional borazine derivatives such as B-diethylborazine (DEB) or B-monoethylborazine, which can be prepared by metal-catalyzed

H N

SiMe3

H

This finding was traced back to the observation that after vacuum fractionation of the volatile polymerization byproducts 0.25 mol of trimethylsilane per mole of borazine could be isolated. However, polymerization again occurred due to BeHeNeH dehydrocoupling reactions. Moreover, the formation of borazylene motifs was discussed, and the following idealized product structure was suggested [109].

H B

H B

HN

NH

B

Si

B N H

N Me3Si

2

B

B N

H N + 2

NH

HN

Si Me2

HN

N

90ºC

HB

BH

- H2

N H H B N

HN Me2 Si N

N

Si Me2 N

(27)

B

B N H

SiMe2

B N

HN

BH

HB N H

n

Si N

Me3Si

BH

B N SiMe3

N H

(26) 2

H

H B SiMe2

N H

N H

Depending on the silazane to borazine stoichiometry, ceramic yields between 38 and 42% were obtained. Treatment of borazine with cyclic silazanes such as hexamethylcyclotrisilazane also seems to be complex. The main reaction however is a dehydrocoupling of borazineeboron atoms with the nitrogen atoms of the silazanes according to Eqn (27), whereas ring-opening reactions with SieN bond cleavage, which may also appear is less likely. Thermolysis up to 1400  C delivers ceramics with 55e62% yield, depending on the borazine to silazane ratio used.

Me2Si

B

B N

HN

H

N

HN

H

NH

HN H B

H B

hydroboration of ethylene by borazine [111]. In contrast to “pure” borazine, these species have only one or respectively two reactive BeH units. The modification by dehydrocoupling of polysilazanes such as HPZ with DEB resulted in a precursor with pendent borazine units (Scheme 13). As a side reaction SieN bond cleaving of HPZ with the formation of DEBeNHSiMe3 was observed. However, the polymeric precursor, which had a sufficient glass transition temperature and an onset of weight loss >200  C, could be melt spun to SieBeCeNeH green fibers which were subsequently transformed into ceramic fibers by thermolysis up to 1400  C. Borazine was also used in the hydroboration of polyalkinylsilanes. Therefore, x equivalents (referred to the monomer unit; x ¼ 1, 2, 3) of poly[1,2-ethynediyl-(methylsilylene)] (PEMS, [HSi(CH3)C^C]n) were mixed with borazine. The mixture was stepwise heated in a Tefloncoated steel autoclave to 200  C [114]. A catalytic hydroboration at an ambient temperature as reported by Evans and Fu [115] was not successful. The polymers were obtained as large redebrown glassy grains. Their color brightened with increasing borazine content. Characterization of the insoluble products was

Chapter | 11.1.10

Precursor-Derived Ceramics

1041

Et

Et

H Si

H Si

N

H N

N

B

H

H

H 60-100ºC

+

Si(CH3)3

NH

B

NH

B

HN

Et B

HN

B

NH

- H2

B H

Si

N

H HN

B

Si

Et

N Si(CH3)3

NH

HPZ B

B N H

Et

Et

side reaction Et H Si

N

H

H

Si

HN

H H + (H3C)3Si

N

B

B HN

NH B Et

SCHEME 13 PHPS [111].

Synthesis of borazine-based SieBeCeN precursors by dehydrogenative coupling of monofunctional borazine derivatives DEB and

performed by IR and solid-state NMR spectroscopy as well as elemental analyses. The latter revealed a loss of borazine in the order of 10e20%, which presumably evaporated while manipulating the precursor mixture. The atomic ratios Si:C ¼ 1:3 and B:N ¼ 1:1 of the starting compounds were maintained. Thus, depolymerization or borazine cleavage did not occur. Hydroboration is expected to appear on C^C bonds of the initial polymer and C]C double bonds resulting after the first hydroboration step. As a side reaction, the C^C unit may also be transformed into an olefinic C]C or aliphatic CeC unit by single or double hydrosilylation, respectively (not shown).

(28)

Dehydrocoupling of SieH with NeH as shown in Eqn (29) is another desirable reaction pathway, which has to be considered.

(29)

IR spectra revealed the presence of SieH n(SieH) and borazine or polyborazylene units (n(NeH), n(BeH), and n(BeN)). The relative intensity of borazine absorption bands increased with increasing borazine content of the materials. Absorption bands of CeH vibrations of sphybridized carbon atoms (terminal C^CeH groups) as present in the starting material PEMS were not detected. Likewise, there was no unambiguous evidence for the formation of C]C double bonds. Solid-state NMR spectroscopy did not give a clear indication about the favored reaction pathway. 1H NMR spectra displayed broad signals without much structural information. 13C NMR spectra of polymers obtained from mixtures with x ¼ 1, 2 exhibited resonances at 114 ppm pointing to a considerable amount of C]C units. Additional broad resonance between 10 and 40 ppm were assigned to carbon atoms with mixed chemical composition (CHSiBC, CH2SiC, CH2BC). The 29Si NMR resonance signals lie within the range expected for the structural components evolving from the proposed reaction pathways (cf. Eqns 28 and 29).

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Handbook of Advanced Ceramics

11

B NMR spectra were identical for all of the three samples. They exhibited line shapes typical for boron atoms in borazine or polyborazylene with a quadrupolar coupling constant CQ ¼ 2.9 MHz and an asymmetry parameter h ¼ 0 [116]. In summary, the spectroscopic results indicated that reaction between PEMS and borazine most probably proceeded via single or double addition of borazine to the C^C bond with the formation of both C]C and/or CeC units. In contrast, there is no hint for depolymerization reactions of the starting compound or ring-opening reactions of borazine.

2.5.2. SieBeCeN Precursors from Ammonolysis and Aminolysis Reactions of Boron-modified Chlorosilanes Totally different approaches to Si-BeCeN ceramic precursors, which do not make use of the borazine derivatives and whichdwith one exceptionddo not produce borazine motifs during the polymerization, step were published by Baldus and Jansen et al. [117e121] and by Riedel et al. [122e126]. Baldus and Jansen reported on ammonolysis and aminolysis reactions of the single-source precursor trichlorosilyl-amino-dichloroborane (TADB) [127], which they obtained starting from HMDS in a two-step reaction with (i) tetrachlorosilane and (ii) trichloroborane. H N Cl 3Si

BCl 2

+ H3C-NH 2

TADB CH 3

Si H3C N Si N N

B H

H3C CH 3 N N Si Si B B N N N N H H CH 3 CH 3

(30)

N

CH 3

IR and NMR data of the meltable and soluble polyborosilazane, which was obtained in 80% yield, as well as elemental analysis, suggested the above polymer structure in which Si3(NCH3)3 six-membered ring systems are connected via HNeBe and N(CH3)B-units. Because of the lack of resonance signals at approximately 30 ppm in the 11B NMR spectrum, the existence of borazine-like structures was excluded. The as-obtained polymer could be transformed into SieBeCeNeH green fibers, which upon thermolysis delivered ceramic fibers with extraordinary high temperature properties (cf. Section 5.5).

FIGURE 3 Crystal structure of ClMe2SieNHeBCl2 (DADB). Two molecules are associated via intermolecular BeN coordination, thus forming planar B2N2 rings [128]. For color version of this figure, the reader is referred to the online version of this book.

Accordingly, derivatives Cl2MeSieNHeBCl2 (methyldichlorosilylaminodichloroborane, MADB) and ClMe2Sie (dimethylchlorosilylaminodichloroborane, NHeBCl2 DADB) as novel processable single-source precursors of amorphous SieBeCeN ceramics have been published [128]. MeSiCl3 and Me2SiCl2 were used as starting compounds instead of SiCl4. Interestingly, DADB crystallizes with two molecules associated via intermolecular BeN coordination, thus forming planar B2N2 rings (Figure 3). MADB and DADB were each polymerized by both ammonolysis and aminolysis using methylamine, releasing more or less highly crosslinked polyborosilazanes. The degree of polymerization was determined by NMR spectroscopy. The results suggested that methyl groups, either present in the molecular precursor (SieMe) or the nitrogen component (NeMe) inhibit quantitative condensation of the initially formed SieNHR (R ¼ H, Me) units. As a consequence, ammonolysis of MADB released a more highly crosslinked precursor compared with that obtained by aminolysis of DADB with H2NMe. The latter was a lowmolecular-weight polyborosilazane with a substantial amount of SieNHMe end groups. Hydroboration of (H2C]CH)(CH3)SiCl2 with borane Lewis base adducts such as borane dimethyl sulfide, H3B
SMe2, results in the formation of tris(methyldichlorosilyl-ethyl)borane [129,130]. Subsequent ammonolysis as reported by Riedel et al. results in SieBeCeN ceramic precursors, which are composed of silazane chains that are crosslinked via CeBeC bridges.

Chapter | 11.1.10

Precursor-Derived Ceramics

1043

An approach that allows for increasing the ceramic yields in the above systems starts from vinyldichlorosilane or vinyltrichlorosilane instead of methyldichlorosilane [133e137].

R H C Cl

CH 2

Si

B BH3* SMe2

Cl

CH 3

CH

R Cl

Si

Cl

- Me2S

- NH 4Cl

CH 3

R'

CH 3 H

R

R = C2H4Si(CH 3)Cl 2

B Si

Cl

Si

CH

R' BH3* SMe2

Cl

CH 3

B

CH 2

C

CH 3

CH

R

NH 3

Cl

Si

NH 3

Cl

- Me2S

- NH 4Cl

R

R

R = H, Cl

R = H, Cl R' = C 2H4Si(R)Cl 2

N

CH 3 H

n

R'

R = [C2H4Si(CH 3)NH]

(31) A crucial subject is the filtration process, which has to be applied for removing the solid coproduct ammonium chloride. This separation can be extremely time intensive, because the precipitate partly appears colloidally distributed within the product mixture. Moreover, ceramic yields are comparably low at around 50%. It will be shown below that structurally similar precursors can be obtained by a simple modification of the reaction coordinate. Using H2BCl
SMe2 or HBCl2
SMe2 as hydroboration reagents instead of H3B
SMe2, results in the formation of bis(methyldichlorosilyl-ethyl)chloroborane and (methyldichlorosilyl-ethyl)dichloroborane [130], which after ammonolysis, release more highly crosslinked precursors, in which the B:Si ratio increases from 1:2 to 1:1 compared to 1:3 in BH3-derived precursors. Ceramic yields increase in the same order from 50% over 56% to 76%. Interestingly, polymers derived from the dichloroborane derivative H3CSiCl2eC2H4eBCl2
SMe2 possess molecular structures that are composed of six- and eight-membered silazane ring systems, SieBeN heterocycles and borazine units [131,132].

N

R

H

In contrast to the methyl-substituted derivative, these species provide latent reactivity, which enables further crosslinking during thermolysis. As a result, ceramic yields are >80%. However, a remarkable disadvantage is the decreased solubility of these precursors compared with the methyl derivative. Consequently, the processing is difficult and polymer yields are low. In the case of the highly crosslinked precursor with R ¼ (NH)0.5, the polymer yield is even <5%, which is a result of its high crosslinking density along with its low solubility. Both polymer yields and processing were significantly improved by switching the above reaction sequence, that is, by first performing ammonolysis followed by hydroboration [133e137].

Me CH Si

B Me NH HN

Si Me

cyclosilazane

H

CH 2

Si

C NH 3

Cl

Si Me N N B

NH B

(32)

B R´

CH

N

R

H

BH3* SMe2 - Me2S n

R = H, (NH) 0.5,CH3

R = H, Cl, CH 3 R´

Me Si

CH 2

Si

- NH 4Cl

R HN

H N

n

(33)

Cl

- NH4Cl - SMe2

Me

HN

Si

NH3

Cl

Me Si

CH 3

R = H, (NH) 0.5 R' = [C2H4Si(R)NH]

C

SMe2 Me B CH Cl Si

CH

H

Cl

Cl

B R'

CH 3

CH Me

NH B SiBN heterocycle

CH

Me borazine

Si

N

R

H

n

R´= C 2H4Si(R)NH R = H, (NH) 0.5, CH3,

= further cross-linking

(34)

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Handbook of Advanced Ceramics

R' CH3

B R'

1. LiAlH 4 2. H3B * SMe2

Cl Si Cl

CH Si

H

H

R R R = H, CH3 R' = C2H4Si(R)H 2

R = Cl, CH3

n NH3, [ BuLi] toluene/thf, 70ºC

H2NCH 3, thf, 60ºC

A

C

B

n H2NCH 3, [ BuLi] toluene/thf, 70ºC

R'

R'

B R'

CH

CH3

Si

N

R

H

B

R'

R'

B CH

R'

CH

CH3

CH3

Si

N

R

CH3 n

n

Si

N

R' = C2H4Si(R)NH R = (NH) 0,5, CH3

R

CH3 n

R' = C2H4Si(R)NCH 3 R = (NCH 3)0,5, CH3

R' = C2H4Si(R)NCH 3 R = (NCH 3)0,5, CH3 SCHEME 14

Dehydrocoupling reactions of tris(hydridosilylethyl)boranes with ammonia and methylamine [54,140,141].

2.5.3. SieBeCeN Precursors from Boron-modified Hydridosilanes

Ammonolysis of vinyl-substituted chlorosilanes (H2C]CH)Si(R)Cl2 (R ¼ H, Cl, CH3) delivers the respective oligo- or poly(vinylsilazane)s (PVSs), which are colorless liquids. Subsequent hydroboration then releases the polymeric precursors in quantitative yields. Whereas the methyl- and the hydrogen-substituted polymers possess sufficient solubility to remain in solution, the more highly crosslinked precursor [B[C2H4Si(NH)1.5]3]n precipitated as a fine-grained solid from the reaction during borane addition. Remarkably, the workup of SieBeCeN precursors obtained in this way is rather simple and only requires the removal of all volatile components by vacuum evaporation. According to this procedure, boron-modified polysilazanes with different boron contents were synthesized and the high temperature properties of their derived materials evaluated (cf. Section 4.4, Figure 26) [138,139].

An alternative method for the formation of SieBeCeN polymers as precursors to ultrahigh temperature ceramics is a dehydrocoupling of ammonia or alkyl amines with tris(hydridosilylethyl)boranes, B[C2H4Si(CH3)3nHn]3 (n ¼ 1e3; C2H4 ¼ CH2CH2, CHCH3) [54,140,141] (Scheme 14). It was observed that the ammonolysis of B[C2H4Si(CH3)3nHn]3 (route A) does not occur directly and requires basic catalysts such as n-butyl lithium similar to a method described by Seyferth and Wiseman in which potassium hydride was used for the crosslinking of cyclic silazanes [52]. A possible mechanism for this base-catalyzed dehydrocoupling was given [54] (Scheme 14).

n

BuLi + NH 3 n - BuH

1

B H

LiNH 2

H2

2

1 NH3

Si

H

R 2x or 3x B

LiH H2N

Si

B NH2

R R = NH2, CH3 proposed intermediate

- NH3 3

Si

N

R

H

n

= C2H4

SCHEME 15 Proposed mechanism for the formation of boron-modified polysilazanes by base-catalyzed dehydrogenative coupling of tris(hydridosilylethyl)boranes with ammonia [54].

Chapter | 11.1.10

Precursor-Derived Ceramics

1045

It was suspected that n-butyl lithium initially deprotonates ammonia with the formation of lithium amide and evaporation of n-butane (1). The more nucleophilic amide then replaces a silicon-bonded hydride which subsequently deprotonates ammonia with the evolution of molecular hydrogen (2 / 1). The proposed silyldiamine or -triamine was not stable under the reaction conditions applied, and by fast elimination of ammonia (3), the polymeric precursors were formed. A precursor system that was designed especially for the production of fiber-reinforced high temperature SieBeCeN ceramic matrix composites was also published [142e144]. Glass-like polymeric precursors were obtained by a thermally induced hydrosilylation reaction of oligo(vinylsilazane) (OVS) using the above-noted tris(hydridosilylethyl)boranes. R´ H 3

B

CH2

C Si

N

H

H

+n

H

n

Si

2.5.4. SieBeCeN Precursors Based on Poly(silylcarbodiimides) According to the procedure described for the synthesis of poly(silylcarbodiimide)s as precursors to ternary SieCeN ceramics described in Section 2.4.2, their boron-modified counterparts are precursors to SieBeCeN materials. They are accessible by different methods: first, by the hydroboration of vinyl-substituted poly(silylcarbodiimide)s with borane dimethyl sulfide. The polymeric starting compounds can be obtained either from the respective vinyl-substituted chlorosilanes with cyanamide, H2NeCN, in the presence of pyridine or from vinyl-substituted chlorosilanes and bis(trimethylsilyl)carbodiimide, Me3SieN]C]NeSiMe3 [146e148].

CH3 H

R'

120-200ºC

C Si

R' = C2H4SiH2R

B CH R Si

B

CH 2

H

R

R´ R´

CH

R´

thermolysis conversion yields/low shrinkage during thermolysis. For further details, see Section 5.6.

CH

R' BH3* SMe2

N C N

R

Si

CH 3 N C N

- SMe2 R

n

CH3

n

R´ = C2H4Si(R)NCN

R = H, CH3

H

R = H, CH 3, (NCN) 0,5

(36) H2C

CH2

Si

N

H

H

n

R' = C2H4SiHRC2H4SiHNH

(35) Remarkably, neither solvents nor catalysts were required for this process nor were byproducts formed. For the crosslinking step, the starting compounds were mixed at room temperature, and the resulting colorless reaction mixture was carefully degassed in a moderate vacuum. During a stepwise heating of the mixtures to 120  C within 2 h and then to 180e200  C within 4e5 h, a remarkable viscosity increase was observed, which resulted in the formation of the transparent, colorless, and glass-like solid. Preliminary experiments have proven the applicability of the silane/silazane mixture for the preparation of fibere matrix composites in the resin transfer molding (RTM, cf. Chapter 5.6) process [145]. Prerequisites for this process, which the precursor mixture fulfills, are, for example, low viscosity (sufficient fluidity) of the starting compounds to allow for sufficient penetration, adequate vapor pressure, appropriate chemistry (no solvents or byproducts), controllable reactivity during crosslinking, and high

IR and NMR results pointed out that boron is attached exclusively to the silicon-bonded vinyl units. An addition of boron to the carbodiimide unit, as was reported for alkyl alumination reactions of monomeric silylcarbodiimides, was not observed [149]. Alternatively, boron-modified poly(silylcarbodiimide)s were obtained in a nonoxide solegel process involving tris(chlorosilylethyl)boranes, which were reacted with excess bis(trimethylsilyl)carbodiimide [146e148]. R' R' n

B Cl

R' CH Si

CH3 Cl

R'

B

exc. BTSC

Si

- (H3C)3SiCl

R

CH

CH3 N C N n

R

R´ = C2H4Si(R)Cl2

R´ = [C2H4Si(R)NCN]n

R = H, Cl, CH3

R = H, (NCN)0.5, CH3

BTSC = (H3C)3Si-N=C=N-Si(CH3)3

(37) Depending on the silicon-bonded substituents R, the viscosity of the solution (sol) increased considerably within 30e60 min (R ¼ Cl, H) or within one day (R ¼ CH3). After approximately three days, the as-obtained reaction mixtures,

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Handbook of Advanced Ceramics

which were then transformed into a gel, began to separate from excess bis(trimethylsilyl) carbodiimide and the byproduct chlorotrimethylsilane. Purification was preformed by removing the volatile components in a high vacuum. Remarkably, this synthetic pathway is rather cheap and easy to perform. Excess starting compound as well as the byproduct chloro-trimethylsilane can easily be recycled. The precursors are obtained in quantitative yields and without the use of solvents. Moreover, time-intensive processing steps are not essentially necessary. Unfortunately, the high temperature properties of ceramics derived from such boron-modified poly(silylcarbodiimide)s were not superior to those obtained from boron-free SieCeN precursors. This was traced back to the chemical composition of the materials, that is, their nitrogen content, which was supposed to be too high. Note that, it was not possible to reduce the nitrogen content in these materials by this process to improve their high temperature performance. For this reason, a synthetic pathway was developed that allows for the synthesis of boron-modified poly(silylcarbodiimide)s with adjustable nitrogen contents while the boron and silicon concentrations were kept almost constant [150,151]. Dehydrocoupling was performed according to Scheme 16 in a one-pot reaction by dissolving the silanes in

tetrahydrofuran and subsequent mixing with different amounts of cyanamide. A catalyst as required for the dehydrocoupling of tris(hydridosilylethyl) boranes and ammonia or amines that is described above was not used. To ensure quantitative turnovers, the reaction mixtures were refluxed for 12 h whereby strong hydrogen evolution and precipitation of the precursors occurred. The above examples show that preparative organometallic chemistry allows for the production of a wide variety of molecular precursors for PDCs. Nevertheless, there is still a need for further development. Future work will focus on precursors that release phase-pure ceramics as well as composites with tunable properties.

2.5.5. SieBeCeN Precursors: attempts Toward Industrial Synthesis For any kind of large-scale testing and application, the availability of preceramic polymers in 10e100 kg amounts is a crucial prerequisite. In a first attempt, Bayer AG (Germany) scaled up the laboratory synthesis of Cl3Si(NH)BCl2, TADB, by a factor of 100, demonstrating the principle feasibility of a large-scale production of this precursor. Unfortunately, major problems were faced, in particular contamination by oxygen, degrading the materials’ high temperature performance. The contaminations were due to R´ B CH

R´

R´

Si(R)H2 H B CH

H2(R)Si

Si

Si

H B

CH3 R n = 0.5

H

CH3

+ n H2NC N

N C N 0.5

H-N1: R = H Me-N1: R = CH3

B CH H

Si

N C N

n = 1.5 R´

Si(R)H2 R´

CH3

n R R´ = C2H4Si(R)NCN H-N6: R = H Me-N6: R = CH3

R´ = C2H4Si(R)H2 R = H, CH3 n=1

= C2H4

Si

thf 12h rflx., - H2

2

R

CH

R´ n=3

B

CH3

CH

R´

N C N

H 0.5

Si

CH3 N C N 0.5

n R R´ = C2H4SiH(R)(NCN)0.5

R n R´ = C2H4SiH(R)(NCN)0.5

H-N2: R = H Me-N2: R = CH3

H-N3: R = H Me-N3: R = CH3

SCHEME 16 Synthesis of boron-modified poly(silylcarbodiimides) by a dehydrogenative coupling of tris(hydridosilylethyl)boranes, B[C2H4e Si(R)H2]3 (R1/4 H, CH3), with cyanamide; H2NeCN. The molar ratio of the starting compounds was 1:0.5 to 1:3. Indexes H and Me differentiate between polymers derived from B[C2H4eSiH3]3 and B[C2H4eSi(CH3)H2]3, respectively, whereas the indexes N1eN6 give the numbers of nitrogen atoms per boron atom. Chemical compositions ranging from (C6.5H20NSi3B) (HeN1) to (C9H15N6Si3B) (HeN6) and (C9.5H26NSi3B) (MeeN1) to (C12H21N6Si3B) (MeeN6) could be realized by this procedure [151].

Chapter | 11.1.10

Precursor-Derived Ceramics

the extreme sensitivity of the molecular and polymeric intermediates against moisture and oxygen in combination with the difficulty of generating sufficiently inert production lines. At Fraunhofer Institute for Silicate Research (Wu¨rzburg, Germany), attempts were made, to process CH3Cl2Si(NH)BCl2, MADB, and the respective polymer batchwise, rather analogous to the laboratory synthesis and the Bayer process [152]. However, similar problems appeared, and the fussy purification and inertization of the apparatus prevented a breakthrough in the industrial synthesis. Synthetic routes that allow for continuous processing are clearly advantageous. They allow the use of miniaturized plants, have facilitated maintenance, and usually exhibit high conversion rates. In this regard, Weinmann et al. published

1047

the first approach toward a non-batchwise synthesis of a single-source precursor to SieB-CeN ceramics. Starting from SiCl4, MeNH2, and BCl3, Cl3SiNMeBCl2 dichloroborylmethyltrichlorosilylamine (DMTA) was obtained in the gas phase in a two-step reaction involving salt elimination reactions via Cl3SiNMeH, trichlorosilylmethylamine (TSMA) as an intermediate [153]. Remarkably, conventional batchwise synthesis is not suitable to produce substantial amounts of DMTA. A scheme and the simplified flow diagram of the continuous process are shown in Figures 4 and 5, respectively. A photograph of the laboratory setup is provided in Figure 6. A constant flow of MeNH2 gas diluted with molecular nitrogen and excess SiCl4 vapor are injected simultaneously

FIGURE 4 Scheme of the continuous two-step synthesis of Cl3SiNMeBCl2, DMTA. For color version of this figure, the reader is referred to the online version of this book.

1st Reactor

2nd Reactor Off-gas

A A

Filter

Excess N2

Condenser

Filter

A A

Off-gas washer

SiCl4 Product N2

MeNH2

BCl3

FIGURE 5 Flow diagram of the reactor for the two-step synthesis of Cl3SiNMeBCl2, DMTA. Straight lines indicate mass flow during the production cycle, whereas doted lines indicate mass flow during the regeneration process. Scheme taken from Ref. [153].

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Handbook of Advanced Ceramics

FIGURE 6 Laboratory setup for the continuous DMTA synthesis including (1) SiCl4, MeNH2/N2 supply, (2) reactor 1 for TSMA production, (3) inductively heated SiC filter, (4) BCl3 supply, (5) reactor 2 for DMTA production, (6) DMTA collection vessel, (7) salt removal unit, (8) cryostat for product condensation, and (9) mass flow controllers and various sensors. Arrows indicate the flow direction during production step. The image is taken from Ref. [153].

into the first reactor that operates at approximately 100  C. Immediate formation of TSMA (gas) and MeNH3Cl (solid) occurs. The hydrochloride precipitates and deposits on the walls of the reactor. N2, TSMA, and excess SiCl4 are completely separated from MeNH3Cl by filtration through a silicon carbide filter and subsequently exposed to a stream of BCl3/N2. Spectroscopically pure DMTA and excess starting compounds SiCl4 and BCl3 are separated from the solid couple product TSMA hydrochloride by filtration. Excess SiCl4 and BCl3 are subsequently removed from DMTA by evaporation under reduced pressure. The ongoing deposition of hydrochlorides in the course of the production decreases the performance of the reactor and finally clogs the filters. Therefore, precursor synthesis based on a cyclic mode, consisting of sequentially occurring production and regeneration steps, is required. Production is performed until the partial clogging of the filters causes a pressure increase within the system, which is monitored online. The supply of reactants is then interrupted and the system subsequently regenerated by vaporizing the amine hydrochlorides in a high vacuum. Fast and quantitative sublimation is enabled by inductively heating the ceramic filters. The maximum production time of the initial setup was 25e30 min. Quantitative regeneration was achieved within another 30 min. Subsequent cooling of the filters took

another 60 min. Accordingly, a full cycle took up to 2 h. The maximum flow rate for the actual laboratory setup (diameter of the reactors was ~3.0 cm) was 4.9 mmol (0.15 g) of MeNH2 per minute, corresponding to a daily synthesis of approximately 440 mmol (108 g) of DMTA. An even more straightforward process is the continuous salt-free two-step synthesis of TADB (cf. Eqn (30)) by silazane cleavage, starting from HMDS, SiCl4, and BCl3 [154]. H

H SiCl4

N Me3Si

SiMe3

- Me3SiCl

BCl3

N Cl3Si

SiMe3

- Me3SiCl

H N Cl3Si

BCl2

.

(38) The reaction is performed in an array of two subsequent high surface area columns filled with Raschig rings. A scheme of the first step, that is, synthesis of N,N,N,0 0 0 trichloro-N ,N ,N -trimethyldisilazane (TTDS) (Cl3SiNHSiMe3) is given in Figure 7. Pressure and temperature are chosen in a way that HMDS and TTDS are liquid and SiCl4 as well as the couple product MeSiCl3 are in the gas phase. Typical reaction conditions were 45  C/160 mbar. A stream of SiCl4 vapor is passed through the system from the

Chapter | 11.1.10

Precursor-Derived Ceramics

1049

FIGURE 7 Scheme of the first step, that is, TTDS synthesis of the continuous two-step synthesis of Cl3SiNHBCl2, TADB. The color gradient mirrors the concentration gradient of the liquid and gaseous species. For color version of this figure, the reader is referred to the online version of this book.

bottom to the top. HMDS is simultaneously introduced on the top of the column. Gravity causes the liquid to flow downward thereby wetting the surface of the Raschig rings. MeSiCl3 elimination and conversion of HMDS to TTDS occur on the surface of the glass rings, causing a concentration gradient of the two liquids within the reactor. The upper part mostly contains HMDS, whereasdon choosing proper reaction conditionsdthe lower part only contains TTDS. The latter is collected in a vessel attached to the bottom of the system. Excess SiCl4 vapor and the vapor of the couple product MeSiCl3 are collected in a cooling trap attached to the upper part of the reactor. The second reaction step is performed accordingly. BCl3 gas is introduced into a column filled with Teflon rings and liquid TTDS given into the column in the reverse flow. The product Cl3SiNHBCl2, TADB, is finally purified by distillation under reduced pressure. However, the determination of proper reaction conditions is of major importance, since otherwise, formation of trichloroborazine as a side product occurs, which precipitates and clogs the column.

According to US 20110028302 A1 [154], TADB is mixed with HMDS, and the mixture is polymerized in a salt-free reaction by a controlled stepwise heat treatment to finally 200  C. Evaporation of MeSiCl3 and completion of the condensation are supported by simultaneously reducing the pressure to finally 10 mbar.

2.6. Precursors to Ternary SieOeC and Quaternary SieMeOeC Ceramics (M [ B, Al, Ti, Zr, Hf, etc.) Poly(organosiloxanes) are preceramic polymers that can be used for the synthesis of silicon oxycarbide-based ceramics on thermal decomposition in inert gas atmosphere. They are generally denoted as silicones, are usually inexpensive and exhibit unique chemical, physical, and electrical properties [155]. The general synthesis method for the preparation of polysiloxanes relies on the reaction of chloro(organo)silanes with water (Scheme 17a).

(a) CH3 Cl

Si

CH3

CH3 Cl

HO

CH3

Si

Si

OH

O

CH3

CH3

n

(b) H3C H3C

Si

Si

CH3

H3C

Si

Si

CH3

H3C SCHEME 17

CH3

O

O

CH3

CH3 CH3 F3CSO3H

Si

Si

CH3 CH3

O n

Synthesis of polysiloxanes via (a) hydrolysis and polycondensation of dichlorosilanes and (b) ROP of cyclic siloxanes.

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Handbook of Advanced Ceramics

In recent years, novel silicon-rich poly(organosiloxanes), named polysilaethers [156], have been synthesized by the polycondensation of a,6-functionalized linear silanes [156] or through the ROP of cyclic siloxanes [157] (Scheme 17b). Furthermore, highly branched poly(organo)silsesquioxanes [RSiO1.5]n have been used as preceramic polymers for the synthesis of SieOeC-based ceramics, since their ceramic yield upon pyrolysis is very high due to their increased crosslinking degree. Crosslinked polysiloxanes or silicone resins can also be prepared by the solegel process via hydrolysis and condensation reactions of hybrid silicon alkoxides. This class of precursors was used for the first time to synthesize silicon oxycarbide glasses [158e160]. They are modified silicon alkoxides of the general formula: 0 RxSi(OR )4x (R ¼ alkyl, allyl, aryl; R0 ¼ methyl, ethyl), which on gelation convert into silicone resins of the composition RxSiO(4x)/2. The solegel process allows for precisely controlling the composition of the silicone resin, because different hybrid silicon alkoxides can be used for cohydrolysis and subsequent polycondensation. Thus, precursors for stoichiometric silicon oxycarbide, as well as for SieOeC materials, showing excess of carbon or silicon have been prepared [161]. Moreover, this preparative technique allows for introducing additional elements within the preceramic network, for example, Al, Ti, B, by using their corresponding metal alkoxides [162]. Recently, the modification of polysilsesquioxanes with transition metal alkoxides (such as zirconium(IV)propoxide or hafnium(IV)butoxide) via solegel-like condensation reactions has been reported (Scheme 18), which leads to preceramic precursors for quaternary SieMeOeC ceramics (M being e.g. Zr, Hf) [163,164]. Thus, a hydroxy- and ethoxy-substituted polysilsesquioxane has been reacted with the alkoxides and subsequently crosslinked and pyrolyzed to furnish SiZrOC and SiHfOC ceramics. The modification of the polysilsesquioxane with the metal alkoxides has been shown to strongly increase the crosslinking degree of the preceramic precursor, which consequently leads to a strong decrease of the open porosity of the resulting ceramics [164].

3. POLYMER-TO-CERAMIC TRANSFORMATION The thermally induced degradation of precursors is a rather complex process, which involves numerous chemical reactions in the solid state as well as in the vapor phase. Due to limitations of monitoring such reactions, there is still a lack of understanding of thermolysis mechanisms in detail. However, because thermolysis significantly influences both the composition and microstructure of PDCs, an understanding of the single steps that occur during the heat treatment is of considerable interest. Accordingly, thermolysis of different types of precursors has been studied [51,55,57,165e182] and reviewed [12,58,183] extensively. As noted in previous sections, preceramic polymers are composed of a polymer backbone to which either single atoms or molecular units are attached. These substituents influence physicalechemical properties, such as solubility, viscosity, and softening or melting, thereby directly determining the processability of the precursor. With respect to thermolysis, it is important that pendent groups may also provide functionality and as a consequence latent reactivity. The latter is known to be a key issue for obtaining high polymer-to-ceramic conversion yields (high “ceramic yields”), because it provides a means for further crosslinking, thus inhibiting polymer backbone degradation and volatilization of low-weight molecular backbone components. In contrast, pendent, nonfunctional groups are eliminated as gaseous byproducts, which do not only induce a mass loss upon the thermal treatment, but are also responsible for retention of impurities or creation of gasgenerated pores and cracks. As a logical consequence, molecular precursors for ceramics should be designed in such a way that exclusively molecular hydrogen is split off during the heat treatment, which would provide the benefit of an easy outgassing because of the high mobility of hydrogen in any ceramic material. Unfortunately, high ceramic yields and good processability are often contradictory. Therefore, compromises are necessary in the precursor design, usually mandating processing with “lowceramic-yield” precursors in order to obtain higher quality ceramic products [183]. A relevant aspect in precursor thermolysis, which until now limits its applicability to the formation of ceramics in

OnBu n

BuO

OH

OEt Si

O

Si CH3

Si

O n

CH3

O

OEt O

Si m OSiR3

Hf(OnBu)4

Si

O

Si CH3

OnBu

Hf

Si

O n

O

CH3

SCHEME 18 Chemical modification of a polysiloxane with hafniumtetra(n-butoxide).

Si m OSiR3

Chapter | 11.1.10

Precursor-Derived Ceramics

small dimensions (thin films, fibers), is the change in density and volume with increasing temperature. Even if ceramic yields are high, dramatic volume shrinkage occurs during thermolysis, which is due to the increase in density. Whereas the density of polymeric precursors is typically in the range of <1 g/cm3 the respective values for fully thermolyzed amorphous ceramics can be at around 2.3e2.8 g/cm3. This means that the increasing density of a specimen during thermolysis causes volume shrinkage to 45% of the initial value or lower. Considering the loss of gaseous byproducts during thermolysis, the volume shrinkage is even higher. For a precursor with a density of 1 g/cm3, which delivers ceramics with a density of 2.3 g/cm3 in 60% yield, the shrinkage is almost 75%. Considering an increase in density up to >3.0 g/cm3 due to the crystallization of the amorphous material at elevated temperature (cf. Section 5.2), the whole situation worsens further. Near net shape manufacturing of PDCs is thus difficult to perform and mostly requires either passive or reactive filler materials to inhibit shape distortion [184,185]. The complexity of the thermolysis process requires the combination of contrasting investigation tools such as thermal and elemental analysis as well as spectroscopic and diffraction methods. Common techniques applied are, for example, simultaneous thermogravimetric analysis (TGA) in which mass changes during the heat treatment are monitored combined with differential thermal analysis (DTA) to receive valuable information on energetic changes during the polymer-to-ceramic conversion. Elemental analysis of the thermolysis intermediates is essential, because it is the only method, which allows for a highly accurate examination of the compositional changes of the precursors that occur during the polymer-to-ceramic conversion. Structural characterization of thermolysis intermediates provides distinct information on chemical reactions that occur during ceramization. Solid-state MAS NMR (in the case of silicon containing precursors especially 29Si-NMR [186]) and FTIR spectroscopy are state of the art. Both are nondestructive methods. NMR spectroscopy represents a tool for probing the local environment of single NMR-active nuclei in amorphous insoluble materials. IR spectroscopy additionally allows for a rapid determination of characteristic vibrational frequencies of structural units, that is, stretching vibrations of OeH, NeH, CeH, SieH, C]N, C]C, N]C]N, etc. or vibrations caused by deformation of three-atom arrays in the precursor or the thermolysis intermediates. A helpful, supplementary method in this regard is Raman spectroscopy.

3.1. Thermolysis of Precursors to Silicon Carbide Among numerous precursors, which are described in the literature for the formation of silicon carbide ceramics, there is a small number of polymers to silicon carbide that

1051

have reached commercial interest. In this regard, the most important systems are poly(dimethylsilane), which is historically seen as the first commercial precursor to SiC/C composites and poly(methylsilane), which is a precursor of near-stoichiometric silicon carbide. As noted in Section 2.2.2 poly(dimethylsilane), [Si(Me)2]n, can be thermally rearranged to yield poly (methyl-silylenemethylene), [(H)Si(Me)eCH2]n. In contrast to [Si(Me)2]n, the latter is a soluble and meltable polymer. [(H)Si(Me)eCH2]n was first applied in 1975 to produce SiC-containing fibers in the so-called Yajima process [26,187]. The procedure is schematically presented in Figure 50 in Section 5.5. Poly(methyl-silylenemethylene) has a molecular weight MW ¼ 1200. It can be melt spun and delivers green fibers with poor tensile strengths of around 5e10 MPa. Even though the precursor is obtained in a rather simple procedure and in high yields, it suffers from the lack of latent reactivity, that is, the ability to thermally crosslink after the spinning process. As a consequence, fiber integrity is lost during thermolysis. To maintain fiber integrity, it is thus necessary to subsequently cure the green fibers. In this step, the fiber surface is chemically modified and the fiber rendered infusible. Curing can be performed by different procedures: according to Scheme 19, heating of the fibers in air results in a partial oxidation of surface SieH (rather than CeH) groups with formation of SieOH units. A subsequent condensation step produces a thin SieCeO film on the surface of the fibers [183]. This silicon oxycarbide film efficiently inhibits melting of the fiber during thermolysis. Oxygen, which is introduced in this step, cannot be removed during the heat treatment and remains in the fiber, limiting the maximum application temperature to approximately 1150  C. At higher temperature decomposition with the formation of volatile species such as CO and SiO as well as crystallization occurs, leaving pores and large crystallites behind, which substantially decrease the mechanical properties. Alternatively, electron beam irradiation can serve for fiber curing. Compared to green fiber curing in air, it is much more expensive and more difficult to perform, but no oxygen is introduced into the green fiber. Consequently, thermolysis delivers oxygen-free SiC/C fibers with enhanced mechanical properties and higher maximum application temperature. Thermolysis of the green fibers has to be performed in an inert atmosphere. Yajima et al. reported on thermolysis of NicalonÔ green fibers in a vacuum at finally 1500  C to obtain 7-mm-thick SiC fibers, which exhibit a density of approximately 2.5 g/cm3, tensile strengths of 1.5e3 GPa, and a Youngs modulus of 170e200 GPa, depending on the processing applied [188e190]. In contrast, Hi-NicalonÔ fibers exhibit a density of 2.74e3.1 g/cm3, tensile strengths

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Handbook of Advanced Ceramics

CH3 Si CH2

CH3

ΔT, O2

Si CH2

H

O

CH3

CH3

Si CH2

Si CH2

+ H

O CH3 Si CH2

+

H

O

Si CH2

CH3

Si CH2

2

H

O

Si CH2 OH

CH3 +

+

O O

CH3

CH3

Si CH2 OH

OH

CH3 Si CH2

CH3 2

Si CH2

O - H2O

OH

Si CH2 CH3

SCHEME 19

Proposed mechanism for the surface oxidation in the curing step of silicon carbide fibers [183].

of 2.8 GPa, and elastic moduli of 270e420 GPa, depending on their (poly-) crystallinity. Moreover, their maximum application temperature is >1400  C compared to 1150  C for that of NicalonÔ fibers. For further details, see Section 5.5. TGA of the thermolysis step, which released ceramics in 60% yield, indicated that the thermal decomposition of the polymer proceeds gradually >300  C and is completed at 800  C. The pyrolytic degradation of the polymer does not take place “stoichiometrically” by elimination of methane and hydrogen according to CH 3 1/n

ΔT

Si CH 2 H

SiC + H2 + CH4

(39)

n

since almost 50% of carbon, which is present in the precursor as methyl groups, is retained in the ceramic fibers. Consequently, SiC/C composite fibers without defined microstructures rather than phase-pure SiC fibers are obtained. To yield phase-pure SiC, alternative improved precursors were developed. The idea behind these concepts was to design polymers in which the final elemental composition of the ceramic, that is, the 1:1 stoichiometry of the constituting elements silicon and carbon is already preformed. Moreover, the synthesis of precursors was aspired, which does not produce gaseous byproducts upon thermolysis other than hydrogen. Polymers, which (at least theoretically) fulfill these requirements, are polymethylsilane, [SiMeH]n [24], and polysilaethylene,

[H2SiCH2]n [27,29], which can be obtained according to procedures described in Section 2.1. Although the precursors exhibit a 1:1 Si:C ratio, it is difficult to obtain stoichiometric or near-stoichiometric SiC. The reasons are (i) that the Kumada rearrangement is a radical reaction, which takes place at elevated temperature with low selectivity and (ii) that depolymerization of the precursors during thermolysis occurs with the volatilization of carbon and/or silicon species, which cause changes in the chemical composition of the precursors. IR and NMR investigations of the thermolysis intermediates provide information on chemical reactions that occur during the heat treatment. In the case of [Si(Me)(H)CH2]n, at least four different steps can be distinguished. Below 400  C, low weight oligomers evaporate. Up to 550  C, molecular weight increases due to the formation of new SiC bonds that can be traced back to a dehydrocoupling of SieH and CeH units [166,191]. In contrast, there is no evidence for the formation of SieSi linking. A Further heating to 800  C results in the degradation of residual SieH, CH3, and SieCH2eSi units with evaporation of low molecular weight organic species, that is, molecular hydrogen and methane. This stage corresponds to the organic-to-ceramic transition and releases an inorganic covalent SiC network. Additionally, graphitic carbon forms, which is evident from 13C MAS NMR investigations [191]. Above 800  C, only very small weight loss is observed, which is due to the evaporation of residual hydrogen. Heating to temperatures exceeding 1100  C finally results in the crystallization of the predominantly amorphous ceramic. Details on this topic are presented in Section 4.

Chapter | 11.1.10

Precursor-Derived Ceramics

3.2. Thermolysis of Precursors to Silicon Carbide/Nitride Ceramics Due to the structural diversity of SieCeN precursors compared to SiC polymers, thermolysis seems to be much more complex. Although for SiC usually polysilanes or PCSs serve as preceramics, polysilazanes or poly (silylcarbodiimide)s are mostly used as precursors for ternary SieCeN ceramics. The ternary phase diagram in Figure 8 distinguishes different types of compositions in the SieCeN system. Poly(silsesquiazane)s [192] or poly(silylcarbodiimide)s [13,89,91,92], which are both comparably rich in nitrogen, usually deliver ceramics with compositions located on the tie-line CeSi3N4 (A). In contrast, ceramics derived from polysilazanes with the general structure [(R1)(R2) SieNHe]n (R1, R2 ¼ single-bonded organic unit). [41,51,55,56,167,170,193e197] have compositions located in the three-phase field SiC/Si3N4/C (B). Laine et al. published an exception from this general rule. They obtained Si3N4/C (type A) compositions by thermolysis of [H2SieN(CH3)e]n rather than silicon nitride and silicon carbide-based composites [177]. Compositions located in the three-phase field SiC/Si3N4/Si (C) were obtained by copyrolysis of PCS and PHPS [198e200]. Even though a comparison of the elemental composition of the precursors and their derived ceramics gives first hints for possible reactions that occur during the heat treatment, it is inevitable to structurally characterize thermolysis intermediates accurately. In general, methods such as NMR and IR spectroscopies, elemental analysis, and (mass-spectra-coupled) TGA are considered for this purpose.

3.2.1. Thermolysis of Polysilazanes Most of the work in this field concentrated on thermolysis of polysilazanes. It turned out that the molecular architecture of the polymer backbone, the pendent functional

1053

groups, and the crosslinking of the precursors significantly influence the polymer-to-ceramic conversion. As mentioned in Section 2.3, ammonolysis of dichlorosilanes usually delivers six- or eight-membered cyclosilazanes, [SiR1R2eNH]n (n ¼ 3, 4). Because of their low molecular weight, as-obtained polymers are mostly not qualified as precursors for ceramics. In the worst case, thermolysis of such volatile molecules results in a 100% mass loss. Moreover, cyclosilazanes are generally liquids, which have poor rheological properties, disqualifying them as precursors for ceramic fibers or bulk materials. A procedure for increasing ceramic yields of lowweight polysilazanes is a catalyst-induced crosslinking by dehydrocoupling of SieH and NeH units. For example, Seyferth and Wiseman described crosslinking reactions of cyclotetrasilazane, yielding a highly branched polymer, which is soluble in tetrahydrofuran [52,53]. Details on the procedure including a proposed mechanism are presented in Section 2.3. The increased crosslinking was directly mirrored in increased ceramic yields from <30% to approximately 84%. The incorporation of reactive sites in oligomeric silazanes, for example, vinyl groups, also results in increased ceramic yields. This is directly reflected by a comparison of ceramic yields of [SiMeHeNH]n (20%) [195] with that of [SiViHeNH]n (Vi ¼ HC]CH2, 83%) [56] or [SiMe2e NH]n (50%) and [SiViMeeNH]n (76%) [55], as determined by TGA. Detailed investigations on the thermolysis of vinylsubstituted oligosilazanes of the general type [SiViReNH]n and the influence of the vinyl group in crosslinking reactions were published by Bill et al. (R ¼ Me, (NH)0.5) [57,201] and by Mocaer et al. (R ¼ H) [176]. Based on solid-state 1H, 13C, and 29Si NMR and FTIR investigations, Bill et al. suggested conclusive thermolysis mechanisms. For example, crosslinking reactions involving SieC coupling accompanied by the elimination of hydrogen (Eqn (40)) or methane (Eqn (41)) were considered.

N

Si

CH 3 + H

Si

Si CH 2 Si

+ H2

(40) Si3N4 (A) Si

(B)

CH 3 + H3C

Si

Si CH 2 Si

+ CH4

(C)

(41) C

SiC

Si

FIGURE 8 Ternary SieCeN phase diagram (T < 1440  C, 1 atm N2). A: compositions on the tie-line CeSi3N4; B, C: compositions in the three phase fields SiC/Si3N4/C or SiC/Si3N4/Si, respectively.

They occur at around 500  C and most likely follow radical mechanisms. Further increasing the temperature, finally results in the formation of CSi4 tetrahedra.

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Handbook of Advanced Ceramics

Raising the temperature to higher values results in SieC bond cleavage and b-hydride shift with the formation of SieH bonds and hydrogenated sp2-carbon atoms. Under the given conditions, the latter eliminate hydrogen and transform into sp2-carbon.

Si 2

Si CH 2 Si

Si

C

Si

+ CH4

(42)

Si

Vinyl groups that are attached to the silazane skeleton enable crosslinking by olefin polymerization at low temperatures, usually around 300  C. It will be discussed below (Figure 9) that 13C NMR spectroscopy with crosspolarization and magic angle spinning is an ideal tool for monitoring the progress in such reactions. The sp2-carbon atoms of the vinyl groups generate NMR signals at approximately 130e140 ppm. During the crosslinking, they are transformed into sp3-carbon atoms, which generate resonance signals at around 10e30 ppm. H C

CH 2

CH

H Si

*

CP

*

SieH units play an important role in that they have reactive sites for hydrosilylation and/or dehydrocoupling. The former additionally requires C]C or C^C units in the polymer and mostly takes place below <200e250  C. SieN bond formation by dehydrocoupling of SieH and 29

Si

1050

CP

800

CP

600

CP

500

CP

CP

400

CP

CP

300

CP

CP

200

CP

dec

Precursor

dec

CP

*

CP

*

*

*

*

CP

200

100 ppm

0

CH CH

sp2-C + H2

T (°C)

C

+

(44)

(43)

13

CH2

Si

CH 2

Si

Si

CH

0

-100 ppm

FIGURE 9 Experimental 13C (left) and 29Si (right) NMR spectra of PMVS, [SiViMeeNH]n at various stages of the thermolysis process. All spectra were obtained under MAS conditions using the crosspolarization (CP) technique [202].

Chapter | 11.1.10

Precursor-Derived Ceramics

1055

NeH units is important, because in combination with the above-discussed C]C polymerization and subsequent SieC bond cleavage, it sufficiently explains the transformation of SieC units, present in the precursor, into SieN (finally SiN4) units.

Si

H + H

Si

N

+ H2

N

(45)

In the pyrolytic degradation of oligo- or polysilazanes of the type [SiR1R2eNH]n, NeH bonds are frequently involved in trans-amination reactions. These take place at temperatures around 200e500  C and result in the formation of NSi3 sites among gaseous ammonia, which is eliminated. Si

H 3

Si

N

2

Si

Si

N

Si

+ NH3

(46)

such a spectral breakdown is represented as an example for PMVS [202]. The 13C CP MAS NMR spectrum of as-synthesized PMVS precursor shows NMR signals at around 130 and 140 ppm, which are generated by the olefin carbon atoms and one at 0 ppm, which is assigned to the silicon-bonded methyl group. The appearance of the spectrum does not change after heating the polymer to 200  C, indicating that in this temperature range, no chemical modifications of the precursor involving carbon atoms take place. On heating the sample to 300  C, drastic changes are observed. The signals at 130 and 140 ppm significantly decrease in intensity, whereas a broad resonance appears at approximately 25 ppm. This observation points to an olefin polymerization of the pendent vinyl groups in the precursor according to Eqn (43). In Scheme 20, this reaction indicates the onset of the thermal degradation of the polymer. In the 13 C NMR spectrum of the sample heated to 400  C, the latter signal decreases in intensity, whereas a new broadened signal at around 130 ppm is found. This points to an SieC bond cleavage and the formation of sp2-carbon atoms as pointed out in Eqn (44). Further raising the temperature

The loss of ammonia leads to the reduction of the nitrogen content in the ceramic materials. The tendency for trans-amination reactions is especially pronounced for poly(silsesquiazane)s of the general type [SiR(NH)1.5]n, because these polymers are rich in nitrogen. Their composition is located in the three-phase field CeSi3N4eN (cf. Figure 8), whereas that of the derived ceramic lies on the tie-line CeSi3N4, as will be discussed below. Finally, SieN coupling with the elimination of methane is often considered. This reaction requires a cleavage of thermodynamically relatively stable SieC bonds. It thus takes place exclusively at elevated temperature.

H C N

CH 2

Si

N

CH 3

CH N

CH 2

Si

N

CH 3 H

Si

CH 3 + H

N

Si

N

+ CH4

C

C

N

N

CH 3

(47) In the following, the pyrolytic degradation of two structurally different oligosilazanes, which yield ceramics with different chemical composition, is discussed in detail. For this purpose, poly(methylvinylsilazane) (PMVS), [SiViMeeNH]n, and nitrogen-rich poly(vinylsilsesquiazane) [SiVi(NH)1.5]n (PNVS) are considered. Both polymers do not only differ in their chemical composition but also in their crosslinking densities. The first is a linear oligomer (or cyclomer), whereas the second constitutes a highly branched network. Reliable information on the pyrolysis mechanism is received especially from NMR spectra, which for this intention, have to be analyzed in detail. In the following,

Si

Si Si N Si

C(sp 2)

C Si C

N

Si

N

N

N

Si

N

Si

Si

C

Si

Si

Si

SiN2C2/CSi4

SiN3C/CSi 4

N

Si

N

N

SiN4

SCHEME 20 Proposed thermolysis intermediates and structural ceramic units that evolve during the polymer-to-ceramic conversion of PMVS, [SiViMeeNH]n [202].

1056

to 500  C, causes only little changes in the appearance of the spectrum. However, signals appear broader, and the resonance of the sp2-carbon atoms is somewhat more intensive. Remarkably, the signal of the methyl groups still remains unchanged. At 600  C, again drastic changes are observed. The intensity of the signal at 130 ppm strongly increases, whereas that at 0 ppm significantly broadens. This temperature marks the decomposition of residual organic functions, that is, the silicon-bonded methyl groups. It is accompanied by the loss of hydrogen and/or methane and the formation of SieC bonds (cf. Eqns (40e42)). Moreover, methane that is split off partly decomposes to graphitic carbon and molecular hydrogen. Further information on the thermal degradation of PMVS is received from 29Si NMR spectroscopy. The spectrum of the precursor displayed in Figure 9 shows two sharp lines at 14 and 17 ppm, which point to a mixture of cyclic trimers and tetramers. Heating to 200  C causes no changes in the signal appearance, whereas raising the temperature to 300  C results in a low field shift of the signals to around 3 ppm. This is in accordance with the transformation of the sp2-hybridized vinyl carbon atoms to sp3 hydrocarbons. At 400  C, a new signal at 21 ppm evolves. This value is typical for C(sp3)SiN2H sites and clearly points to a rearrangement with SieC bond cleavage and SieH bond formation as is suggested in Eqn (44). However, it is also possible that in this stage, C(sp3) SiN3 sites form by a dehydrocoupling reaction of SieH and NeH units. Their resonance appears in a similar chemical shift range as that of C(sp3)SiN2H sites. At 500  C, a new signal at 45 ppm emerges, which can unequivocally be assigned to SiN4 sites, most probably due to the reactions according to Eqns (44) and (45). Above 500  C, the NMR signals broaden and overlap substantially. However, it can be concluded without doubt that SiC2N2, SiCN3, and SiN4 sites generate the broad signal observed at and >600  C. According to the observations in the 13C and 29Si NMR spectra of the pyrolysis intermediates of PMVS and the proposed pyrolysis mechanism in Scheme 20, a ceramic evolves, which is composed of sp2-carbon and an amorphous SieCeN phase, in which silicon is bonded in SiC4xNx (x ¼ 2e4) environments. This is in accordance with the finding of the elemental analysis, which confirms that the elemental composition of PMVS-derived ceramic is located in the three-phase field CeSiCeSi3N4 (cf. Figure 8). The lack of aliphatic silicon-bonded hydrocarbons and the higher crosslinking in PNVS, in contrast, result in a ceramic material with different composition. Chemical reactions that take place during the heat treatment are depicted in Scheme 21. They are basically comparable with those of PMVS. This is reflected by a similar onset of the decomposition temperature, marked by a polymerization of the silicon-bonded vinyl groups at around 300  C.

Handbook of Advanced Ceramics

H C N

Si

CH 2 N

N

CH N

Si

CH 2 N

N

H C

C

N

Si

N

N N N

Si

N

N

C(sp 2)

SiN4

SCHEME 21 Proposed thermolysis intermediates and structural ceramic units that evolve during the polymer-to-ceramic conversion of nitrogenrich poly(vinylsilsesquiazane) [SiVi(NH)1.5]n (PNVS) [201].

During the further heat treatment, the aliphatic hydrocarbon chains are split off with the formation of SiHN3 and SiN4 sites among sp2-carbon that segregates. Subsequent dehydrocoupling of SieH and NeH units as well as transamination reactions result in the formation of a ceramic material, which is composed of sp2-carbon and silicon nitride. This is nicely reflected in both 13C and 29Si NMR spectra of the final thermolysis products (not shown here), which give no evidence for any SieC units [201,202]. Consequently, the compositions of ceramics obtained from PNVS are located on the tie-line CeSi3N4, which is verified by elemental analysis. Recently, amorphous ceramics with compositions along the tie-line SiCeSi3N4 could be obtained [62]. For the development of such materials, the above-discussed findings were considered carefully. The synthesis of the precursor is discussed in detail in Section 2.3 (cf. Scheme 10). In this regard, the main objective was to design a precursor in which the final elemental composition of the ceramic, that is, the Si4N4C stoichiometry of the constituting elements is already preformed. It was presumed that the polymer-to-ceramic conversion does not suffer from the elimination of gaseous thermolysis byproducts other than hydrogen.

Chapter | 11.1.10

Precursor-Derived Ceramics

1057

Therefore, the following idealized thermolysis reactions were considered, taking only hydrogen evolution on heat treatment into account. l

l

sufficient latent reactivity due to the presence of both SieH and NeH units. These enable thermally induced dehydrocoupling reactions according to Eqn (46) and efficiently avoid depolymerization and volatilization of low molecular species during the heat treatment. The observed mass loss of 5.5% corresponds to the hydrogen content of the precursor determined by elemental analysis and points to the fact that exclusively hydrogen evaporates upon thermolysis. This assumption was confirmed by the results of the mass spectroscopic investigations. Hydrogen elimination starts at approximately 200  C and continues up to 1000  C. Above 1000  C, the concentration of hydrogen increases again, parallel with that of a species with m/z ¼ 18. The latter corresponds to water, which in this temperature range is eliminated from the alumina crucible in which the measurement was performed. The amount of gaseous species <1000  C with an m/z ratio >2, however, is negligible, suggesting that the Si:C:N ratio is retained during the polymer-to-ceramic transformation. However, a very small amount of ammonia and methane and their corresponding fragments can be detected unequivocally, though in low intensity, at approximately 380  C and approximately 580  C, respectively. A precise comparison of the relative intensities of methane and ammonia with that of hydrogen is not yet possible, because the large amount of hydrogen exceeded the detector’s capacity.

thermolysis of PHPS [SiH2eNH]n (sum formula Si4N4H12 for n ¼ 4) results in a composite Si3N4 þ Si; thermolysis of poly(methylsilazane) [H3CSiHeNH]n (sum formula Si4N4C4H20 for n ¼ 4) delivers a material with a chemical composition Si3N4 þ SiC þ 3C.

Accordingly, copolymers composed of (SiH2eNH) and (H3CSiHeNH) building blocks in a 3:1 ratio such as [(SiH2eNH)3(H3CSiHeNH)]n (sum formula Si4N4CH14) should thus deliver Si3N4 þ SiC ceramics without “free” carbon or silicon. Consequently, precursor synthesis was performed by coammonolysis of H2SiCl2 and H3CSiHCl2 in a 3:1 ratio at 10  C using tetrahydrofuran as solvent according to Scheme 10. To avoid depolymerization and volatilization of low molecular species during the polymer-to-ceramic conversion, the as-obtained precursor was further crosslinked by base-catalyzed dehydrocoupling according to a method described earlier by Seyferth et al. using 0.1 mol-% of n-butyl lithium as a catalyst. The theoretical ceramic yield of the crosslinked precursor, presuming only hydrogen evaporation, is around 95%. The polymer-to-ceramic conversion was monitored by means of TGAeMS as shown in Figure 10. Thermolysis up to 1200  C takes place in a two-step decomposition in the 200  Ce900  C range. The ceramic yield of 94.5% is the highest value observed for SieCeN precursors so far. Thermolysis of the non-cross-linked precursor, in contrast, delivers ceramics in only 78% yield [203]. Accordingly, the high ceramic yield of the crosslinked precursor is mainly a consequence of the high crosslinking density of the polymer combined with

3.2.2. Thermolysis of Polysilylcarbodiimides In contrast to the SieN skeleton in polysilazanes, which is retained during the thermal degradation of the polymers, the SieN]C]N backbone in poly(silylcarbodiimide)s is decomposed quantitatively. Neither findings in NMR nor IR investigations gave a hint for the presence of such structural units in the amorphous ceramics. The only

1,00E-010

98

8,00E-011

.

TGA

H2+

96

6,00E-011

94 +

.+

NH 3 NH +

92

CH 3

.

+ CH4

4,00E-011 H2O.

2

+

90

2,00E-011

88

0

Rel. Intensity [a.u.]

Sample W eight [%]

100

0,00E+000 200

400

600

800

1000

1200

Temperature [°C] FIGURE 10 The TGAeMS of crosslinked [(SiH2eNH)3(H3CSiHeNH)]n. Fragments ranging from m/z ¼ 2 to m/z ¼ 84 are considered. Heating rate: 5  C/min; atmosphere: flowing Ar [62].

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exception are two metastable crystalline phases, SiC2N4 and Si2CN4 (cf. Figure 2), which were published by Riedel et al. [95]. As already pointed out in Section 2.4.2, these phases were obtained by a reaction between tetrachlorosilane and bis(trimethylsilyl)carbodiimide and subsequent heating to 400 and 920  C, respectively. Basically, there are two different degradation mechanisms that additionally have to be considered in the thermolysis of poly(silylcarbodiimide)s. The first mechanism involves the elimination of acetonitrile, H3CeCN, or HCN from poly(silylcarbodiimide)s exhibiting (NCN)SieCH3 or (NCN)SieH units, respectively. R Si

Si

N C N

N

+ R-C N

R = H, CH 3

(48) 13

Both CH3CN or HCN elimination result in the formation of novel SiN motifs. Acetonitrile that splits off subsequently decomposes with the formation of sp2-carbon. Simultaneous elimination of cyanogene, N^CeC^N, and nitrogen takes place at elevated temperatures.

6

Si

N C N

Si

Si

4

Si + 3 (CN) 2 + N2

(49) According to the formerly mentioned degradation, SieNeSi linkages form. In addition, this type of reaction is responsible for the thermal degradation of SiC2N4 to Si2CN4, as already mentioned above (Figure 2 [95]). Even though structurally different, polysilazanes and poly(silylcarbodiimide)s follow comparable degradation 29

T (°C)

C

Si N

Si

CP 1000

CP

CP

800

CP

CP

600

CP

CP

500

CP

CP

400

CP

CP

300

CP

CP

200

dec

Precursor

200

100 ppm

0

CP

dec

0

-100 ppm

FIGURE 11 Experimental 13C (left) and 29Si (right) NMR spectra of PMVSC, [SiViMeeNCN]n at various stages of the thermolysis process. All spectra were obtained under MAS conditions using the CP technique [182,202].

Chapter | 11.1.10

Precursor-Derived Ceramics

routes. As an example, the progress in the thermal decomposition of poly(methylvinylsilylcarbodiimide), [SiViMeeNCN]n (PMVSC), which was investigated by solid state 13C and 29Si MAS NMR is discussed (Figure 11, Scheme 22) and compared with that of PMVS [SiViMeeNH]n (see Scheme 20) [182]. Similar to the 13C CP MAS NMR spectrum of PMVS, PMVSC shows NMR signals at around 135 ppm, which are generated by the vinyl carbon atoms, as well as a line at 0 ppm, which is assigned to the silicon-bonded methyl group. An additional signal at 125 ppm can be attributed to the carbodiimide carbon atom, which vanishes after heating to 200  C. According to the thermal degradation of PMVS, further heating to 300  C causes polymerization of the pendent vinyl groups and results in the appearance of a broad resonance at around 20e25 ppm, which can be attributed to aliphatic carbon atoms. The latter disappears after heating to 400  C. Moreover, the line width of the N] C]N resonance at 125 ppm increases considerably, most probably because of the formation of “free” carbon, which also shows resonance in this chemical shift range. At 500  C, signals broaden further and the low field signal increases in intensity, indicating further formation of graphite. The absence of CeH resonance signals after

H C

N C N

CH 3

CH N C N

CH 2

Si

N C N

CH 3 N C

C

N C N

Si

N

CH 3

N N

Si

N

N

C(sp 2)

heating to 600  C points to the fact that the thermal degradation of the precursor is almost completed. Similar insights are obtained from 29Si MAS NMR spectroscopy. Olefin polymerization at temperatures up to 300  C causes a low field shift of the 29Si signal from 35 to 20 ppm, which is a consequence of the transformation of SiC(sp2) C(sp3)(NCN)2 sites into SiC(sp3)2(NCN)2 motifs. Further heating causes multiple reactions, resulting in the formation of at least 4 differentiable 29Si sites. According to Scheme 22, the formation of SiN2C(sp3)(NCN) (d ¼ 32 ppm) and SiN3C(sp3) (d ¼ 16 ppm) sites is preferred. Above 500  C, all low field signals vanish, whereas a broad signal forms that can unequivocally be assigned to SiN4 sites. In conclusion, a comparison of the reactive intermediates in the thermolysis of [SiViMeeNH]n (PMVS) and [SiViMeeNCN]n (PMVSC) clearly shows that the polymer structure and chemical composition significantly influence the polymer-to ceramic conversion and therefore the chemical composition of the PDCs. In contrast to PMVS which delivers a ceramic with a composition located in the three-phase field CeSiCeSi3N4 (Figure 8), thermolysis of PMVSC results in a material with a composition located on the tie-line CeSi3N4.

3.3. Thermolysis of Precursors to SieBeCeN Ceramics

CH 2

Si

1059

SiN4

SCHEME 22 Proposed thermolysis intermediates and structural ceramic units that evolve during the polymer-to-ceramic conversion of PMVSC [SiViMeeNCN]n [182,202].

Until now, not much is published on the thermolysis behavior of precursors for SieBeCeN ceramics. So far, only precursors of the general structure [B(C2H4SiReXe)3]n (R ¼ single bonded organic ligand; X ¼ NH, NCN) have been investigated. It was Prof. Dr. Klaus Mu¨ller, together with his students, who for the first time in detail examined the conversion of the multinary precursors to quaternary SieBeCeN ceramics by means of solid-state NMR spectroscopy [54,201,202,204]. The extension of the ternary SieCeN system with boron results in additional reactions during the thermal degradation of the polymeric precursors. In this regard, rearrangements that involve BeC bond cleavages and BeN bond formations have to be considered especially. In this regard, two different reactions must be taken into account: BeC bond cleavage and BeN bond formation with (Scheme 23, reaction pathway a) or without (Scheme 23, reaction pathway b) SieN bond cleavage. In the first case (a), SieCH(CH3)eSi motifs form, whereas simultaneous BeC and NeH bond cleavages according to reaction pathway (b) result in the formation of pendent SieCH2eCH3 units. Such reactions are best monitored by 11B, 15N, or 29Si MAS NMR spectroscopy, whereas 13C MAS NMR does not allow for a distinct differentiation of individual resonance signals, because of a significant signal overlap of all

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H3C

CH B

H3C

CH

Si CH

B

Si

H

+

(a)

N

CH 3

H

N

H a)

H

H3C

b)

CH 2 Si

B

(b)

N

H SCHEME 23 Possible polymer skeleton rearrangements during the thermolysis of boron-modified polysilazanes [B(C2H4SiHeNHe)3]n due to BeC bond cleavage: (a) including SieN bond cleavage and SieC bond formation; (b) without SieN bond cleavage.

aliphatic carbon atoms. Unfortunately, the natural abundance of 15N is too low for obtaining NMR spectra with a sufficient signal-to-noise ratio. For performing 15N NMR spectroscopy, 15N-enriched samples must therefore be synthesized. As an example, the single steps in the thermal degradation of boron-containing [B(C2H4SiHeNH)3]n, which can be obtained from the above mentioned Poly(hydridovinylsilazane) (PHVS) [SiViHeNH]n in a hydroboration reaction using borane dimethyl sulfide [136], will be discussed on the basis of 13C, 29Si (Figure 12) 13

29

T (°C)

C

Si

and 11B MAS NMR spectroscopy (left spectrum in Figure 13). Additionally, the results of 15N MAS NMR investigations (Figure 13, right spectrum) will be reviewed. The 13C NMR spectrum of the polymeric precursor [B(C2H4SiHeNH)3]n (Figure 12, bottom left) exhibits two strong signals at 12 and 28 ppm, which refer to BeC2H4eSi groups along with two very weak resonance signals at 124 and 139 ppm. The latter corresponds to remaining C]C units that due to steric effects did not react during the hydroboration of PHVS. After heating to 200  C, the weak olefin signals observed in the precursor vanished completely. This points to the fact that increasing the temperature enables quantitative hydroboration of the C]C units. Upon heating to 400  C, the signal at 28 ppm

1050

*

11

15

N

B T (°C)

*

600

*

1050

500

600

*

*

500

400

400

200 200

Precursor *

200

*

100

Precursor

100

0

0

-100 ppm

ppm 13

29

FIGURE 12 Experimental C (left) and Si (right) NMR spectra of boron-containing [B(C2H4SiHeNH)3]n [133], which was obtained by hydroboration of PHVS [SiViHeNH]n [56], at various stages of the thermolysis process. All spectra were obtained under MAS conditions using the CP technique [202,204].

200

0 ppm

-200

-100

-300 ppm

-500

FIGURE 13 Experimental 11B (left) and 15N (right) NMR spectra of [B(C2H4SiHeNH)3]n [133], which were obtained by hydroboration of PHVS [SiViHeNH]n [56], at various stages of the thermolysis process. All spectra were obtained under MAS conditions [202,204].

Precursor-Derived Ceramics

(aliphatic CHCH3) disappeared, whereas aliphatic signals at 12 ppm (generated by CH2 and CHCH3) are still present. Further heating to 600  C results in a remarkable signal broadening and the appearance of a new, very broad resonance signal at around 130 ppm, which can be attributed to graphite-like domains. At higher temperature, the intensity of the latter spectral component increases further at the expense of the signal intensity of the aliphatic carbon atoms, which is now centered at around 15 ppm, a value, which is typical for CSi4 environments. The 29Si NMR spectrum of the precursor in which silicon has a (SiHC(sp3)N2 environment) exhibits a broad line centered at 13 ppm, which is in accordance with theoretical calculations. Upon heating to 400  C, two new signals emerge. One of these signals with very low intensity appears high field shifted to 37 ppm. This refers to SiC(sp3) N3 sites that must have been formed in dehydrocoupling reactions involving SieH and NeH units. In contrast, the second signal is observed at lower field (4 ppm), corresponding to SiHC2(sp3)N units, which most likely have been formed by a polymer skeleton rearrangement including SieC bond formation by BeC and SieN bond cleavage according to reaction pathway (a) in Scheme 23. Further raising the temperature results in a broad resonance signal centered at 30 ppm, which is caused by additional dehydrogenative SieN coupling, finally resulting in the formation of a mixture of SiC(sp3)2N2 and SiC(sp3)N3 sites. FTIR spectra of the thermolysis intermediates confirm the above assumption. Due to the quantitative dehydrogenative SieN coupling, NeH and SieH stretching vibrations, which are observed in the IR spectra of the polymeric precursors and which appear as very strong signals at around 3380 and 2124 cm1, respectively, completely disappear after annealing the samples to 600  C. Accordingly, 11B and 15N MAS NMR spectra show pronounced changes in the 400e600  C range. The experimental 11B NMR spectra shown in Figure 13 refer to the excitation of the central transition (m1 ¼ 1/2 to m1 ¼ þ1/2). Owing to the quadrupolar moment of the boron nucleus, the signals of all 11B NMR spectra are significantly broadened (a second-order broadening, which cannot be eliminated by fast rotation at the magic angle). The resonance signals in the 11B NMR spectra of the polymeric precursor and the thermolysis intermediates <500  C are centered at around 0 ppm. They appear broadened, featureless, and without any fine structure. Their line shapes point to a significant heterogeneity in the local environment of the boron nuclei, presumably due to the presence of a mixture of trigonal and tetrahedral coordination spheres. However, exceeding 500  C results in a low field shift to 15 ppm and the formation of a signal with distinct special features (quadrupolar coupling constant ¼ 2.8e2.9 MHz), which are typical for trigonalcoordinated boron nuclei. These spectra are almost

1061

identical to the spectrum of h-BN. This observation confirms the above-proposed rearrangement (Scheme 23) of the precursor at 500  C, which is accompanied by a BeC bond cleavage and BeN bond formation. 15 N NMR spectra of 15N-enriched samples are given in the right column of Figure 13. The NMR spectrum of the precursor exhibits a line at around 360 ppm arising from NSi2H units, whereas the 400  C sample exhibits an additional broad signal at 245 ppm. The latter can be attributed to NHB2 and NHBSi units, which are generated by the above-discussed rearrangement. After heating the sample to 600  C, a very broad signal is seen in the 15N NMR spectrum, which can be traced back to a superposition of NHB2, NHBSi, NSi3, and NSi2H sites. Finally, after heating to 1050  C, only one signal remains at around 320 ppm, which clearly reflects the formation of an amorphous ceramic network including NB3 and NSi3 structural units. Remarkably, the ceramic yield of [B(C2H4SiHeNH)3]n is >86%, which is an exceptionally high value. Figure 14 shows the result of thermogravimetric investigations of the precursor. Thermal degradation of the precursor occurs in one step and starts at around 280e300  C (the insignificant mass loss <200  C is due to the evaporation of residual solvent, which was used in the synthesis of the polymer). This is exactly the temperature range in which the onset of the BeC bond cleavage followed by the abovementioned rearrangement (cf. Scheme 23) was observed by a solidstate MAS NMR spectroscopy. The mass loss of 15% is due to the volatilization of both methane, which is eliminated from SieCH(CH3)eSi motifs and hydrogen. The latter evaporates due to the elimination from hydrocarbons and dehydrogenative SieN coupling.

0 Weight Change [%]

Chapter | 11.1.10

-5

-10

-15

250

750 1000 500 Temperature [°C]

1200

FIGURE 14 TGA of [B(C2H4SiHeNH)3]n [133], which was obtained by hydroboration of PHVS [SiViHeNH]n [56]; heating rate: 2  C/min; atmosphere: flowing argon [133].

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3.4. Thermolysis of Precursors to SieOeC and SieMeOeC Ceramics

Si

In the case of polysiloxanes, crosslinking is achieved via condensation, transition metal-catalyzed addition reactions (e.g. hydrosilylation) as well as free radical initiation techniques. Polymers containing methyl or vinyl groups can be crosslinked thermally by using peroxides [205]. Furthermore, silanol groups can be reacted with a moisturesensitive silane crosslinker. Silicon hydride functionalities in combination with vinyl groups allow for hydrosilylation reactions and the synthesis of infusible materials, which are resistant toward hydrolysis and thermal degradation (Scheme 24) [206,207]. The crosslinking process of polysiloxanes comprising functional groups, for example, hydroxy or alkoxy groups, can be achieved by the condensation of silanol groups with in situ water release and subsequent hydrolysis reactions of the alkoxy substituents and SieOeSi bond formation (Scheme 24). Using appropriate catalysts, such as a tetrakis(pentafluorophenyl) borate in the case of a polysiloxanol or [bis(2-ethylhexanoate)tin] in the case of poly(methoxymethylsiloxane) or poly(methylsiloxane), these reactions can occur already at room temperature [208,209]. The pyrolysis of crosslinked polysiloxane resins lead to the formation of amorphous silicon oxycarbide ceramics [210e212]. Mineralization occurs primarily through the evolution of hydrocarbons (essentially CH4) and hydrogen. In parallel, a variety of redistribution reactions between the various bonds SieO, SieC, SieH may occur (Scheme 25) [213], leading to the evolution of silanes (usually between 400 and 600  C), thus decreasing the final ceramic yield.

R R

Si

H C

(a)

R

R

Si

+ H

O

(b)

Si

R = Me: D; R = H: DH, DH2

Si

CH2 O n

OR + HO

Si

Si

O

Si

+ ROH

2 HO

Si

Si

O

Si

+ H2O

Si

OH + ROH

Si

OR + H2O

SCHEME 24 Crosslinking processes in polysiloxanes: (a) hydrosilylation reaction (as for vinyl-substituted polyhydridosiloxanes); (b) condensation and hydrolysis reactions.

At higher temperatures, that is, 600e1000  C, processes involving extensive CeH, SieC, and SieO cleavage leads to ceramic materials consisting of an amorphous silicon oxycarbide phase and residual free carbon. The structural investigation of the polymer-to-ceramic transformation has been essentially characterized by IR spectroscopy [214], solid-state NMR experiments (29Si, 13 C) [215e217], and TGA/MS [216]. The onset of thermal degradation strongly depends on the type of organic groups present in the polysiloxanes. The decomposition of their organic substituents on pyrolysis leads to the formation of a silicon oxycarbide phase consisting of SiCxO4x (0  x 4) sites. Their relative amount

O

O

cat.

Si

R

O

R = Me: M; R = H: MH, MH2, MH3

Si

H2C

n

R O

CH2

Si

O O

O

Si

O

O

R = Me: T; R = H: TH

T

O

D + D (Si-O/Si-O exchange) M+ T D + Q (Si-C/Si-O exchange) M+ Q

2D 2D 2T D+ T 2 TH

DH2 + Q

T + TH

DH + Q

DH2 + TH

MH3 + Q

DH + TH

MH2 + Q

MH3 + TH

SiH4+ Q

MH2 + TH

MeSiH3+ Q

successive Si-H/Si-O exchange SCHEME 25 Structural units in polysiloxanes and possible redistribution reactions thereof during thermal treatment.

Chapter | 11.1.10

Precursor-Derived Ceramics

can be extracted from the integration of the various resonance signals, which allows one to exactly determine the composition of the oxycarbide phase. A comparison with the elemental analysis allows for the estimation of the amount of segregated C phase present within the ceramic (so-called free carbon). 13 C MAS-NMR spectroscopy can also be used to follow the polymer-to-ceramic transformation: thus, SieCH3 groups were found to be stable up to 600  C, whereas at temperatures around 800  C, the formation of aromatic C sites can be observed. These findings are in agreement with those obtained by means of TGA coupled with mass spectrometry: the polymer-to-ceramic transformation is accompanied by the evolution of H2 and CH4 [160], which form due to the cleavage of SieC and CeH bonds. A large number of polysiloxane compositions have so far been studied [161,218]. It appears that the composition of the final SieOeC phase (SiO2xCx/2) can be predicted because the O/Si molar ratio remains almost constant during the pyrolysis process. Consequently, the elemental composition of the polysiloxane can be adjusted (C/Si ¼ [2 O/Si]/2) to finally obtain SiCO glasses with minimized contents of free carbon. This approach was applied by Soraru et al., who introduced a proper amount of SieH groups in the polysiloxane network. The authors modulated the content of free carbon and thus obtained a transparent SieOeC glasses [161]. Indeed, the concurrent presence of SieH and SieCH3 units promotes the formation of SieCH2eSi bonds, and thus the insertion of C into the SiCO framework [218]. The chemical modification of polysiloxanes with the main group and transition metal alkoxides has been shown to be a convenient access to SieMeOeC ceramics (M ¼ B [219,220], Al [221], V [222], Nb [223], Ta [224], Ti [225], Zr [163,226], Hf [164]). In the case of a chemical modification of polymethylsilsesquioxane (PMS) with hafnium tetra(n-butoxide), the polymer-to-ceramic transformation occurs differently compared to that of the nonmodified polysilsesquioxane. TGA/MS studies revealed the release of small amounts of water as well as of ethanol and butanol. The release of ethanol was attributed to the condensation reaction of SieOH and SieOEt groups (which are present as functional groups in PMS), while butanol is released via the condensation reactions of SieOH and HfeOnBu groups, thus revealing the chemical modification of PMS via hafnium alkoxide (Scheme 26) [164]. The same behavior was observed also for a zirconium alkoxidemodified polysilsesquioxane [226]. In both cases, FTIR spectra of the as-prepared alkoxidemodified polysilsesquioxane-based materials revealed the presence of SieOeM units (M ¼ Hf, Zr; absorption bands at 959 cm1 for Hf and 950 cm1 for Zr, Figure 15).

1063

Si

OH + EtO

Si

OH +

Si

- EtOH

n BuO Hf(O Bu)3

Si

n

- nBuOH

O

Si

Si

O

Hf(OnBu)3

SCHEME 26 SieOH/SieOEt and SieOH/HfeOnBu condensation reactions during crosslinking of the hafnium alkoxide-modified PMS.

The ceramic yield of the hafnium-alkoxide PMS (ca. 81%) was found to be lower than that of pure PMS (ca. 86%). This can be explained by the high mass loss induced by the condensation reactions and decomposition of Hf(OnBu)4. Considering its theoretical ceramic yield of 44.7% in combination with a ceramic yield of PMS of 86%, and weight fractions of 10.1% and 89.9%, respectively, a ceramic yield of 81.9% is expected, which is in agreement with the recorded value. The polymer-to-ceramic transformation of hafnium alkoxide-modified PMS was also investigated by means of MAS NMR techniques. The 29Si NMR spectrum of hafnium alkoxide-modified PMS (Figure 16) was found to be rather similar to that of the pure polymer. The 29Si signals expected for the silicon units ]Si(OEt)(CH3) and ]Si(OH)(CH3) as well as for the terminal ^SieCH3 groups appear at 65 (denoted by label e),57 (d) and18 (a) ppm, respectively. The observed experimental chemical shifts are in agreement with the data reported recently for PMS [227]. The NMR spectra of the SieHfeOeC ceramic nanocomposites are significantly broadened. At 800  C, the sample consists of an amorphous SiCxOy phase and the 29Si MAS NMR spectrum shows a distribution of silicon sites, namely, SiC3O (6 ppm, (a’)), SiC2O2 (25 ppm, (b’)), SiCO3 (65 ppm, (c’)), and SiO4 (106 ppm, (d’)). The formation of SiC2O2 and SiO4 units was seen as a result of rearrangement processes occurring during pyrolysis [228]. The 29Si NMR signal of tetrahedral SiO4 units (106 ppm) of the 800  C sample was found to be shifted by approximately 4 ppm to lower field compared with that of pure silicon oxycarbide prepared at the same temperature (110 ppm) [215,229]. This low-field shift was attributed to the presence of SieOeHf bonds within the ceramic backbone. It is generally accepted that modification of silicon oxides by the formation of SieOeM mixed bonds causes a low-field shift of the respective 29Si resonance; furthermore, the low-field shift increases with the number of M atoms in the second coordination sphere of silicon [230]. Downfield shifts up to 10 ppm have been reported, for example, for aluminosilicates [231], zeolites [232], mixed oxides SiO2eZrO2 [233,234], and SieOeCeZrO2 ceramics [226] prepared by solegel techniques.

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FIGURE 15 (a) FTIR spectra of PMS, and as-prepared PMS modified with 10 and 30 vol% hafnium alkoxide (arrows indicate the SieOeHf absorption band); (b) FTIR spectra of PMS and zirconium alkoxide-modified PMS.

An increasing pyrolysis temperature induces a phase separation within the material and the formation of silicaand SiC-rich domains. The SiO4 resonance increases in intensity and dominates the 29Si NMR spectra of the samples pyrolyzed at 1000 and 1200  C. In contrast, the SiCnO4n (n ¼ 1, 2, or 3) resonances decrease considerably in intensity SiCO3 (68 ppm, (a”)), or disappear completely (SiC2O2 and SiC3O). The spectrum of the 1300  C sample exhibits a SiO4 resonance with a chemical

shift of 109 ppm (c”’). The amount of mixed SiCO3 sites (b”’) decreases further, while a new signal at 12 ppm (a”’) evolves, corresponding to SiC4 units. This finding indicates the initiation of the phase separation process within the SieOeC matrix [235,236]. The information obtained by thermal analysis, FTIR and MAS NMR spectroscopy revealed that Hf/PMS up to 800  C consists of an amorphous SiHfOC phase (Figure 17). Especially the chemical shift of SiO4 centers

Chapter | 11.1.10

Precursor-Derived Ceramics

1065

and the presence of a SieOeHf band in the FTIR spectrum are strong indications for the presence of a single amorphous phase in the materials. At higher temperatures hafnia begins to precipitate as revealed by a slight high field shift of the SiO4 signal in the 29Si MAS NMR spectrum. The segregation of hafnia within the SiHfOC amorphous matrix was confirmed by transmission electron microscopy (TEM) studies. At temperatures >1300  C, additional phase separation of the silicon oxycarbide phase with the formation of amorphous silica and amorphous silicon carbide regions occurs, as indicated by the presence of SiC4 sites in the 29 Si MAS NMR spectra.

4. HIGH TEMPERATURE PROPERTIES 4.1. General Comments

FIGURE 16 29Si MAS NMR spectra of PMS modified with 10 vol% Hf(OnBu)4. The spectra at room temperature (r.t.) and 800  C were recorded under CP conditions, while the spectra at 1000, 1200, and 1300  C were recorded by means of SP experiments. Peaks labeled with asterisks are spinning sidebands. The insert is a magnification of the r.t. spectrum.

The choice of a material basis for high temperature structure materials suffers from a principal dilemma. Oxide materials show creep affinity at high temperatures, limiting their application in energy facilities and aerospace industry [237]. Nonoxide materials, which possess increased creep stability at high temperatures, are not thermodynamically stable in air and react with oxygen and/or moisture [238e240]. Those that form protective oxide layers either contain oxide type grain boundary phases promoting creep and increased internal oxygen diffusion at higher temperatures. Promising candidates with oxidation resistance and mechanical stability at high temperatures are covalent ceramics, which are prepared without sintering additives. According to the above sections, such materials can be obtained by the thermolysis of suitable organometallic precursors in an inert gas atmosphere [241e243]. Due to low self-diffusion coefficients in covalent ceramics, conventional sintering of powder compacts without

+ Hf(OR)4

FIGURE 17

Structural evolution of hafnium-alkoxide-modified PMS during polymer-to-ceramic transformation.

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Handbook of Advanced Ceramics

additives is not suitable for the production of composites based on covalent ceramics. Multinary nonoxide ceramics are of special interest because these materials combine the properties of the constituting binary phases. In this context, amorphous SieBeCeN ceramics obtained from organosilicon precursors became of fundamental interest, because these materials recently demonstrated a short-time thermal stability up to 2000  C without mass loss and only minor crystallization as will be discussed at the end of this section. These attributes, coupled with a high resistance toward oxidation even at elevated temperatures, give rise to extensive studies.

4.2. Thermodynamic Aspects A major aspect in understanding the high temperature behavior of precursor-derived amorphous ceramic composites is thermodynamics, that is, the consideration of phase reactions of the constituting thermodynamically stable unary, binary, or ternary phases. Solid-state reactions of such phases may in the worst-case result in the decomposition of the ceramic composites. For example, the stable and metastable solid phases of the quaternary system SieBeCeN are shown in Figure 18. Aside from the pure components silicon, boron, carbon, and nitrogen, a number of well-characterized binary phases such as SiC, Si3N4, BN, B4þdC as well as intermetallic silicon boron phases are known. Besides that, several metastable binary and ternary phases have been revealed in the subsystems. For example, in the binary CeN system, b-C3N4 has been predicted by Liu and Cohen [246]. In the system SieCeN b-SiC2N4, which decomposes >900  C and Si2CN4 (see also Sections 2.3 and 3.2), which is stable up to around 1000  C, have been reported [95]. Crystalline, B SiBn SiB6

B4+δC

SiB3 BN

B2CN2

N

BC2N BC3N BC4N

Si3N4

“C3N4“ SiC2N4

Si2CN4

SiC C

Si

FIGURE 18 SieBeCeN concentration tetrahedron including stable and metastable binary and ternary solid phases in the quaternary SieBeCeN system [244,245]. Note that the compositions BCxN are stoichiometric but not crystalline; C3N4 was calculated in detail [246] but so far not isolated in crystalline form.

phase-pure ternary BeNeC phases are to our knowledge not known so far. However, because of the almost similar lattice parameters of h-BN and graphite (both are layered materials with sheets of trigonally coordinated atoms) they build solid solutions. Stoichiometric (but not phase-pure) compositions of BCN shown in Figure 18 have usually been obtained by the reactions of trichloroborane with N-containing hydrocarbons such as acetonitrile, polyacrylnitrile, CVD processes using boranes, ammonia, and hydrocarbons or by thermolysis of boraneepyridine complexes [247e255]. Moreover, several amorphous compositions have been disclosed in as-thermolyzed PDCs, for example, amorphous carbon, amorphous silicon nitride, and amorphous silicon carbonitride [57,256e260]. A set of consistent thermodynamic data of the stable phases of the quaternary SieBeCeN system has been developed [244,245,261], which enables a detailed description of phase equilibria. With respect to possible applications of precursorderived nonoxide ceramics, SieCeN is the most important subsystem in the quaternary system. It was stated above that such ternary compositions can be obtained by the thermolysis of polysilazanes or poly(silylcarbodiimide)s at 1000e1400  C. As-obtained SieCeN materials are usually amorphous. In other words, they are in a nonequilibrium state, where phase configuration cannot be reproduced by equilibrium calculations. However, spectroscopic methods such as solid-state MAS NMR, and IR spectroscopy (IR) in combination with wide-angle scattering reveal a partial segregation of the amorphous ceramics into structural units that refer to thermodynamically stable phases. At higher temperature, segregation proceeds and the amorphous state is transformed into crystalline composites, consisting of Si3N4, SiC, and/or C, depending on the molecular structure and composition of the precursor and the conditions applied during thermolysis. An important tool for predicting thermally induced decomposition reactions are calculations of phase equilibria by the CalPhaD (Calculation of Phase Diagrams) approach using the set of thermodynamic data mentioned above. A number of thermodynamic calculations in the SieCeN system have already been published [244,245,262e265]. It turned out that the decomposition of such ceramics is to a significant extent controlled by the decomposition reactions of silicon nitride either by a carbothermal reduction according to the following. Si3 N4 þ 3C

1484  C

!

3SiC þ 2N2

ð1013 mbar N2 Þ (50)

or by decomposition into the elements: Si3 N4

1841  C

/

3Si þ 2N2

ð1013 mbar N2 Þ

(51)

Chapter | 11.1.10

Precursor-Derived Ceramics

1067

FIGURE 19 Partial pressure diagram for C/SiC/Si3N4 composites. Of major interest is the dependency of the decomposition temperature of Si3N4 and C with the formation of SiC and N2. At 1 bar p(N2), this temperature is 1484  C (1757 K), whereas at 10 bar p(N2), this value is shifted to 1700  C (1973 K). Assuming 104 bar in argon atmosphere, the decomposition of such a composite material should occur at 954  C (1227 K) [265].

It can be concluded from Figure 19 that the onset of the above decomposition reactions strongly depends on the nitrogen partial pressure. In 1 bar N2 atmosphere, the above reactions take place at 1484  C (1757 K) and 1841  C (2114 K), respectively. Increasing the nitrogen partial pressure to 10 bar results in a shift of the decomposition temperature due to carbothermal reduction to as much as 1700  C (1973 K), whereas at 1 mbar, one woulddconsider the results of the CalPhaD computationsdexpect decomposition already at around 1050  C. Thermochemistry studies have also been performed on polymer-derived SieOeC ceramics to assess their thermodynamic stability relative to crystalline phases. Calorimetric measurements of heats of dissolution in a molten oxide solvent indicated that these ceramics possess a negative enthalpy of formation (15 to 72 kJ/mol)

relative to their crystalline constituents (Table 1) [266,267]. Therefore, their crystallization resistance does not only rely on kinetic reasons but also on the thermodynamic characteristics. In contrast, calorimetric measurements performed on poly(silylcarbodiimide)-based ceramics provided evidence that amorphous carbon-rich SieCeN PDCs have slightly positive or close to zero heats of formation relative to graphite and the binary crystalline components Si3N4 and SiC (Table 1). Therefore, SieCeN PDCs are less stable than SieOeC ceramics. However, amorphous ceramics have a higher entropy than the corresponding crystalline phases, which may stabilize the amorphous PDCs at high temperatures considering the relatively small (positive) enthalpies of formation in Table 1. Therefore, the experimentally observed thermal stability of the SieCeN ceramics may still reflect at least in part thermodynamic stability.

TABLE 1 Enthalpy of formation from elements as well as from carbon and crystalline binary systems (SiO2 and SiC for SieOeC; Si3N4 and SiC for SieCeN) for one SieCeNe and two SieOeC-based ceramics pyrolyzed at 1100  C in inert atmosphere [268e271] DHf from elements at 25  C [kJ/mol]

Sample

Composition

Si1O1.1C0.73

SiO2 [mol%]

SiC [mol%]

C [mol%]

43.10

34.98

21.92

SiO2 [mol%]

SiC [mol%]

C [mol%]

14.73

25.97

59.30

Si3N4 [mol%]

SiC [mol%]

C [mol%]

Si1O0.72C2.10

Si1C2.80N1.20

14.30

4.80

80.90

DH From binaries and carbon at 25  C [kJ/mol]

242  3.6

53.9  3.6

149.9  4.4

51.9  4.4

50.6  6.4

2.8  6.4

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In conclusion, the unique thermal stability of the PDCs and their exceptional crystallization resistance relies on (i) kinetic stabilization by the “free” carbon, which suppresses diffusion processes and (ii) thermodynamics that is, negative enthalpy formation of (amorphous) PDCs compared to their crystalline binary counterparts.

(a)

N VT50

1414°C

NCP 200

1484°C Gas + Graphite + Si3N4

Si3N4

4.3. High Temperature Properties of SieCeN Ceramics As an example to illustrate the influence of chemical composition on high temperature properties, the decomposition behavior of two different SieCeN ceramics obtained by the thermolysis of commercially available PVS VT50 [57] and polymer-hydridosilazane NCP200 [59], which possess different chemical compositions, is briefly discussed. In Figure 20, calculated isothermal sections in the ternary SieCeN system are shown. Figure 20a shows phase equilibria between 1414 and 1484  C. In this temperature range, three three-phase fields exist: C/Si3N4/ Gas, C/Si3N4/SiC, and Si3N4/SiC/Liquid. Below 1414  C (melting point of silicon), the latter is composed of Si3N4/ SiC/Si. The composition of VT50 ceramic (symbolized by a triangle) is located on the tie-line Si3N4eC. Presuming quantitative segregation into the pure (crystalline) phases, the sample is composed of Si3N4 and graphite. In contrast, the composition of NCP200 ceramic is located in the three phase fields C/Si3N4/SiC. Accordingly, a fully crystalline sample would be composed of SiC, Si3N4 and graphite. CalPhaD predicts that raising the temperatures to >1484  C (Figure 20b) initiates the carbothermal reduction of Si3N4 according to Eqn (50). The expected quantitative loss of nitrogen in the case of VT50-derived ceramics should result in the formation of a SiC/C composite. In contrast, thermally induced degradation of ceramics obtained from NCP200 (symbolized by a rhombus) should deliver a composition, which is located on the tie-line Si3N4/SiC. The calculations are confirmed by XRD of annealed ceramic samples. A comparison of XRD patterns of VT50 and NCP200-derived ceramics after annealing at 1800  C in a nitrogen atmosphere (5 h, 1 bar) is shown in Figure 21. Whereas in the XRD pattern of annealed VT50 ceramic only b-SiC reflections appear, the pattern of annealed NCP200-derived ceramic exhibits both a-Si3N4 and b-SiC reflections. As expected, an increase of the temperature to 2000  C (not shown in the figure) results in the dissociation of a-Si3N4 into the elements. After cooling to room temperature, XRD patterns of the respective sample now clearly show a-Si reflections besides those of b-SiC. The contrasting crystallization of VT50 and NCP200 ceramics is also reflected in their different mass-loss

Si3N4 + SiC + Liquid

SiC + Si3N4 + Graphite

C

Si

SiC

(b)

N

1484°C Gas + SiC + Si3N4

1841°C

Si 3N 4

Si3N4 + SiC + Liquid

SiC + Gas + Graphite

C

Si

SiC

N

(c) 1841°C

SiC + Gas + Graphite

C

SiC + Gas + Liquid

SiC

Si

FIGURE 20 Ternary Si/C/N phase diagrams including compositions of ceramics derived from VT50 (triangles) and NCP200 (rhombs) at 1 atm N2 and various temperatures: (a) 1414  C < T < 1484  C; (b) 1484  C < T < 1841  C; (c) T > 1841  C. Decomposition of as-obtained VT50 ceramic >1484  C occurs in one step delivering SiC/C composites, whereas decomposition of NCP200-derived ceramic proceeds in two steps [57].

Chapter | 11.1.10

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1069

Si3N4. Due to the carbothermal reduction of Si3N4 at 1484  C, graphite is quantitatively consumed and bonded as SiC. Correspondingly, it is expected that the amount of Si3N4 is reduced to 38 mass-% with the formation of 12 mass-% of gas (N2). The second decomposition step at 1841  C consumes residual Si3N4. As a consequence, additional 15 mass-% of gas form along with 22% of a liquid (Si). As expected, the amount of SiC is not affected by this reaction.

4.4. High Temperature Properties of SieBeCeN Ceramics FIGURE 21 XRD patterns of VT50 and NCP200-derived ceramics after annealing to 1800  C in 0.1 MPa N2 for 5 h [269].

behavior. A suitable method for investigating mass lossaccompanied degradation is high temperature TGA. Figure 22 shows a comparison of the TGA of VT50 (top) and NCP200 ceramics (bottom). It is evident from the TGA curves that VT50 decomposes in a one-step reaction with the loss of 30% of its original sample weight. The determined value exactly fits with the nitrogen content of the ceramic determined by elemental analysis and confirms the results of the CalPhaD calculations shown in Figure 20. The onset of decomposition, which is at around 1500  C, is in accordance with the thermodynamic calculations. The corresponding DTA curve displays a strong endothermic peak, centered at around 1665  C. In contrast, NCP200-derived ceramics decompose in a two clearly separated steps. The onset of the first decomposition step caused by a carbothermal reduction of Si3N4 is at around 1550  C and thus shifted by 60  C to higher temperature compared with the result of the calculations. The second step is caused by Si3N4 dissociation into the elements and starts at around 1800  C. The mass loss of the individual decomposition steps is 12 and 15 mass-%, respectively. The total of 27 mass-% exactly matches the nitrogen content of the ceramic material. In addition, the values of the mass-loss of the single steps are purported by the amount of “free carbon.” It is interesting to note that indeed all as-thermolyzed ceramics with an overall composition within the three-phase field C/Si3N4/SiC decompose like VT50 or NCP200-derived ceramics in one- or two-step reactions, only depending on whether their C:Si ratio is >1 or <1, respectively. In this context, phase fraction diagrams are very suitable to quantitatively predict phase reactions. As an example, Figure 23 shows such a diagram of NCP200-derived ceramic. The total pressure was considered to be 1 bar. As-pyrolyzed NCP200-derived ceramic is composed of 8 mass-% of graphite, 19 mass-% of SiC, and 73 mass-% of

Organometallic polymers containing the elements silicon, boron, carbon, nitrogen, and hydrogen are precursors for silicon nitride, silicon carbide, and boron nitride-based SieBeCeN ceramics, which possess significantly improved thermal stability, and higher oxidation resistance compared with ternary SieCeN ceramics [120,124,133]. Results of detailed investigations of the high temperature properties of such materials by high temperature TGA, XRD, and TEM of thermally annealed samples will be shown below. A comparison of the thermal stability by HTTGA in an Ar atmosphere of an SieCeN ceramic derived from poly(vinylsilazane) (PVS) and an SieBeCeN material obtained from a boron-modified polysilazane is shown in Figure 24. It is clearly seen that PVS-derived ceramics start to decompose below 1500  C. As has been mentioned, this breakdown is primarily a consequence of the decomposition reaction of SieN units, which are present in the amorphous material together with “free” carbon causing the formation of silicon carbide and the elimination of nitrogen. SieBeCeN ceramics obtained from [B(C2H4SiHeNH)3]n [133], in contrast, resist thermal degradation up to 1950  C. This example nicely shows that by modification with around 5 mass-% of boron the onset of decomposition in SieCeN ceramics can be shifted from around 1500  C by roughly 450  C to higher temperature. From the thermodynamic point of view, the unusual high temperature stability of the SieBeCeN ceramic cannot be understood. According to the calculated phase fraction diagram shown in Figure 25 (composition Si: 24.0, B: 8.0, C: 44.0, N: 24.0 atom%), the SieBeCeN ceramic should at temperatures <1484  C, be composed of BN (16 mol-%), Si3N4 (28 mol-%), SiC (24 mol-%), and 32 atom-% of graphite [133]. In analogy to the abovementioned SieCeN ceramic, it is expected that in the case of the quaternary material, Si3N4 also reacts above 1484  C with graphite according to Eqn (50). Consequently, the original sample weight should decrease due to the loss of gaseous nitrogen. The amount of free carbon in this case decreases from 32.0 to 20.0 atom-%,

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Handbook of Advanced Ceramics

(a) -1800

5 0

-2000

TG -2200

-10 -2400 -15 -2600 -20

DTA

-2800

-25

DTA/μv endo

Mass loss (%)

-5

-3000

-30 -35

-3200

1200

1300

1400

1500

1600

1700

1800

1900

2000

Temperature (°C)

(b)

5

-2100

0

-2300

TG -2500

DTA -10

-2700

-15

-2900

-20

-3100

-25

-3300

-30

-3500 1200

1300

1400

1500

1600

1700

1800

1900

DTA/μv endo

Mass loss (%)

-5

2000

Temperature (°C) FIGURE 22

High temperature TGA/DTA investigations of VT50 (a) and NCP200-derived ceramics (b) in 0.1 MPa N2. Heating rate: 5  C/min [269].

whereas the amount of silicon carbide increases from 24.0 to 48.0 mol-%. However, from the high temperature TGA investigations (cf. Figure 24), it can be assumed that in contrast to the ternary ceramics, the thermodynamically expected decomposition reaction does not take place in the case of the SieBeCeN ceramic. In addition, boron, which is present as boron nitride, does not thermodynamically directly participate in any decomposition reactions <2000  C. Systematic investigations on the connection of chemical composition and thermal stability of precursor-derived SieBeCeN ceramics reveal that both the boron and

nitrogen content significantly influence thermal properties. By stepwise increasing the amount of boron in SieBeCeN precursor-derived materials, Mu¨ller et al. showed that at least 4 weight-% of boron are needed to efficiently avoid thermally induced decomposition [138,139]. For example, a series of SieBeCeN ceramics obtained from PVS, which was hydroborated with different amounts of borane dimethyl sulfide, is discussed in more detail in the following section. Table 2 gives the chemical composition of the ceramics derived from the above precursors in weight-%; values in

Chapter | 11.1.10

Precursor-Derived Ceramics

1071

FIGURE 23 Calculated SieCeN phase fraction diagram (ptot ¼ 1bar) for NCP200-derived ceramics [265]. A carbothermal reduction of Si3N4 at 1484  C (1757 K) results in the formation of SiC and N2 (Gas). Decomposition of remaining Si3N4 into the elements takes place at 1841  C.

Weight Change [%]

0

Si-B-C-N -10

-20

-30

Si-C-N

1250

1750

1500

2000

Temperature [°C] FIGURE 24 Comparison of the thermal behavior of ternary SieCeN and quaternary SieBeCeN ceramics [133] by high temperature TGA. Heating rate: 5  C, 1 bar argon.

FIGURE 25 Calculated SieBeCeN phase fraction diagram (ptot ¼ 1bar) for ceramics derived from [B(C2H4SiHeNH)3]n. Chemical composition: Si: 24.0, B: 8.0, C: 44.0, N: 24.0; atom-%. It is expected that due to the carbothermal reduction of Si3N4 decomposition of the ceramic composite occurs at 1484  C (1757 K), resulting in the formation of SiC and N2 (Gas) [133].

parentheses are atom-%. Phase compositions, which were obtained using the CalPhaD calculations on the basis of the results of the chemical compositions obtained by elemental analysis, are given in Table 3. Oxygen values are <1.5% and neglected. From the chemical composition of the ceramics, which was determined by elemental analysis, it is obvious that the increasing boron concentration in the precursors is directly reflected in an increasing amount of boron in the ceramic materials. It is the highest in 5c, which was derived from quantitatively hydroborated PVS 5 (cf. Scheme 27) and decreases in the row 4c > 3c > 2c. Accordingly, the concentration of Si, C and N increases in that manner. The results of the thermodynamic calculations in Table 3 point out that with increasing boron content the phase fraction of boron nitride increases from 7.3 mol-% in 2c to 18.2% in boron-rich 5c. Consequently, the amount of nitrogen available for the formation of silicon nitride is R´

CH 2

H C x

Si

N H

H 1

B R´ + BH3 * SMe2 n

- Me2S

CH 3

CH Si

N

H

H

H C

n

CH 2

Si

N

H

H

(x-3)n

R´= C 2H4Si(H)NH 2: 3: 4: 5:

x x x x

= = = =

8 5 4 3

Increasing Boron Content

SCHEME 27 Synthesis of SieBeCeN polymers with adjustable boron concentration [138].

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TABLE 2 Composition of ceramic materials 1ce5c (weight-%, atom-% in brackets) obtained from precursors shown in Scheme 27 [138] Ceramic

1c

2c

3c

4c

5c

Si

43.0 (26.8)

42.9 (26.0)

43.0 (26.0)

42.7 (25.3)

41.6 (24.7)

C

29.2 (42.5)

29.6 (41.9)

29.0 (41.0)

29.0 (40.2)

27.6 (38.4)

N

24.3 (30.3)

22.3 (27.1)

22.1 (26.8)

22.1 (26.3)

22.3 (26.6)

O

0.4 (0.4)

1.3 (1.4)

1.1 (1.2)

1.3 (1.4)

1.3 (1.4)

2.3 (3.6)

3.2 (5.0)

4.5 (6.9)

5.8 (9.0)

B

e

Empirical Formula


SiC1.6N1.1B0.00

SiC1.6N1.0B0.14

SiC1.6N1.0B0.19

SiC1.6N1.0B0.27

SiC1.6N1.1B0.36


O-content is neglected.

TABLE 3 Calculated relative phase fractions of crystalline 1ce5c materials (atom-%, O-content is neglected) and predicted (Dmp) and observed mass loss (Dmo at 1870  C) due to Si3N4 decomposition [138] Ceramic

1c

2c

3c

4c

5c

BN

0

7.3

10.2

14.0

18.2

Si3N4

53.3

41.6

38.5

34.2

31.3

SiC

8.1

17.0

19.6

22.2

23.3

C

38.6

34.0

31.7

29.5

27.2

Dmp

25%

20%

18%

17%

15%

Dmo

24%

19%

13%

8%

0.8%

reduced. This has two consequences: the silicon nitride phase fraction decreases from 41.6 mol-% in 2c to 31.3 mol% in 5c, whereas the silicon carbide phase fraction increases from 17.0 to 23.3%. The relative increase of silicon carbide again results in a decrease of the amount of graphite. In conclusion, increasing boron concentration results in increased boron nitride and silicon carbide but decreased silicon nitride and graphite phase fraction and vice versa. According to Scheme 27 and Tables 2 and 3, the total mass loss due to the carbothermal reduction of silicon nitride increases with increasing phase amounts of silicon nitride in the row 1c > 2c > 3c > 4c > 5c. This conclusion is confirmed by high temperature TGA as shown in Figure 26.

0 5c

weight change/%

-5

4c

-10

3c

-15

-20

2c 1c

-25 1000

1200

1400

1600

1800

2000

2200

temperature/°C FIGURE 26 Comparison of the thermal behavior of quaternary SieBeCeN ceramics with different boron concentration by high temperature TGA. Heating rate: 5  C/min, 1 bar argon. For details concerning the molecular structure of the polymeric precursors and the elemental and phase composition of the derived ceramics, consider Scheme 27 and Tables 2 and 3, respectively.

Chapter | 11.1.10

Precursor-Derived Ceramics

1073

1800°C α/β-SiC

+ 2c

β-Si3N4

+

+ +

++

++ + + + ++ + +

+

3c 4c 5c

10

20

30

40

50

60

70

80

2Θ FIGURE 27 Partial pressure diagram for C/SiC/Si3N4/BN composites. In contrast to the ternary C/SiC/Si3N4 system (cf. Figure 19), additional carbothermal reduction of boron nitride must be considered in the high temperature chemistry of the ceramic composites [133,269].

2000°C α/β-SiC

+ As expected, boron-free 1c starts to decompose at around 1550  C and loses 24% (¼Dmo) of its original sample weight. This is in accordance with the thermodynamically calculated (predicted) value of 25% (Dmp) given in Table 3. The incorporation of even small amounts of boron such as in 2c results in a shift of the onset of decomposition by 80  C to 1630  C. However, the total mass loss of 19% corresponds to the expected value (Dmp ¼ 20%). Further increased boron contents in 3c and 4c result in an additional shift of the decomposition temperature to around 1700  C. Moreover, a significant deviation of found (3c: Dmo ¼ 13%, 4c: Dmo ¼ 8%) and calculated (3c: Dmp ¼ 18%, 4c: Dmp ¼ 17%) mass losses is observed. Compound 5c, which possesses the highest boron content, shows no signs of decomposition in the expected temperature range. Thermal degradation is significantly retarded and starts at around 2000  C. At this temperature, carbothermal reduction of boron nitride is expected: 4BN þ C / B4 C þ 2N2

(52)

In analogy to the carbothermal reduction of silicon nitride (cf. Eqn (50)), the onset of BN decomposition also depends on the nitrogen partial pressure (Figure 27). The different behavior in the high temperature TGA shown in Figure 26 is mirrored in the crystallization/phase composition of the materials after an additional heat treatment at elevated temperature. For example, after heating to 1800  C (1 bar argon) for 5 h, XRD patterns of 2c and 3c correspond to the diffraction pattern of silicon carbide (Figure 28). This is in agreement with the high temperature TGA investigation that indicated quantitative (2c) or almost quantitative thermal degradation (3c) of silicon nitride at this temperature.

β- Si 3N 4 C

2c

+

+ +

++

+ ++ + + + ++ + +

3c 4c 5c

10

20

30

40

50

60

70

80

2Θ FIGURE 28 XRD patterns of the ceramics obtained from 2ce5c (cf. Scheme 27) after annealing to 1800  C (top) and 2000  C (bottom) in 1 bar argon for 5 h each [138]. It is evident that after heating to 1800  C in the case of less thermally stable materials 2c and 3c (cf. Figure 26), only reflections of silicon carbide are observed, whereas in the case of ceramics derived from 4c and 5c, additionally silicon nitride reflections are found. In the case of 4c, the latter vanish after annealing the material at 2000  C.

In the XRD patterns of the boron-richer materials 4c and 5c, b-Si3N4 reflections are distinctly detected besides SiC peaks. Their line widths suggest a nanocrystalline arrangement. At 2000  C, the thermodynamically expected cleavage of SieN bonds in 4c proceeds, resulting in the absence of Si3N4 reflections in the XRD diagram. In contrast, the diffraction pattern of 5c, which was shown by high temperature TGA to be the only thermally stable composite at this temperature, still represents b-Si3N4 reflections. In addition, all samples heat treated to 2000  C exhibit an additional broad reflection at 26.6 , which is a characteristic of graphitic carbon. Investigations on the influence of the nitrogen content on the thermal stability on SieBeCeN ceramics were also published [150]. Therefore, boron-modified poly (silylcarbodiimide)s with adjustable nitrogen content were

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Handbook of Advanced Ceramics

synthesized by a dehydrocoupling reaction of tris (hydridosilylethyl)boranes, B(C2H4SiRH2)3 (R ¼ H, CH3), and cyanamide (cf. Scheme 16) [151]. Investigations of the polymer-to-ceramic conversion of the precursors in an Ar atmosphere revealed that they deliver SieBeCeN ceramics in 65e85% yield, depending on the elemental composition and the molecular structure, that is, the crosslinking of the polymers. For example, TGA of a series of polymers obtained by the reaction of B(C2H4SiH3)3 [140] with different amounts of H2NeCN is shown (Figure 29). The calculated chemical composition of the precursors (idealized sum formula) is given in Table 4. Thermolysis of HeN1 e HeN6 releases ceramics in a 75e84% yield. The polymers with the lowest degree of crosslinking HeN2 and HeN3 have the highest ceramic yields of approximately 84%. The more highly crosslinked precursors HeN4 and HeN5 deliver ceramics in an 81%

TABLE 4 Idealized sum formulas of boron-modified silylcarbodiimide HeN1 and poly(silylcarbodiimide)s HeN1 e HeN6 obtained from B(C2H4eSiH3)3 and H2NeC^N [151] Silane: cyanamide

Sum formula (monomer unit)

HeN1

1:0.5

C6.5H20NSi3B

HeN2

1:1

C7H19N2Si3B

HeN3

1:1.5

C7.5H18N3Si3B

HeN4

1:2

C8H17N4Si3B

HeN5

1:2.5

C8.5H16N5Si3B

HeN6

1:3

C9H15N6Si3B

H-N3

Weight Change [%]

0 -5 -10

H-N2 H-N3.5

-15

H-N2 H-N1

-20

H-N5 H-N6

-25

250

500

750

1000

H-N4

1250

Temperature [°C] FIGURE 29 TGA of HeN1 e HeN6 (cf. Scheme 15). Heating rate 2  C/min, argon atmosphere [158].

and 80% yield, respectively, whereas the most highly crosslinked precursor HeN6, delivers ceramics in a 75% yield. Mass-spectra-coupled TGA revealed that neither silicon-, boron-, nor nitrogen-containing species evaporated during thermolysis. It can thus be concluded that the Si:B:N ratio which was adjusted in the precursors was maintained in the ceramics. High temperature properties of ceramics derived from HeN1 e HeN6 can clearly be correlated with their nitrogen content. Ceramics obtained from HeN5 and HeN6, which possess the highest nitrogen content among the materials investigated, decompose earliest at around 1670  C and finally loose 25% and 30% of their original sample weight. The decomposition temperature of HeN4 is shifted by 50  C to 1720  C, the final mass loss of 10% is significantly lower than that of HeN5 and HeN6. Compound HeN2, HeN3, and HeN3.5 behave like 5c ceramics (cf. Figure 26). Neither of these materials decomposes below 1950  C. The difference in the thermal mass loss behavior is again reflected in the different crystallization behavior of the ceramic materials. For example, the phase evolution observed by XRD of HeN2 and HeN5, which were annealed under similar conditions, is shown in Figure 31. As-obtained ceramics (1400  C) are predominantly amorphous. However, in the patterns of HeN2 ceramics, broadened reflections at 2Q ¼ 36 and 61 point to the presence of nanocrystalline silicon carbide. The appearance TABLE 5 Composition of ceramic materials HeN2 e HeN6 (a) weight-% and (b) calculated phase fraction in mol-% [mass-%]

Si

B

C

N

HeN2

37.9

6.0

37.2

18.9

HeN3

37.8

5.6

36.1

20.5

HeN3,5

37.2

4.8

35.7

22.3

HeN4

35.5

4.8

35.1

24.6

HeN5

33.1

4.5

35.3

27.1

HeN6

31.5

4.9

32.2

31.4

SiC

[mol-%]

BN

Si3N4

HeN2

17.5

21.9

23.7

36.9

C

HeN3

16.7

23.4

23.2

36.7

HeN3,5

14.0

31.8

14.6

39.6

HeN4

13.9

36.0

8.7

41.4

HeN5

12.9

41.1

1.2

44.8

HeN6

14.0

48.2

L6.8

44.6

Chapter | 11.1.10

Precursor-Derived Ceramics

1075

of this XRD pattern remains unchanged, even after heating the sample to 1800  C indicating that there is no further crystallization up to this temperature. After annealing to 1900  C, reflections of both a-SiC and b-Si3N4 appear. Remarkably, the latter do not decrease in intensity even after heating the samples to 2000  C. The crystallization behavior of HeN5 ceramic is significantly different and reflects the decomposition behavior observed by TGA shown in Figure 30. No crystallization of phase formation is observed <1600  C. Above this temperature, only b-SiC forms, whereas there is no evidence for the crystallization of silicon nitride. This finding again indicates the decomposition of SieN units due to a carbothermal reduction either before or during crystallization. The above examples clearly state that, under certain conditions, SieBeCeN ceramics are thermally much more stable, than thermodynamically predicted. To understand in detail the thermal stability of such ceramics and to find a universal relationship between the thermal stability and the compositions of the materials, different types of SieBeCeN ceramics were investigated in detail by means of TGA and XRD. Additionally, spectroscopic and microscopic methods were incorporated to receive information on the atomic structure of the amorphous and nanocrystalline assemblies. Thermodynamic calculations reveal that almost all compositions of SieBeCeN ceramics investigated so far are located in the four-phase equilibrium field Si3N4 þ SiC þ C þ BN (cf. Figure 18). Thermally stable compositions in this equilibrium field are located close to the three-phase equilibrium SiC þ C þ BN, whereas thermally less stable compositions are located close to the

H-N2

three-phase equilibrium Si3N4 þ C þ BN (Si3N4-enriched area). This fact is illustrated in Figure 32 schematically, showing a calculated isothermal section in the SieBeCeN system at a constant boron content of 10 atom-%. Structural analysis of as-obtained amorphous ceramics by means of small and wide angle scattering as well as by FTIR and solid-state NMR spectroscopy reveal a homogeneously ordered atomic array originating from the polymer structure. Remarkably, the structural units of the thermodynamically stable phases are already preformed within the amorphous states on an atomic and a medium range scale. Crystallization of these amorphous ceramics initially yields metastable amorphous phases Si3þy/4CyN4y (Y ¼ 0e4) and BN/C. The former is related to the composition line Si3N4eSiC (Figure 18). As already shown in the X-ray studies of annealed HeN2 and HeN5 in Figure 31,

H-N3.5

Weight Change [%]

0

H-N3

-10

H-N4 -20

H-N5 -30

H-N6

1200

1600 1800 1400 Temperature [°C]

2000

FIGURE 30 Comparison of the thermal behavior of quaternary SieBeCeN ceramics derived from HeN1 e HeN6 with different nitrogen concentrations by high temperature TGA. Heating rate: T < 1200 10  C/min, T > 1200 2  C/min, 1 bar argon atmosphere. For details concerning the molecular structure of the polymeric precursors and the elemental and phase composition of the derived ceramics, consider Scheme 16 and Table 5, respectively [150].

FIGURE 31 XRD patterns of the ceramics obtained from HeN2 and HeN5 (cf. Scheme 16) annealed at 1400e2000  C (100  C steps, 1 bar N2) each for 3 h [150].

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Handbook of Advanced Ceramics

N Thermally stable Thermally less stable

G + Si3N4 + SiC + BN C + Si3N4 + SiC + BN

Si3N4 + C + BN SiC + C + BN

Si + Si3N4 + SiC + BN SiC + C + BN + B4C

SiC + BN + B4C + L

Si

C

FIGURE 32 Calculated phase equilibria in the quaternary SieBeCeN system at 10 atom-% boron and a total pressure of 1 bar. Thermally stable SieBeCeN ceramics with compositions in the four-phase system C/Si3N4/SiC/BN are usually poor in nitrogen, whereas thermally less stable materials are often located close to the three-phase field system C/Si3N4/BN (on the tie-line CeSi3N4).

crystallization of such composites delivers both silicon nitride and silicon carbide. Detailed studies on the phase evolution of amorphous SieBeCeN ceramics by solid-state NMR using MAS were published by Schuhmacher et al. [203,204]. They compared the crystallization behavior of high temperature stable SieBeCeN ceramic, which was obtained from [B(C2H4SiHeNH)3]n (MW33 [133]) with that of a less stable material, derived from nitrogen-rich boron-modified poly(silsesquiazane) [B(C2H4Si(NH)1.5)3]n (MW60 [133]). The representative 29Si MAS NMR spectra given in Figure 33 cover the temperature range from 1050 to 2000  C. After thermolysis to 1050  C, 29Si NMR spectra

of both ceramics appear almost similarly (the NMR spectrum of the 1050  C MW60 sample is not shown). They are each characterized by a broad resonance centered at around 40 ppm. Apart from a line narrowing in the spectrum of MW60-derived ceramic, further heat treatment to 1400  C causes only minor changes in the signal appearance. The chemical shift range in combination with the broadened resonance signals at and below this temperature point to the presence of a mixture of SiCxN4x sites (x ¼ 1e4), suggesting a fully amorphous structure. After heating to 1600  C, significant progress in the phase evolution of both materials occurs. In the case of MW33-derived material, two new resonance signals centered at 19 and 49 ppm MW60

MW33 SiC4

SiN4

SiC4

FIGURE 33 Experimental 29Si NMR spectra of ceramics derived from boron-modified polysilazanes [B(C2H4SiHeNH)3]n MW33 (left) [133] and boron-modified poly(silsesquiazane) [B(C2H4Si(NH)1.5)3]n MW60 (right) [133] after annealing up to finally 2000  C (atmosphere: 1 bar nitrogen). All spectra were obtained under MAS conditions and single pulse excitation [202,204].

Chapter | 11.1.10

Precursor-Derived Ceramics

evolve, which indicate the formation of SiC4 and SiN4 sites, respectively. However, a broad signal, which appears in between, still points to a remarkable amount of SiCxN4x sites. In contrast, annealing of MW60 ceramic to 1600  C causes the disappearance of high field shifted signals. The only resonance signal in the 29Si NMR spectrum of MW60 ceramic is found at 19 ppm and suggests quantitative degradation of SieN motifs due to a carbothermal reduction with the formation of silicon carbide. Heating MW33 ceramics further results in additional demixing. As a consequence, only two, well-separated signals remain after heating to 1800  C, which exhibit considerably smaller line widths than those obtained from ceramics annealed at lower temperatures. It was claimed that this is a consequence of the presence of SiC and Si3N4 crystallites. Obviously, all structural components with silicon in mixed coordination, such as SiCxN4x sites, are disintegrated at this temperature, as also found by X-ray studies on the same samples. Even though obtained from structurally different precursors, the crystallization behavior of ceramics obtained from MW33 and MW60 shown in Figure 34 is comparable with that obtained from boron-modified polysilylcarbodiimides depicted in Figure 31. Ceramics with “high” nitrogen contents decompose around 1550e1600  C with crystallization of silicon carbide, whereas in the case of ceramics with “low” nitrogen content, neither thermal degradation nor crystallization is observed <1750  C. At and above this temperature, such materials transform into nanocrystalline silicon nitride/silicon carbide composites. Investigations of the structural evolution of SieBe CeN ceramics by high-resolution transmission electron microscopy (HR TEM) were also the subject of several publications. For example, Jalowiecki et al. investigated by HR TEM the microstructure of a boron-doped silicon carbonitride composite, which was derived from a borondoped polyhydridomethylsilazane [259]. As-obtained amorphous material was subsequently annealed at 1800  C for 50 h in an argon atmosphere. HR TEM investigations revealed that the annealed ceramic possessed a quite unusual microstructure. It was segregated into a turbostratic BNCx phase that occurred along the grain boundaries of nanosized silicon carbide and silicon nitride crystals as shown in Figure 35. It was supposed that the BNCx phase inhibits diffusion processes and consequently retards the crystal growth. Moreover, it binds free carbon, thus decreasing the carbon activity. The connection of carbon activity and decomposition temperature of silicon nitride due to the reaction with free carbon, which was calculated thermodynamically [264], is given in Figure 36. Obviously, a decreasing carbon activity shifts the onset of the carbothermal reduction of silicon nitride according to Eqn (50) to higher temperatures (at 1 bar N2, this reaction

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MW33

α/β-SiC β-Si 3N4

2000°C 1950°C 1900°C 1850°C 1800°C 1750°C 1700°C 1650°C 1600°C 1550°C 1500°C 1400°C 10

20

30

40

50

60

70

80

2Θ[°] MW60 α/β-SiC

2000°C 1950°C 1900°C 1850°C 1800°C 1750°C 1700°C 1650°C 1600°C 1550°C 1500°C 1400°C 10

20

30

40

50

60

70

80

2Θ[°] FIGURE 34 XRD patterns of the ceramics obtained from boron-modi(MW33, top) and fied polysilazanes [B(C2H4SiHeNH)3]n [B(C2H4Si(NH)1.5)3]n (MW60, bottom) after pyrolysis at 1400 and annealing at 1500e2000  C (50  C steps, 1 bar nitrogen) each for 3 h. In the case of the MW33-derived ceramics, a/b-SiC (n) and b-Si3N4 (l) reflections are observed, whereas in the case of MW60 ceramic only a/bSiC (n) reflections are found [133].

takes place at 1484  C). Presuming a carbon activity of 0.5 results in a shift by 90  C, whereas an activity of 0.1 would result in a shift of the decomposition temperature of around 330  C. Apparently, a further decrease of the carbon activity does not result in a further shift of the onset of Si3N4 decomposition to higher temperature since the decomposition of silicon nitride into the elements (at 1848  C), which is indicated by the horizontal line, becomes the dominant reaction. A quantitative encapsulation of the silicon nitride grains by the BNCx grain boundary phase presumably prevents the

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SiC

BNCX Si3N4

2 nm FIGURE 35 HR TEM image of SieBeCeN ceramics obtained from a boron-doped poly(methylsilazane) after annealing to 1800  C for 50 h in a nitrogen atmosphere. It is composed of SiC and Si3N4 nanocrystals that are separated by a turbostratic BNCx phase [259].

FIGURE 36 Calculated phase diagram displaying the relationship of decomposition temperature of silicon nitride due to carbothermal reduction vs. nitrogen partial pressure and carbon activity. At 1 bar N2 and aC ¼ 1 decomposition occurs at 1448  C. Assuming aC ¼ 0.2 results in a shift of this value to approximately 1700  C. Shifting the nitrogen partial pressure at aC ¼ 0.2 to 10 bar results in a further increase of the silicon nitride decomposition temperature to approximately1980  C. For color version of this figure, the reader is referred to the online version of this book.

release of nitrogen that forms by the silicon nitride degradation [265]. The correlating nitrogen partial pressure within this “shell” can thus be allowed to rise to higher values than atmospheric pressure. According to Figure 19, the onset of both the carbothermal reduction of silicon nitride and decomposition into the elements is directly connected to the partial pressure of nitrogen. Whereas at 1 bar the carbothermal reduction of silicon nitride occurs at 1484  C, a decomposition temperature of 1700  C can be calculated for a nitrogen pressure of 10 bar.

By combining the single topics discussed above, the unusual high temperature stability of SieBeCeN ceramics can be understood unequivocally. The BNCx grain boundary phase decreases the carbon activity and, due to an encapsulation effect, increases the internal nitrogen partial pressure considerably. Both effects result in a shift of the onset of the carbothermal reduction of silicon nitride to higher values, that is, to 2030  C, at 10 bar nitrogen pressure and aC < 0.2 (cf. Figure 36, upper line). However, as mentioned above, the stabilization of silicon nitride is only possible, if the amount of boron (i.e. the boron nitride phase fraction) exceeds a certain value, which was empirically found to be at around 8 atom%. This amount approximately corresponds to 16 mol% BN in the final ceramic, if derived from boron-modified polysilazanes [B(C2H4SiReNH)3]n or boron-modified poly(silylcarbodiimide)s [B(C2H4SiRe N]C]Ne)3]n. If the boron content in the materials is lower, there is obviously no sufficient dissolution of free carbon within the turbostratic BN phase. The second critical point is the relative phase amount of Si3N4. Even though there is a lack of more detailed studies, it turned out that the relative phase amount of silicon nitride should not exceed 30 mol%. If the respective value is higher, an efficient encapsulation of silicon nitride crystals that evolve during a heat treatment is not possible. Several publications report in detail on the high temperature mass stability of SieBeCeN by considering the relative ratios of (SiC þ BN þ C)/Si3N4, Si3N4/(Si3N4 þ SiC), and BN/ (BN þ C) [54,138,139,150,151]. However, this issue needs by far more detailed investigations and the development of new synthetic approaches to novel precursors for SieBeCeN ceramics, which allow for the preparation of ceramics with adjustable chemical compositions.

4.5. High-Temperature Properties of SieOeC and SieMeOeC Ceramics Silicon oxycarbide undergoes a phase separation at temperatures beyond 1100  C to form amorphous nanocomposites consisting of silica, silicon carbide and excess carbon [235,236]. Two processes occur in the SieOeC system: (i) crystallization of phase separated silicon carbide; (ii) carbothermal reaction of phase separated silica with excess carbon, accompanied by the formation of crystalline SiC and release of gaseous CO. Above 1500  C, severe decomposition of the SieOeC ceramic due to the reaction of silica with SiC to form silicon monoxide and CO is thermodynamically expected [270]. The incorporation of additional elements into the SieOeC network has a substantial effect on the decomposition and crystallization behavior. In SieBeOeC ceramics, silicon carbide crystallizes at lower temperatures than in boron-free SieOeC [271,272].

Chapter | 11.1.10

(a)

Precursor-Derived Ceramics

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FIGURE 37 (a) TEM image of a SieOeC/HfO2 ceramic nanocomposite pyrolyzed at 1100  C in argon atmosphere. Dark spherical particles of amorphous hafnia, which precipitated throughout the amorphous SieOeC matrix, are clearly visible; (b) Microstructure of a SieOeC/ HfO2 ceramic nanocomposites annealed at 1600  C; dark particles identified as HfSiO4 (hafnon) are well distributed throughout an amorphous SiO(C) matrix (reprinted with the permission of Wiley Blackwell) [273].

(b)

Recently, the high-temperature behavior of Zr- and Hfmodified SieOeC has also been investigated. The materials were prepared by the chemical modification of a polysilsesquioxane with metal alkoxide, followed by crosslinking and pyrolysis in argon atmosphere at 1100  C [163,164,273]. Crystallization and microstructural evolution occur in four temperature regimes: (i) at around 800  C the materials consist of a single amorphous SieMeOeC phase (M ¼ Zr, Hf); (ii) between 800 and 1100  C the metal oxide phase precipitates throughout the matrix; the metal oxide nanoparticles (particle sizes of 2e5 nm) are mainly amorphous (Figure 37a); (iii) >1100  C the SieOeC matrix starts to phase separate to silica, silicon carbide, and excess carbon, whereas the metal oxide nanoparticles crystallize to t-ZrO2 and t-HfO2; (iv) >1400  C a solid-state reaction between silica and the metal oxide phase occurs leading to the formation of a crystalline MSiO4 (M ¼ Zr, Hf) phase in the form of well-dispersed nanoparticles (20e30 nm) throughout the amorphous PDC matrix (Figure 37b) [163,273]. Systematic studies on the thermal stability of SieOeC/ ZrO2 and SieOeC/HfO2 ceramic nanocomposites upon annealing at T >> 1100  C have shown that the incorporation of the metal oxide phase within SieOeC strongly suppresses decomposition processes (e.g. carbothermal decomposition) even at temperatures as high as 1600  C. In contrast to the SieOeC sample, which exhibits a mass loss of 48.8% if annealed for 5 h in Argon atmosphere at 1600  C, significantly lower mass losses are recorded for the Hf-modified samples (Table 6). The modification of the polysiloxane with 30 vol% of zirconium or hafnium alkoxide leads to SieOeC/ZrO2 and SieOeC/HfO2

TABLE 6 Mass loss of SieOeC, SieOeC/HfO2 (two different contents of hafnia, as 10 vol% and 30 vol% hafnium alkoxide has been used for the chemical modification PMS) and SieOeC/ZrO2 (two different contents of zirconia, similar to SieOeC/HfO2) upon annealing at 1300, 1400, and 1600  C for 5 h in an argon atmosphere Mass Loss [%] Sample

1300  C

1400  C

1600  C

SieOeC

0.60

1.00

48.80

SieOeC/HfO2 (10 vol%)

1.19

1.88

18.28

SieOeC/HfO2 (30 vol%)

0.22

1.33

8.56

SieOeC/ZrO2 (10 vol%)

e

2.2

65.7

SieOeC/ZrO2 (30 vol%)

e

0.1

4.7

ceramic nanocomposites with mass losses of 4.7 wt% and 8.6 wt%, respectively, if annealed under similar conditions [163,273]. The strong improvement of the thermal stability of the ceramic nanocomposites by the incorporation of zirconia or hafnia into the SieOeC matrix relies on the suppression of the carbothermal reaction of silica with carbon. Instead, a reaction between silica and metal oxide occurs and leads to the formation of MSiO4 (see Scheme 28, as for SieOeC/ HfO2) [273]. SCHEME 28 Possible processes occurring in SieOeC/HfO2 ceramic nanocomposites upon annealing at 1600  C.

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Handbook of Advanced Ceramics

Preceramic Compound

- Ceramic/ Metal Powder + Preceramic Compound

- Milling Sieving Shaping

- Milling Sieving - Precursor aerosols - Gas phase Pyrolysis

- Melt Spinning - Dry Spinning

- Dip Coating - CVD

Infiltration

Ceramization

Monoliths

Sintering

Powders

Fibers

Infiltrated Material

Coatings

FIGURE 38 Preparation of ceramic materials by precursor processing.

The carbon contents measured in the samples obtained upon annealing at temperatures >1300  C (as for SieOeC/HfO2, see Table 7) support the fact that the carbothermal reaction of silica is suppressed [273]. Thus, the carbon content was shown to remain practically constant upon annealing, indicating that it is not consumed on reaction with silica.

5. APPLICATIONS

TABLE 7 Carbon and oxygen content of hafniamodified silicon oxycarbide (SieOeC/HfO2) annealed at 1300, 1400, and 1600  C Temperature

C [wt%]

O [wt%]

1300  C

9.78

41.60

1400  C

10.66

40.71

1600  C

9.70

41.35

5.1. General Comments The wide variety of available organometallic polymers, suitable as preceramic materials, offers exceptional opportunities for the development of PDCs. According to Figure 38, bulk parts, fibers, coatings, infiltrated media, fiber matrix composites, and near net shape manufactured ceramic components are available [3,243,274e276]. The preformsdalso referred to as green partsdcan be produced using techniques well known from polymer process engineering. The different approaches and techniques illustrated in Figure 38 will be discussed in the following sections topologically.

5.2. Monolithic Ceramics Bulk ceramics can be obtained by polymer powder compaction and subsequent ceramization according to the flow scheme shown in Figure 39. In contrast to conventional sintering processes, no sinter additives are required for the densification. Powder compaction is usually achieved by cold, warm, or hot isostatic or uniaxial pressing,

Polymer

Crosslinking Crosslinked Polymer

Homogenization, Powder Compaction Polymeric Monolith

Thermolysis Ceramic Monolith FIGURE 39 Flow scheme for the preparation of precursor-derived bulk ceramics.

Chapter | 11.1.10

Precursor-Derived Ceramics

depending on the physical/chemical properties of the precursor. A key step is the initial transformation of volatile or liquid precursors into highly crosslinked preceramic infusible networks. PDCs suffer from a high volumetric shrinkage (>50 vol percent) and residual porosity that evolves during pyrolysis of the preceramic polymers. This is a consequence of the loss of organic substituents and an increasing density from 1 g/cm3 (precursor) to approximately 3 g/cm3 (ceramic). Thus, the conversion into ceramic is accompanied by the generation of mechanical stresses that lead to defects and cracking of the ceramic parts [277]. Many studies in the last two decades were directed to reduce shrinkage and porosity by optimizing precursor compositions and finding suitable processing conditions, for example, with respect to crosslinking and ceramization. Fillers can be used to minimize the overall shrinkage (Figure 40). There are two types of fillers for this purpose. Passive fillers serve as a space holders, they are not involved in reactions during pyrolysis (their size and composition remain constant). In contrast, active fillers react during pyrolysis either with the polymer matrix, the decomposition volatiles, or the pyrolysis atmosphere, thus changing/tailoring the shrinkage and the composition of the final ceramic [184,185]. Inert fillers are, for example, metal oxides, carbides, or nitrides (e.g. Al2O3, SiO2, Y2O3, SiC, B4C, Si3N4, BN). The volumetric shrinkage of a system consisting of a preceramic polymer and an inert filler (uV(IF)) can be described as uV ðIFÞ ¼ ð1  VF =VF* ÞuV ðPÞ, where VF is the volume fraction of the inert filler, VF
is the critical volume fraction of the inert filler and uV(P) represents the volumetric

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shrinkage of the preceramic polymer. The critical volume fraction of the inert filler is the volume fraction at which the filler particles start to form a rigid network, that is, further shrinkage of the system is not possible. Hence, at volume fractions VF
, the total system shrinkage is zero. Active fillers are mainly used to produce near net shape parts because they nearly compensate the shrinkage of the precursors during pyrolysis. Usually, reactive powders are used, such as pure metals, intermetallics, metal hydrides or metal carbonyl complexes. The active filler usually reacts with the decomposition products or with the reactive pyrolysis atmosphere forming new phases that expand in volume. Therefore, the volumetric shrinkage of a system consisting of a preceramic polymer and an active filler, uV(AF), can be described as: uV ðAFÞ ¼ ½ð1  VF =VF* ÞuV ðPÞ þ VF uV ðFÞ, where uV(F) is the shrinkage/expansion value resulting from the filler transformation reaction. Thus, the critical volume fraction of an active filler, VAF
, to achieve net zero shrinkage in the bulk composite part is defined * ¼ u ðPÞ=ðu ðPÞ=V * Þ  u ðAFÞ [184,185]. as VAF V V V F The near net shape processing of PDC parts using active fillers has been denoted by Greil as Active Filler Controlled Pyrolysis [185]. Composite systems with near zero shrinkage can be designed with both passive and active fillers. However, in the case of active fillers, lower filler volume fractions are necessary to achieve zero shrinkage. The use of filler particles not only reduces shrinkage during pyrolysis but may also endow tailored properties to the composite. In this way, special properties such as mechanical strength, thermal and electrical conductivity, magnetism, or surface features can be adjusted.

FIGURE 40 Dense and crack-free green body (left) and ceramic part (right) prepared by crosslinking and pyrolysis of an alumina-filled polysilsesquioxane (reprinted with the permission of Elsevier) [278].

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Handbook of Advanced Ceramics

FIGURE 41 TMA of cold isostatically pressed green bodies obtained from a PVS [281].

20 Open Porosity [%]

The polymer powder compaction process for the preparation of PDC bulk parts by cold isostatic pressing was developed by Riedel et al. [60,279,280] in the early 1990s. The final density of ceramic materials obtained by this procedure was in the range of 90% maximum. For obtaining dense monolithic bodies, however, warm pressing of suitable (nonvolatile) precursors is the state of the art. Warm uniaxial pressing as described by Seitz and Bill of PVS VT50 [281] delivered glassy green bodies with 7.5% of open porosity. Subsequent thermolysis released ceramic bulk parts with relative densities of around 97%. Optimized conditions for this process were determined by means of thermomechanical analysis (TMA). Figure 41 shows such a diagram obtained from a cold-isostatically pressed PVS body (precursor: VT 50, Hoechst AG) [281]. TMA investigations give information on the shrinkage (or expansion) of materials upon heat treatment. If it is performed in an oscillating modus, it additionally allows for the determination of the softening temperature of the polymeric precursor (cf. inset in Figure 41). This value is of major importance because an increasing temperature results in a decrease in viscosity of the polymeric precursor, which enables a facilitated powder compaction. On the other hand, the maximum temperature applied for the powder compaction must be considered carefully. Exceeding this value (which correlates to the maximum of the TMA curve) results in the shrinkage of the green part due to a beginning ceramization during compression, causing the formation of shrinkage cracks which usually appear parallel to the pressing direction. As a consequence, shape integrity is lost, and in the worst case, the monoliths break into pieces. Haug et al. investigated in detail plastic forming of boron-modified poly(silylcarbodiimide)s for the production of SieBeCeN ceramic monoliths and correlated the conditions applied for the powder compaction with the open porosity of the derived bulk parts [282]. It turned out that the open porosity of bulk ceramics obtained from green

100°C

15

10 120°C

5

0 30 35 40 45 Pressure applied during Warm Pressing [MPa] FIGURE 42 Open porosity of ceramics derived from boron-modified poly(silylcarbodiimide) [B(C2H4Si(CH3)NCN)3]n [146,147]. Powder compaction was performed by plastic forming of the polymer powder at different temperatures and pressures and subsequent pyrolysis at 1400  C for 2 h [282]. For color version of this figure, the reader is referred to the online version of this book.

bodies prepared at 100  C/20 MPa could be decreased from about 20% to 12% by increasing the applied pressure to 43 MPa. If the same polymer was plastically formed at 120 , the porosity decreased from 6.6% at 20 MPa to 5% if 43 MPa were applied for the densification of the polymeric powder (Figure 42). It is obvious that the temperature applied during plastic forming of the precursors has more influence on the open porosity than the pressure applied. However, in the above case, further raising the temperature did not result in an additional reduction of the open porosity, because at and above 150  C, shape integrity was totally lost as is evident from Figure 43 [282]. The absence of low-melting grain boundary phases in precursor-derived materials leads to extraordinarily high temperature mechanical properties that can be very different from those of conventionally sintered ceramics.

Chapter | 11.1.10

140°C

Precursor-Derived Ceramics

150°C

FIGURE 43 Limits in the densification of bulk ceramics derived from boron-modified poly(silylcarbodiimide) [B(C2H4Si(CH3)NCN)3]n [146] by means of plastic forming. After polymer powder compaction at 140  C, sufficient open porosity is provided for releasing gaseous thermolysis byproducts, whereas after compaction at 150  C, the green body is too dense to efficiently release such gases [282].

Recent studies of the high temperature mechanical properties of as-thermolyzed SieCeN and SieBeCeN materials have shown that the amorphous state reveals an outstanding mechanical stability, even at rather high temperatures [283e285]. However, the deformation mechanisms are not yet well understood. Christ et al. characterized amorphous SieBeCeN ceramics using isothermal compression creep testing in the temperature range of 1200e1500  C [283,284]. It was found that the deformation rate contains a stress-independent section as well as a stress-dependent component that is proportional to the applied stress, indicating that this portion of the deformation mechanism is based on viscous flow. As shown in Figure 44 for SieBeCeN ceramics derived from [B(C2H4SiHeNH)3]n (MW33) and [B(C2H4SiMeeNH)3]n (T2-1), a continuous decrease in the deformation rate from approximately 106 to 108 s1 with time was found, which was mainly explained with a reduction of free volume in the amorphous material. Stationary creep was not observed, even after 450 h, indicating that the materials improve their creep stability during the creep tests: It is obvious and interesting to note that the curves of both materials can be divided into two ranges. The first one referred to as range I up to approximately 2
105 is similar

FIGURE 44 Deformation rates of amorphous ceramics derived from [B(C2H4SiHeNH)3]n (MW33) [133] and [B(C2H4SiMeeNH)3]n (T2-1) [124] during isothermal creep testing at 1400  C under a load of 100 MPa [283]. Atmosphere: air.

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to primary creep in conventional materials, which can often be described using the Norton power law [286]. Range II is characterized by an accelerated decrease of the deformation rate as compared to range I. Christ et al. [283,284] as well as Thurn et al. [287,288] also observed that the temperature dependence of the creep behavior of both SieCeN and SieBeCeN ceramics is comparably low and qualitatively the same for all temperatures. For example, the influence of the temperature on the deformation rate of T2-1-derived SieBeCeN ceramic is shown in Figure 45. Among the investigated samples, deformation rate was the lowest at 1350  C and the highest at 1500  C. The range I to range II transition occurred earlier at elevated temperatures. This shift was explained by a stress-independent reduction of free volume, which proceeds diffusioncontrolled and thus accelerates with increasing temperature. All creep tests mentioned above were performed in air and thus represent long time oxidation experiments. Once tested, the specimen usually possessed porous oxide surface layers consisting of silicon and oxygen as determined by energy-dispersive X-ray spectroscopy (EDX) and wavelength dispersive X-ray spectroscopy (WDX). Although the thickness of such oxide layers increased with testing time, a saturation value was reached. Although after short time exposures, even up to 1700  C, the thickness of the surface layers was <2 mm [133], the thickness of such layers was around 50 mm after annealing in air for 300 h at 1400  C and did not increase even after annealing for 500 h.

5.3. Ceramic Coatings Preceramic polymers solutions and slurries can be applied to surfaces to produce ceramic coatings. The coatings are prepared using various application techniques such as spin coating, spraying, or dip coating followed by crosslinking and pyrolysis. The choice of the coating procedure mainly depends on (i) slurry properties (composition, rheology, etc.), (ii) the substrate (size, shape complexity, roughness),

FIGURE 45 Deformation rates of [B(C2H4SiMeeNH)3]n (T2-1) [124] derived amorphous ceramics during isothermal creep testing at various temperatures under a load of 50 MPa [283]. Atmosphere: air.

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FIGURE 46 SEM micrographs of SieCeN-based coatings with a thickness of 20e30 mm, prepared via pyrolysis at 1100  C in inert atmosphere on alumina substrates: (a) without use of fillers; (b) with 4.2 vol% alumina; (c) with 15 vol% alumina. One can observe that the addition of fillers suppresses crack formation upon pyrolysis [298].

Handbook of Advanced Ceramics

(a)

(b)

(c)

and (iii) the desired coating thickness [289]. Thin layers can be deposited on planar, smooth surfaces by spin coating; the thickness can easily be adjusted by varying the spin rate [290,291]. Dip coating provides thicker layers and complex-shaped substrates can be coated via this technique [292,293]. Spray coating or screen printing can be used for layers with thickness up to 50 mm [294]. Independent of the procedure applied, one should note that there is a critical thickness (hCCT) above which delamination and cracking of the coating occur during pyrolysis [295,296]. With respect to delamination processes, hCCT depends on the adhesive strength between coating and substrate (wa), on the coating shear modulus (G) and on the shrinkage ratio a: hCCT ¼ wa =Gað1 þ a6  2a3 Þ. For thick coatings of up to several hundred micrometers, multiple layer deposition (with or without intermediate layer treatment) or deposition of polymer/filler particle systems may be applied to reduce shrinkage and cracking (Figure 46). Ceramization in an inert (nitrogen, argon) or reactive (ammonia, air) atmosphere may be used to furnish the PDC coatings. In the case of low melting substrates, advanced pyrolysis techniques such as laser pyrolysis or ion-beaminduced crosslinking and ceramization may be used. For instance, laser pyrolysis has been used to obtain ceramic coatings on plastic or aluminum substrates [297]. Furthermore, using focused laser beams, well-defined ceramic patterns on different substrates can be prepared [298]. PDC-based coatings can protect metal or carbon-based substrates from thermal loading and oxidative degradation [299,300]. Coatings consisting of SieOeC filled with TiSi2 particles can protect stainless steel surfaces toward

thermal oxidation up to 800  C. PDC-based coatings on silicon wafers [301,302], titanium, graphite, quartz [303], glass rods [304], alumina [305], and ceramic porous parts have a beneficial influence on the high-temperature behavior of the substrates used. Recently, silver-filled SieCeN-based PDC coatings have been shown to exhibit antibacterial functionality [306]. A multifaceted field is the protection of carbon-fiberreinforced ceramic composites from oxidation, hydrolysis, and erosion by PDC coatings. Because of their high maximum application temperature and low density, carbonfiber-reinforced composites are very promising materials for realizing mechanically chargeable, light weight structures with a high potential for aerospace applications [307]. However, under the harsh re-entry conditions, carbon fibers in such composites are oxidized and erode. As a consequence, shape integrity of exposed components is lost, which in the worst case results in the total damage of the composite. Figure 47 shows such a carbon fiber-reinforced silicon carbide substrate coated with a polymer-derived SieBeCeN ceramic after one coatingethermolysis cycle. The precursor was obtained by ammonolysis of B(C2H4SiMeCl2)3 according to Ref. [124]. Additionally, 42 vol.-% of Si powder was added to the polymer solution to avoid crack formation during the polymer-to-ceramic conversion [308,309]. It is evident that the ceramic coating homogeneously covered the surface of the substrate. The right part in the image in which the fiber structures are visible represents noncoated substrate, which was fixed to the dip-coating equipment. Figure 48 displays a comparison by TGA of the

Chapter | 11.1.10

Precursor-Derived Ceramics

FIGURE 47 Carbon-fiber-reinforced silicon carbide composite covered with a ceramic SieBeCeN coating. The coating was obtained by a dipcoating process using a slurry of 42 vol-% of silicon in a toluene solution of [B(C2H4SiMeNH)3]n [124] and subsequent thermolysis [308,309].

Weight Change [%]

0 -10 -20 -30 -40 -50

% (Si-B-C-N coated C/SiC) % (C/SiC Reference)

100

200 300 Time [min]

400

FIGURE 48 Comparison of the oxidation behavior at 1400  C in air of C/SiC composites with (squares) and without (circles) SieBeCeN coating. The mass loss of around 55% in the case of the uncoated CMC corresponds to the amount of carbon fibers in the material [308,309]. For color version of this figure, the reader is referred to the online version of this book.

oxidation behavior of the pure and the coated composite at 1400  C. Obviously, the ceramic coating efficiently prevented the substrate from degradation. Whereas an exposure of the uncoated substrate to air after 500 min resulted in the loss of approximately 55% of the original sample weight, there was nearly no mass change observed in the case of the protected substrate. The observed 55% mass loss roughly corresponded to the amount of carbon fibers in the substrate and was due to a total “fiber-burnout.”

5.4. Ceramic Membranes Microporous ceramic membranes are excellent candidates for gas separation applications. For steam reforming of natural gas, high-temperature reactors that operate at approximately 800  C are generally used. High-pressure membrane reactors operating at significantly lower

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temperatures (ca. 500  C) can potentially achieve conversion efficiency similar to that obtained in conventional reactors [310e313]. Furthermore, the membrane reactors have the advantage of synthesizing and purifying hydrogen gas simultaneously, thus leading to a highly efficient hydrogen production. Two types of membranes are used for high temperature membrane reactors: (i) dense palladium-based membranes and (ii) microporous ceramic membranes. Palladium-based membranes exhibit excellent hydrogen permeability. However, they are very expensive, susceptible to poisoning by sulfur and sensitive to cracking or pinhole formation due to hydrogen embrittlement. Microporous PDC membranes exhibit relatively high gas permeability, but it is their exceptional high temperature stability that makes them superior to palladium-based membranes [314]. Remarkably, the choice of the precursor and the technology applied for its deposition allow to control the microstructure development of the membrane and to adjust the desired microporous/mesoporous structural features [315]. Microporous amorphous-silica-based membranes can be prepared as thin films on permeable alumina porous supports having graded and layered porous structure, with pore sizes ranging from several hundred nanometers to a few nanometers [316,317]. The gas transport properties of amorphoussilica-based membranes deposited on mesoporous anodic alumina capillary were investigated. The gas permeability for small molecules (such as He or H2) was much higher than those for larger gas molecules. This behavior emphasizes the potential of these membranes for use in gas separation properties [318]. Oyama et al. synthesized amorphous-silicabased membranes by means of CVD techniques. Whereas He and H2 permeate through the membrane, the permeation of molecules with a kinetic diameter > 0.3 nm such as CO, CO2, or CH4 is suppressed. This emphasizes the great potential of these membranes in applications related to the hydrogen purification. An amorphous-silica-based membrane was also synthesized by pyrolysis in air of a polysilazane deposited on a silicon nitride porous support. It exhibited a hydrogen permeability of 1.3 108 mol/m2 s Pa at 300  C and a H2/N2 selectivity of 141, which is comparable with the selectivity of other amorphous-silica- or silicon oxycarbide-based membranes [319]. Amorphous ceramic membranes were also prepared in nonoxide systems, such as SieC, SieN, SieCeN, and SieBeCeN. The possibility of a molecular sieve amorphous SiC ceramic membrane was first demonstrated for a polysilastyrene derived composite membrane on a porous Vycor glass [320]. Amorphous silicon carbide ceramic membranes were prepared by thermal [317], e-beam [321], or chemical [322] curing of PCSs followed by pyrolysis in inert atmosphere. Hydrogen-selective SieOeC-based membranes were prepared via air curing of PCSs and subsequent pyrolysis in argon [323e325].

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FIGURE 49 Polymer-derived SieBeCeN ceramic membrane deposited on a porous alumina substrate [327].

Amorphous silicon-nitride-based ceramic membranes were prepared by the pyrolysis of a polysilazane in ammonia atmosphere at 650  C. The as-synthesized membrane showed a hydrogen permeability of 1.3 108 mol/m2 s Pa at 200  C and a H2/N2 selectivity of 165, whereas after hydrothermal treatment at 300  C, the permeability was >1.0 107 mol/m2 s Pa at 300  C with H2/N2 selectivity beyond 100. Recently, novel polymer precursors have been synthesized for the preparation of PDC membranes in the systems SieCeN and SieBeCeN. The SieCeN-based ceramic membrane was prepared via pyrolysis of a novel polysilylcarbodiimide precursor, which was synthesized via a nonoxide solegel process based on reactions of bis (trimethylsilyl)carbodiimide with chlorosilanes. Nitrogen sorption isotherm analysis of a 500 nm SieCeN layer deposited on a porous support surface indicated the existence of pores with sizes in the range 2e5 nm [326]. Amorphous SieBeCeN-based ceramic membranes with a thickness of approximately 1.75 mm were prepared by dip coating of a polyborosilazane on a macroporous alumina support, followed by pyrolysis in inert atmosphere (Figure 49) [327]. Pore analysis revealed a trimodal distribution of the pore sizes with maxima at 0.6, 2.7, and 6 nm. Materials based on SieCeN or SieBeCeN are highly promising amorphous nonoxide ceramics that can withstand extremely high temperatures and are, therefore, of great interest for hot gas separation or filtration applications.

5.5. Ceramic Fibers The fabrication of high-modulus refractory ceramic fibers requires controllable polymer rheology and adjustable polymer reactivity. Precursor fibers are obtained by either melt or dry spinning. Different principles for fiber processes such as extrusion, downdrawing from preforms,

Handbook of Advanced Ceramics

or updrawing from precursor melt or solutions are known [328]. Once spun, the precursor fiber is rendered infusible by subsequent curing, which represents an additional chemical surface reaction. Without this step, fiber integrity is lost because of melting and/or creep of the green fiber before full transformation into the ceramic state. The first nonoxide silicon-based ceramic fibers were obtained from poly(dimethylsilane) by the Yajima process [26]. Synthesis of polydimethylsilane through alkali-metalpromoted dehalocoupling of dichlorodimethlysilane and its thermal (Kumada) rearrangement into processable poly (methylsilylene-methylene), [HSi(CH3)eCH2]n, also referred to as polysilapropene, were already discussed in detail in Section 3.1. Polysilapropene fibers were obtained by melt spinning in an inert gas atmosphere such as nitrogen or argon at around 300  C. According to Figure 50, curing of the green fibers was performed by controlled surface oxidation in air at approximately 150  C, which resulted in the oxidation of SieH units accompanied with the formation of SieOeSi linkages on the fiber surface. Heating of such partially oxidized fibers in an inert gas atmosphere to 1200  C delivered SieCeO fibers, which are commercially available under the trademark NicalonÔ [329]. As evidenced from HR TEM and X-ray photoelectron spectra, such fibers are composed of free carbon, nanocrystalline b-silicon carbide and an amorphous nonstoichiometric SieCeO matrix. Heating of the fibers to temperatures exceeding the final thermolysis temperature results in their decomposition caused by the elimination of SiO and CO. A further crucial subject is the hydrogen content of the ceramic fibers which is progressively reduced when heating the fibers to 1200e1300  C or higher. All these effects are accompanied by a coarsening of the fiber morphology and result in H

PCS

Si CH2 Me

n

Spinning

PCS

Green Fiber Electron Beam Irradiation

Oxidation Cured Fiber Thermolysis Si-C-O Fiber NicalonTM Fibers

Si-C Fiber Hi-Nicalon Fibers

TM

FIGURE 50 Flow scheme of the preparation of SiC NicalonÔ and Hi-NicalonÔ fibers [26].

Chapter | 11.1.10

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a dramatic deterioration of the high temperature mechanical properties. Alternatively, electron beam (2 MeV) or g-ray (60Co) irradiation served for the chemical modification of the fiber surfaces, by which contamination of the fibers with oxygen could almost be avoided. The mechanism behind theses processes is based on radical reactions. Noticeably, free radicals, which remain in the green fibers, have to be removed by additional annealing steps under inert atmosphere to reduce the reactivity of the extremely sensitive radical species in order to avoid oxidation and/or hydrolysis when exposing the fibers to air. Commercially available silicon carbide fibers, which are obtained by melt spinning and subsequent electron beam curing, are, for example, HiNicalonÔ fibers [329]. Even though oxygen-free, these fibers also crystallize at comparably low temperatures (<1300e1400  C). In contrast to SieCeO fibers, the formation of nanocrystalline b-SiC and random segregation of carbon occur, whereas there is no evidence for an amorphous SieCeO phase. Thermodynamic calculations predict that in an inert atmosphere oxygen-free SiC/C, fibers are thermally stable up to approximately 2500  C. Unfortunately, such calculations only consider phase transformation or decomposition reactions that involve evaporation of gaseous species, whereas crystal growth cannot be taken into account. The latter is driven by a reduction of the surface energy and results in a grain coarsening with increasing temperature. Both silicon carbide grain coarsening and the increase of the size of the carbon domains negatively influence the mechanical properties of the fibers. A critical aspect is the oxidation behavior of SiC fibers [330]. At increased temperature, either passive or active oxidation occurs, depending on the oxygen partial pressure. Passive oxidation is accompanied with the elimination of CO or CO2 and appears at lower temperatures and high oxygen pressure. According to Eqn (53) or Eqn (54), it results in the conversion of (surface-near) silicon carbide into silica, which forms a protective layer on the ceramic fiber. SiC þ 2O2 / SiO2 þ CO2

(53)

SiC þ 1:5O2 / SiO2 þ CO

(54)

Even though carbon-containing species are evaporated in these reactions, the ceramic fibers gain in weight. In contrast, oxidation at low oxygen (partial) pressure according to Eqn (55) SiC þ O2 / SiO þ CO

(55)

results in the volatilization of carbon and silicon-containing species. It is consequently accompanied by a weight loss and in the worst-case results in the total damage of the ceramic fiber.

Recent publications also report on the synthesis and high temperature properties of silicon nitride-based ceramic fibers obtained from polysilazanes (see Section 3.2) or polyborosilazanes (Section 3.3). In previous sections, it was discussed that the derived ternary or quaternary ceramics resist crystallization up to temperatures of 1500  C and approximately 1800  C, respectively. This is mainly a consequence of the random amorphous ceramic network in which silicon atoms are present in mixed coordination spheres. Even though numerous SieCeN precursors have been published since the mid-1970s, there are only a very limited number of procedures described in the literature, which so far deal with the conversion of such polymers into ceramic fibers [197,331e337]. For example, SieCeN green fibers were obtained by the melt spinning of a polymeric silazane HPZ, which was synthesized by reacting trichlorosilane with HMDS [333]. The green fibers were chemically cured by exposing to a stream of argon/trichlorosilane at moderate temperature and subsequently thermolyzed at 1200  C. Alternatively, curing using trichloroborane served for rendering polymeric SieCeN fibers rendering infusible [336]. In contrast to as-obtained NicalonÔ fibers, which are composed of free carbon, nanocrystalline b-silicon carbide and an amorphous nonstoichiometric SieCeO matrix, SieCeN fibers derived from HPZ consist of an amorphous nonstoichiometric SiNxC1(3x/4) phase and free carbon. Remarkably, grain growth in SieCeN fibers is retarded compared to NicalonÔ fibers, but at temperatures exceeding 1500  C in an Ar atmosphere, degradation by carbothermal reduction of the SieN units leads to a total breakdown of the fiber. Baldus and Jansen obtained nonoxide silicon-nitridebased ceramic fibers, which do not decompose <1800  C and which were claimed to additionally possess superior oxidation resistance [117e120]. According to Eqn (30) (cf. Section 2.5.2), TADB was reacted with methylamine to deliver a polyborosilazane, which after crosslinking appeared glass-like but which was still soluble and fusible. The as-obtained precursor could be melt spun at 100e120  C using a multifilament spinneret, which delivers 200 single fibers. Subsequent curing and continuous heating to 1500  C released almost oxygen-free (<1 mass-%) ceramic fibers with an approximate elemental composition close to SiBN3C. Remarkably, oxidation experiments in air at 1500  C for 50 h resulted in the formation of an amorphous 1.7 mm SieCeO layer, which protected the fiber from further oxidation. Interestingly, tensile strengths of such oxidized fibers were identical to those of the as-obtained oxygen-free fibers. Bernard et al. reported in detail the preparation of hightemperature SieBeCeN fibers from polymeric precursors of the general type [B[C2H4SiR(NR’)]3]n (R ¼ H, CH3, (NR’)0.5; R’ ¼ H, CH3) [338e342]. The authors

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Handbook of Advanced Ceramics

FIGURE 51 TMA of [B[C2H4SiR(NMe)]3]n (1M(Me): R ¼ H; 3M(Me), R ¼ Me). [B[C2H4-SiH(NMe)]3]n does not soften below 160  C and is thus not suitable for melt spinning, whereas softening of [B[C2H4SiMe(NMe)]3]n takes place at 75e85  C [339,341]. For color version of this figure, the reader is referred to the online version of this book.

investigated the influence of the different substituents on viscoelastic properties, crosslinking chemistry, and structure as well as performance of the ceramic fibers. A means to determine the principal applicability of a precursor for extrusion processes, that is, spinning from the melt, is differential scanning calorimetry (DSC) in combination with TMA. The latter provides proper process conditions, whereas the former gives valuable information of the thermodynamics (i.e. melt-stability) under the desired processing conditions. For example, the results of TMA investigations of two structurally very similar precursors [B[C2H4SiR(NMe)]3]n (R ¼ H, Me) are shown in Figure 51. The polymers only differ in their silicon-bonded substituent R, which is either hydrogen (1M) or methyl (3M). Nevertheless, the different functionalities of these groups have a significant influence on the processability of the precursors. Compound 1M does not melt below 160  C and is thus not suitable for fiber spinning. Softening occurs at a significantly higher temperature, but as determined by TGA/DSC (not shown), simultaneous decomposition reactions occur. Thus, the melt does not possess sufficient stability for being melt spun. In contrast, 3M softens in the temperature range 75e85  C. Obviously, the siliconbonded methyl group provides sufficient flexibility to the polymer structure to allow a processing at ambient conditions in a temperature range, in which decomposition (>140  C), that is, a change in physical properties and chemical composition does not occur. Melt spinning was performed using a laboratory scale piston extrusion equipment and a windup device setup in a nitrogen-filled glove box. The polymer melt was forced through a 200-mm single-hole spinneret and the filament stretched to approximately 50 mm in diameter and collected

on the spool (50 m/min). To avoid degradation during the polymer-to-ceramic conversion and to maintain fiber integrity, the green fibers were subsequently cured in an ammonia atmosphere at 200  C for 60 min. Such a process requires sufficient latent reactivity of the polymer, which is provided by a small amount of silicon-bonded terminal NHMe units, which are replaced by NH2 groups in a transamination reaction. The newly formed NH2 units condense, thus forming SieNeSi bridges, which cause a surface hardening of the fibers. Pyrolysis in an N2 atmosphere releases ceramics in an approximately 30% yield. Remarkably, the yield may be increased to 57% when applying the abovementioned curing process. Figure 52 shows scanning electron micrographs of green fibers and as-obtained ceramic fibers. The fibers have smooth surfaces. Cracks or pores are not visible, neither in the green fibers nor in the ceramic fibers. During thermolysis, the fiber diameter decreases from approximately 50 mm to approximately 20e25 mm. This shrinkage corresponds to a decrease in volume of 75% (since the fibers were pyrolyzed on the spool on which they were collected, no shrinkage in fiber direction occurred). This value corresponds to the weight loss during thermolysis in combination with an increase in relative density from approximately 0.9 g/cm3 to approximately 2.1 g/cm3. Such SieBeCeN fibers possess excellent thermal stability. High temperature TGA in an N2 atmosphere points to the fact that no decomposition takes place at temperatures up to 1700  C. At 1750, degradation with volatilization of gaseous species and crystallization occurs, which accelerates at higher temperatures. The change in the fiber morphology with temperature can nicely be monitored by Scanning electron microscopy (SEM) (Figure 53).

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Precursor-Derived Ceramics

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FIGURE 52 SEM of melt-spun green fibers (left) of [B[C2H4SiMe(NMe)]3]n (3M) and ceramic fibers (right) obtained by subsequent curing in ammonia and thermolysis in an N2 atmosphere.

FIGURE 53 SEM of fracture surfaces of as-obtained ceramic 3M fibers (left) and after annealing in an N2 atmosphere at 1700  C, 1750  C, and 1900  C for 2 h [339].

The as-obtained 3M ceramic fiber shows typical glasslike fracture. The fracture surface is smooth, and there is no evidence for the presence of pores, cracks, or inhomogeneities. The microstructure is preserved after annealing for 2 h at 1700  C in a N2 atmosphere. An increasing temperature results in a beginning degradation. The fracture surface of a fiber annealed at 1750  C display a remarkable coarsening of the microstructure and extensive grain growth. The beginning degradation is also visible on the surface of the fiber, which is still smooth, yet it reveals certain irregularities. The fiber annealed at 1900  C is totally deteriorated. It is characterized by a very dense shell of approximately 0.5-mm thickness, which encloses a porous and crystalline-coarse interior. Such fibers are extremely brittle and no longer qualified as reinforcement in fiber-reinforced ceramics.

5.6. Fiber-reinforced Ceramics Fiber-reinforced ceramic matrix composites (CMCs or FRCs) have been shown to exhibit extraordinary thermomechanical properties combined with low density. They are consequently highly interesting candidates for manifold applications under severe conditions especially in aerospace [307]. CMCs can be obtained by different methods. The most established procedures are liquid phase routes, gas phase routes, or their combination, referred to as hybrid processes [307].

Liquid phase routes occur by infiltration of liquids such as melts, solutions, or slurries (which will not be discussed here in detail) into a fiber preform. The first one, usually a reactive melt infiltration, requires matrix materials with sufficiently low melting point, such as silicon (1410  C) or aluminum (660  C). During the infiltration step, the molten elements react either with the fiber preform, for example, carbonaceous fiber prepregs, or the atmosphere. Even though these methods allow for a rapid preparation of near net shape CMCs, there are some drawbacks, which have to be considered carefully. For example, the preparation of carbon-reinforced SiC composites is performed by capillary infiltration of liquid silicon. Molten silicon readily wets carbon and under the applied conditions reactsdon the expense of the carbon fibersdto SiC, even in the presence of amorphous carbon. To avoid fiber degradation, it is essential to protect the fibers, for example, by depositing boron nitride coatings, which is an additional timeintensive and expensive procedure. The preparation of SieC-based CMCs with inert fibers, such as Nicalon, is also possible. However, in this case, it is inevitable to carefully adjust the ratio of amorphous carbon in the prepreg and silicon, which is infiltrated in order to avoid segregation of residual nonreacted elemental silicon, since such a segregation usually worsens thermomechanical properties of the CMC dramatically. Precursor-derived CMCs, which can be obtained by polymer impregnation and pyrolysis (PIP) process, do not suffer from these drawbacks. Although rather simple, the

1090

process cannot be applied for every kind of polymeric precursor. Prerequisites for an effective polymer infiltration and loading of the preforms are adequate fluidity, that is, low viscosity of the polymer solution to allow for an efficient penetration into the preform, sufficiently high concentration of the solution to prevent extensive shrinkage during the drying step, appropriate crosslinking chemistry of the polymeric precursor, and high polymer-to-ceramic conversion yields. Even if a precursor fulfills all requirements, a large number of infiltration/thermolysis cycles may be required to obtain fully dense CMCs. Weinmann et al. published the first approaches for the fabrication of ultrahigh temperature CMCs with precursorderived SieBeCeN ceramic matrix using two-component precursors. In contrast to the abovementioned system, this route avoids the use of solvents [142e144]. According to Eqn (35) (Section 2.5.3), the polymeric matrix is obtained by a thermally induced hydrosilylation of OVS, with tris(hydridosilylethyl)boranes, B[C2H4eSi(CH3)3nHn]3 (n ¼ 1e3, in the literature referred to as Tris-SiH1eTrisSiH3). Remarkably, polymer synthesis occurs quantitatively without requiring catalysts and without the formation of byproducts. The physical/chemical properties of the starting compounds and the straightforward synthetic proceduredthe starting compounds were simply mixed and reacted at moderate temperaturedallowed for the preparation of fiber-reinforced SieBeCeN green parts using the RTM process. The single steps of the RTM process, as was performed for the preparation of the ceramic fiber-matrix composites, is shown schematically in Figure 54. Since the monomeric starting compounds are extremely sensitive to oxygen and/or moisture, special equipment, which is shown in Figure 55, was developed in order to allow for a handling of all substances under inert gas conditions. It consists of a Schlenk flask that serves as a reservoir for the precursor mixture and a brass mold, which contains the preforms. The flask and mold are linked via an interconnection glass tube. The valves attached allow separate evacuation of the two main parts. Initially, the brass mold is filled with the fiber material and the whole system is thoroughly evacuated. The central valve is closed and the Schlenk flask flooded with inert gas. The reaction mixture is introduced and the Schlenk flask evacuated again. After the central valve is opened, the precursor mixture is injected into the brass mold by slowly introducing argon through the valve attached to the Schlenk tube. After injection, the central valve is closed again and the brass mold heated within an oil bath to crosslink the precursors. The fiber-reinforced green body is removed from the brass mold within a glove box and subsequently pyrolyzed. Suitable conditions during the crosslinking process are of major importance with respect to obtaining a sufficiently dense matrix. Otherwise, an excessive amount of defects

Handbook of Advanced Ceramics

Monomers

1. Vacuum injection 2. Crosslinking Fiber-reinforced Greenbody

Thermolysis CMC

1. Reinfiltration 2. Thermolysis Reinfiltrated CMC FIGURE 54 Flow scheme of the preparation of fiber-reinforced ceramic matrix composites by vacuum injection method. Repetition of the reinfiltration-thermolysis process is required to obtain dense CMCs [142e144].

FIGURE 55 RTM vacuum injection equipment consisting of a brass mold that contains fiber fabrics or a three-dimensional weave and a Schlenk flask that serves as reservoir for the polymeric precursor [142e144].

will appear, degrading the mechanical properties of the final composite. The initially applied two-step heating program reported for the crosslinking chemistry of OVS and Tris-SiH2 (cf. Figure 56), that is, heating to 120  C within 2 h, isothermal annealing for 60 min and further heating to 200  C using a heating rate of 20  C/h, caused a foaming of the lowviscous reaction mixture. Accordingly, the fiber-reinforced composites displayed bubbles and voids and possessed an unsatisfactory green density (Figure 56a) [342]. For this reason, systematic experiments were performed to improve the crosslinking process with respect to

Chapter | 11.1.10

(b)

Crosslinking

(a)

Precursor-Derived Ceramics

Crosslinking time

FIGURE 56 Influence of the heating rate applied during crosslinking of OVS and Tris-SiH2 on the morphology of the precursor. (a) As-described earlier [142e144] and (b) modified procedure (straight line) [342]. For color version of this figure, the reader is referred to the online version of this book.

FIGURE 57 [342].

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obtaining a dense crack-free polymer matrix. Optimized heating schedules were explored empirically. According to the drawn-through line in Figure 56, a temperature regime was chosen, in which the samples were heated to 170  C within 5 h, isothermally annealed for 6 h, and subsequently heated to 175  C and 180  C (isothermal annealing each 3.5 h). Crosslinking was completed by heating the precursor to 200  C for 6 h. Finally, a dense glass-like and crack-free polymer matrix was obtained (Figure 56b). Thermolysis in an Ar atmosphere at 1400  C released ceramic matrix composites. Generally, their density is a function of various parameters such as the volume fraction of fiber and matrix, the efficiency of the impregnation step, weight loss during crosslinking and the ceramic yield of the precursor. However, even if assuming full impregnation and very high ceramic yields, a significant matrix shrinkage during pyrolysis appears, which is due to an increasing matrix density from approximately 1 g/cm3 (precursor) to approximately 2.1e2.5 g/cm3 (bulk ceramic). Accordingly, cracks and pores appear during the heat treatment. Such voids, however, can be subsequently filled by repetitive polymer impregnation and pyrolysis (PIP). The microstructure of a specimen obtained after 16 PIP cycles (89% relative density) is shown in Figure 57. It displays a dense, crack- and pore-free matrix. The integrity of the carbon fibers, which are visible as elliptical gray

SEM of a polished cross-section of a Cf/SieBeCeNmatrix composite after 16 PIP cycles. Dark: C fibers; bright gray: SieBeCeN matrix

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FIGURE 58 SEM of the lateral surface of the Cf/SieBeCeNmatrix composite after creep testing at 1400  C for 60 h applying a constant load of 100 MPa. Arrows indicate crack growth from the exterior of the sample, triangles point out transverse microcracks between the fiber bundles [342].

spots, is maintained (note: fibers are spherical but appear oval because cutting and polishing was not performed perpendicular to the fiber direction). Neither oxidation on the fiber surface nor within the fiber-matrix interface was visible. Moreover, it is evident that the voids between the individual carbon strands were filled by the matrix material with sufficient homogeneity. Bending strength of as-obtained FRC was 255 MPa. The thermal stability of the Cf/SieBeCeNmatrix composite was investigated by TGA and XRD. It retained its microstructure and mechanical properties up to 1500  C. After annealing at higher temperatures, that is, 1700, 1900, and 2000  C, a weight loss appeared and the matrix crystallized. The strength thereby decreased to <100 MPa. Brittle fracture of the FRC did not take place, regardless of the heat treatment conditions applied; the as-prepared sample and the sample investigated at 1500  C displayed a very well-developed fiber pull-out. In contrast to SieBeCeN bulk ceramics in which the creep rate under compression decreases continuously [283], the Cf/SieBeCeNmatrix composite possessed secondary creep up to at least 60 h at 1400  C and 100 MPa load. However, the creep strain, which measured 0.55%, was significantly lower than values typically measured for FRC made by chemical vapor infiltration (CVI) method. Most probably, transverse microcracks in the transverse bundles and interply microcracks mainly contributed to the creep strain (Figure 58).

6. CONCLUSIONS PDCs are unique materials with excellent properties such as high temperature stability against decomposition, crystallization, oxidation, corrosion, and creep. Furthermore, the

Handbook of Advanced Ceramics

chemical and phase composition as well as the microstructure and the properties of the PDCs can be tailored by appropriate designing of the preceramic polymers and by choosing specific crosslinking and pyrolysis/annealing conditions. In this chapter, the synthesis and properties of PDCs has been considered with an emphasis on the design of the preceramic polymers, their processing and their transformation into ceramics. Also the high-temperature behavior of PDCs has been reviewed and discussed in the light of the strong relationship between the architecture of the preceramic polymer and the microstructure and properties of the resulting ceramic. Despite the extensive research performed in the past decades, there are still open questions related for instance to the microstructure-property relations of PDCs. Thus, further studies are needed, which will require novel interdisciplinary approaches to both basic research and materials development.

REFERENCES [1] Colombo P, Mera G, Riedel R, Soraru GD. Polymer-derived ceramics: 40 years of research and innovation in advanced ceramics. J Am Ceram Soc 2010;93:1805e37. [2] Riedel R, Mera G, Hauser R, Klonczynski A. Silicon-based polymer-derived ceramics: synthesis, properties and applicationsda review. J Ceram Soc Jpn 2006;114:425e44. [3] Riedel R. From molecules to materialsda novel route for the synthesis of advanced ceramics. Naturwissenschaften 1995;82: 12e20. [4] Bill J, Aldinger F. Precursor-derived covalent ceramics. Adv Mater 1995;7:775e87. [5] Greil P. Polymer derived engineering ceramics. Adv Eng Mater 2000;2:339e48. [6] Colombo P, Riedel R, Soraru GD, Kleebe H-J, editors. Polymer derived ceramics. From nano-structure to applications. Lancaster, PA, USA: DeSTech Publications; 2009. [7] Bill J, Wakai F, Aldinger F, editors. Precursor-derived ceramics. Weinheim: Wiley-VCH; 1999. [8] Introduction to the Special Topical Issue on Ultrahigh-Temperature Polymer-Derived Ceramics. In: Raj R, Riedel R, Soraru GD, editors. Special issue on ultrahigh temperature polymer derived ceramics. J Am Ceram Soc, 84;2001:2158. and literature cited therein. [9] Babonneau F, Miele P, Riedel R, Soraru GD. Editorial. J Eur Ceram Soc 2005;25:89. Special Issue on “Polymer Derived Ceramics”. [10] Special triple issue on “preceramic polymers”. In: Riedel R, Ionescu E, editors. Soft Materials, 4; 2006. p. 105e299. [11] Birot M, Pillot JP, Dunogues J. Comprehensive chemistry of polycarbosilanes, polysilazanes, and polycarbosilazanes as precursors of ceramics. Chem Rev 1995;95:1443e77. [12] Kroke E, Li Y, Konetschny C, Lecomte E, Fasel C, Riedel R. Silazane derived ceramics and related materials. Mater Sci Eng, R 2000;26:97e199.

Chapter | 11.1.10

Precursor-Derived Ceramics

[13] Riedel R, Kroke E, Greiner A, Gabriel AO, Ruwisch L, Nicolich J. Inorganic solid-state chemistry with main group element carbodiimides. Chem Mater 1998;10:2964e79. [14] Papendorf B, Nonnenmacher K, Ionescu E, Kleebe H-J, Riedel R. Strong influence of polymer architecture on the microstructural evolution of hafnium-alkoxide-modified silazanes upon ceramization. Small 2011;7:970e8. [15] Riedel R. Handbook of ceramic hard materials. Wiley-VCH, ISBN 3-527-29972-6; 2000. and literature cited therein. [16] Suzuki K, Sasaki M. Pressureless sintering of silicon carbide. In: Somiya S, Bradt RC, editors. Fundamental structural ceramics. Tokyo: Terra Scientific; 1987. p. 75. [17] Tanaka H. Sintering of silicon carbide. In: Somiya S, Inomata Y, editors. Silicon carbide ceramics, vol. 1. Barking: Elsevier; 1991. p. 213e38. [18] Henshall JL, Rowcliffe DJ, Edington JW. Fracture parameters in Refel silicon carbide. J Mater Sci 1974;9:1559e61. [19] Narushima T, Goto T, Hirai T, Iguchi Y. High-temperature oxidation of silicon carbide and silicon nitride, materials transactions. JIM 1997;38:821e35. [20] Kipping FS, Sands JE. Organic derivatives of silicon. Part XXV. Saturated and unsaturated silicohydrocarbons, Si4Ph8. J Chem Soc 1921;119:830e47. [21] Miller RD, Hofer D, Rabolt J, Fickes GN. Anomalously longwavelength electronic transition in conformationally locked organosilane high polymers. J Am Chem Soc 1985;107:2172e4. [22] Miller RD. Polysilanesda new look at some old materials. Angew Chem Adv Mater 1989;101:1773e80. [23] Burkhardt CA. Polydimethylsilanes. J Am Chem Soc 1949;71: 963e4. [24] Chang LS, Corey JY. Dehydrogenative coupling of diarylsilanes. Organometallics 1989;8:1885e93. [25] Woo H-G, Walzer JF, Don Tilley T. A sigma-bond metathesis mechanism for dehydropolymerization of silanes to polysilanes by d0 metal catalysts. J Am Chem Soc 1992;114:7047e55. [26] Yajima S, Hayashi J, Omori M. Continuous silicon carbide fiber of high tensile strength. Chem Lett 1975:931e4. [27] Whitmarsh CK, Interrante LV. Synthesis and structure of a highly branched polycarbosilane derived from (chloromethyl)trichlorosilane. Organometallics 1991;10:1336e44. [28] Smith Jr TL. Process for the production of silicon carbide by the pyrolysis of a polycarbosilane polymer, US-4631179; 1986. [29] Wu H-J, Interrante LV. Preparation of poly(dichlorosilaethylene) and poly(silaethylene) via ring-opening polymerization. Macromolecules 1992;25:1840e1. [30] Procopio LJ, Berry DH. Dehydrogenative coupling of trialkylsilanes mediated by ruthenium phosphine complexes: catalytic synthesis of carbosilanes. J Am Chem Soc 1991;113: 4039e40. [31] Boury B, Carpenter L, Corriu RJP. Ein neuer Weg zu SiC-keramik-vorstufen durch katalytische synthese von prkeramischem polymer. Angew Chem 1990;102:818e20. [32] Boury B, Corriu RJP, Leclercq D, Mutin P, Planeix J-M, Vioux A. Poly(vinylsilane): a precursor to silicon carbide. 1. Preparation and characterization. Organometallics 1991;10:1457e61. [33] Suam JC, Speier JL. The addition of silicon hydrides to olefinic double bonds. Part III. The addition to non-terminal olefins in the presence of chloroplatinic acid. J Am Chem Soc 1958;80:4104e6.

1093

[34]

[35]

[36]

[37]

[38]

[39] [40] [41] [42] [43]

[44] [45]

[46] [47] [48] [49]

[50] [51]

[52]

[53]

Lay U. Alkinylsubstituierte Hauptgruppenelementverbindungen: Synthese, Reaktivita¨t und Verwendung als Werkstoffe. PhD Thesis, University of Heidelberg; 1991. Weinmann M. Vinyl-functionalized chlorosilanes as precursors for ceramics. Diploma Thesis, Germany: University of Heidelberg; 1992. Corriu RJP, Gerbier P, Gue´rin C, Henner B. From preceramic polymers with interpenetrating networks to SiC/MC nanocomposites. Chem Mater 2000;12:805e11. Kunai A, Toyoda E, Nagamoto I, Horio T, Ishikawa M. Polymeric organosilicon systems. 25. Preparation of branched polymers by regiospecific hydrosilylation of poly[(silylene)diethynylenes] and their properties. Organometallics 1996;15:75e83. Narsavage DM, Interrante LV, Marchetti PS, Maciel GE. Condensation polymerization of tetrakis(ethylamino)silane and its thermal decomposition to silicon nitride (Si3N4/silicon carbide (SiC)) ceramics. Chem Mater 1991;3:721e30. Schutzenberger H, Colson CR. Silicon. C R Hebd Seances Acad Sci 1885;92:1508e11. ¨ ber den Thermischen Abbau von SiliGlemser O, Naumann P. U ciumdiimid Si(NH)2. Z Anorg Allg Chem 1959;298:134e41. Seyferth D, Wiseman GH, Prud’homme C. A liquid silazane precursor to silicon nitride. J Am Ceram Soc 1983;66:C-13e4. Arai M, Sakurada S, Isoda T, Tomizawa T. Preceramic polysilazane to silicon nitride. Polym Prep 1987;28:407e8. Blanchard CR, Schwab ST. X-ray diffraction analysis of the pyrolytic conversion of perhydropolysilazane into silicon nitride. J Am Ceram Soc 1994;77:1729e39. Rochow EG. Polymeric methylsilazanes. Pure Appl Chem 1966;13:247e62. Funayama O, Arai M, Tashiro Y, Aoki H, Suzuki T, Tamura K, et al. Tensile strength of silicon nitride fibers produced from polyperhydridosilazane. Nippon Seramikkusu Kyokai Gakujutsu Ronbunshi 1990;98:104e7; Chem Abstr 1990;112. 103719. Wannagat U. Dreifach silyl-substituierte Amine. Fortschr Chem Forsch 1967;9:102e44. Bailey DL, Sommer LH, Whitmore FC. Some reactions of trialkylaminosilanes. J Am Chem Soc 1948;70:435e6. ¨ ber das Siliciumanalogon des KrisWannagat U, Brandmair F. U tallvioletts. Z Anorg Allg Chem 1955;280:223e40. Wilden J, Wank A, Asmann M, Heberlein JV, Boulos MI, Gitzhofer F. Synthesis of SieCeN coatings by thermal plasmajet chemical vapour deposition applying liquid precursors. Appl Organomet Chem 2001;15:841e57. ¨ ber Umsilylierungen an DisMooser JP, No¨th H, Tinhof W. U ilazanen. Z Naturforsch 1974;29b:166e73. Blum YD, Schwartz KB, Laine RM. Preceramic polymer pyrolysis: part 1 pyrolytic properties of polysilazanes. J Mater Sci 1989;24:1707e18. Seyferth D, Wiseman GH. X-ray diffraction analysis of the pyrolytic conversion of perhydropolysilazane into silicon nitride. J Am Ceram Soc 1984;67:C-132e3. Seyferth D. Organosilicon polymers as precursors for siliconcontaining ceramics, in: transformation of organometallics into common and exotic materials: design and activation, NATO ASI Ser. In: Laine RM, editor. E: Appl. Sci.-No., vol. 141. Dordrecht: Kluwer Publ; 1988. p. 133e53.

1094

[54] Weinmann M, Nast S, Berger F, Mu¨ller K, Aldinger F. Dehydrocoupling of tris(hydridosilylethyl)boranes with ammonia or amines: a novel route to SieBeCeN preceramic polymers. Appl Organomet Chem 2001;15:867e78. [55] Lu¨cke J, Hacker J, Suttor D, Ziegler G. Synthesis and characterization of silazane-based polymers as precursors for ceramic matrix composites. Appl Organomet Chem 1997;11:181e94. [56] Chong Kwet Yive s, Corriu RJ, Leclerq D, Mutin PH, Vioux A. Polyvinylsilazane: a novel precursor to silicon carbonitride. New J Chem 1991;15:85e92. [57] Bill J, Seitz H, Thurn G, Du¨rr H, Canel J, Janos BZ, et al. Structure analysis and properties of SieCeN ceramics derived from polysilazanes. Phys Stat Sol 1998;166A:269e96. [58] Laine RM, Sellinger A. Si-containing ceramics precursors. In: Rappoport Z, Apeloig Y, editors. The chemistry of organic silicon compounds, vol. 2. London: J. Wiley & Sons; 1998. p. 2245e316. [59] NCP 200, Nichimen Inc.; Chisso, Japan. For detailed investigations see also Ref. [57]. [60] Konetschny C, Galusek D, Reschke S, Fasel C, Riedel R. Dense silicon carbonitride ceramics by pyrolysis of cross-linked and warm pressed polysilazane powders. J Eur Ceram Soc 1999;19:2789e96. [61] NN710, Tonen Corp., Japan. [62] Weinmann M, Zern A, Aldinger F. Stoichiometric silicon nitride/ silicon carbide composites from polymeric precursors. Adv Mater 2001;13:1704e8. [63] Andrianov KA, Kotrelev GV. Synthesis and properties of organocyclosilazanes. J Organomet Chem 1967;7:217e25. [64] Voronkov MG, Lavent’yev VI. Polyhedral oligosilsesquioxanes and their homo derivatives. Top Curr Chem 1982;102:199e236. [65] Day VW, Klemperer WG, Mainz VV, Millar DM. Molecular building blocks for the synthesis of ceramic materials: [Si8O12](OCH3)8. J Am Chem Soc 1985;107:8262e364. [66] Calzaferri G, Herren D, Imhof R. Monosubstitution von octa (hydridosilasesquioxan) H8Si8O12 zu R0 H7Si8O12 mittels hydrosilylierung. Helv Chim Acta 1991;74:1278e80. [67] Ra¨ke B, Roesky HW, Uso´n I, Mu¨ller P. Synthese und Struktur von (CH3Si)6(NH)9: ein offener SieN-Ka¨fig aus Methyltrichlorsilan und Ammoniak. Angew Chem 1998;110:1508e10. [68] Blum YD, Laine RM. Catalytic methods for the synthesis of oligosilazanes. Organometallics 1986;5:2081e6. [69] Laine RM, Blum YD, Chow A, Hamlin RD, Schwartz KB, Rowcliffe DJ. Catalytic synthesis of novel polysilazanes including precursors to silicon nitride. Polym Prep 1987;28: 393e5. [70] Youngdahl KA, Laine RM, Kennish RA, Cronin TR, Balavoine GG. New catalytic routes to preceramic polymers: ceramics precursors to silicon nitride and siliconecarbide nitride. Mat Res Soc Symp Proc 1988;121:489e96. [71] Blum YD, Schwartz KB, Crawford EJ, Hamlin RD. New structures of perceramic polysilazanes synthesized by transition metal catalysis. Mat Res Soc Symp Proc 1988;121:565e70. [72] Liu HQ, Harrod JF. Dehydrocoupling of ammonia and silanes catalyzed by dimethyltitanocene. Organometallics 1992;11:822e7. [73] Seyferth D, Schwark JM. Stoichiometric and catalytic ring opening of hexaalkylcyclodisilazanes by organoalkali reagents. Organometallics 1989;8:1980e6.

Handbook of Advanced Ceramics

[74] Duguet E, Schappacher M, Soum A. High molar mass polysilazane: a new polymer. Macromolecules 1992;25:4835e9. [75] Bruzaud S, Soum A. Anionic ring-opening polymerization of cyclodisilazanes, 1. General aspects. Macromol Chem Phys 1996;197:2379e91. [76] Bruzaud S, Mingotaud A-F, Soum A. Anionic ring-opening polymerization of cyclodisilazanes, 2. Influence of additives on the kinetics of polymerization. Macromol Chem Phys 1997; 198:1873e84. [77] Cazalis C, Mingotaud A-F, Soum A. Anionic ring-opening polymerization of cyclodisilazanes, 3. Influence of the silicon substituent on the kinetics of polymerization. Macromol Chem Phys 1997;198:3441e50. [78] Bouquey M, Brochon C, Bruzaud S, Mingotaud A-F, Schappacher M, Soum A. Ring-opening polymerization of nitrogen-containing cyclic organosilicon monomers. J Organomet Chem 1996;521:21e7. [79] Ebsworth EA, Mays MJ. The preparation and properties of disilylcyanamide. J Chem Soc 1961:4879e82. [80] Ebsworth EA, Mays MJ. Zur Reaktion zwischen Halogensilanen und Silbercyanamid. Angew Chem 1962;74(3):117. [81] Pump J, Wannagat U. Bis-(trimethylsilyl)-carbodiimid. Angew Chem 1962;74:117. [82] Pump J, Wannagat U. Siliconenitrogen compounds. XIV. Bis (trimethylsilyl)cyanamide. Ann Chem 1962;652:21e7. [83] Birkofer L, Ritter A, Richter P. N,N0-bis-trimethylsilyl-carbodiimide. Tetrahedron Lett 1962;5:195e8. [84] Pump J, Rochow EG, Wannagat U. Chemistry of siliconenitrogen compounds. XXI. silyl-substituted carbodiimides. Monatsh Chem 1963;94:588e98. [85] Pump J, Rochow EG. Silylcarbodiimide. IV. Sila-polycarbodiimide. Z Anorg Allg Chem 1964;330:101. [86] Klebe JF, Murray JG. Organosiliconcarbodiimide polymers and process for their preparation, US-3352799; 1968. [87] Kienzle A, Obermeyer A, Riedel R, Aldinger F, Simon A. Synthese und Struktur des ersten oligomeren cyclischen Dimethylsilyl-substituierten Carbodiimids. Chem Ber 1993;126: 2569e71. [88] Obermeyer A, Kienzle A, Weidlein J, Riedel R, Simon A. Die Kristall- und Moleku¨lstrukturen von (Me2SiNCN)4 (1) und Me3SiNCNSiMe3 (2). Z Anorg Allg Chem 1994;620:1357e63. [89] Gabriel AO, Riedel R. Preparation of non-oxidic silicon ceramics by an anhydrous solegel process. Angew Chem Int Ed Engl 1997;36:384e6. [90] Gabriel AO, Riedel R. A novel nonoxide solegel process to SieCeN ceramics. In: Proc. 21st annual cocoa beach conference and exposition, The American Ceramic Society 1997, ISSN 0196e6219, Florida USA, 18; 1997 713e720. [91] Gabriel AO, Riedel R, Storck S, Maier WF. Synthesis and thermally induced ceramization of a non-oxidic poly(methylsilsesquicarbodi-imide) gel. Appl Organomet Chem 1997;11: 833e41. [92] Riedel R, Gabriel AO. Synthesis of polycrystalline silicon carbide by a liquid-phase process. Adv Mater 1999;11:207e9. [93] Gabriel AO, Riedel R, Dreßler W, Reichert S, Gervais C, Maquet J, et al. Thermal decomposition of poly(methylsilsesquicarbodiimide) to amorphous SiCN ceramics. Chem Mater 1999;11:412e20.

Chapter | 11.1.10

Precursor-Derived Ceramics

[94] Balan C, Vo¨lger KW, Kroke E, Riedel R. Viscoelastic properties of novel silicon carbodiimide gels. Macromolecules 2000;33: 3404e8. [95] Riedel R, Greiner A, Miehe G, Dreßler W, Fueß H, Bill J, et al. The first crystalline solids in the ternary SieCeN System. Angew Chem Int Ed Engl 1997;36:603e9. [96] Mera G, Riedel R, Poli F, Mu¨ller K. Carbon-rich SiCN ceramics derived from phenyl-containing poly(silylcarbodiimides). J Eur Ceram Soc 2009;29:2873e83. [97] Balan C, Riedel R. Rheological investigations of a polymeric precursor for ceramic materials: experiments and theoretical modelling. J Optoelectronics Adv Mater 2006;8:561e7. [98] Li Y-L, Kroke E, Klonczynski A, Riedel R. Synthesis of monodisperse spherical silicon dicarbodiimide particles. Adv Mater 2000;12:956e61. [99] Takamizawa M, Kobayashi T, Hayashida A, Takeda Y, Joetsu N. Borsiliciumorganisches Polymeres und Herstellungsverfahren dafu¨r, DE-3444306A1; 1983. [100] Takamizawa M, Kobayashi T, Hayashida A, Takeda Y. Method for the preparation of an inorganic fiber containing silicon, carbon, Boron and Nitrogen, US-4604367; 1986. [101] No¨th H. Beitra¨ge zur Chemie des Bors, XIII. Zur Spaltung der SieN-Bindung durch Lewis-acide Bor-Verbindungen. Z Naturforsch 1961;B16:618e20. [102] Narula CK, Schaeffer R, Paine RT. Synthesis of boron nitride ceramics from poly(borazinylamine) precursors. J Am Chem Soc 1987;109:5556e7. [103] Seyferth D, Plenio H. Borasilazane polymeric precursors for borosilicon nitride. J Am Ceram Soc 1990;73:2131e3. [104] Seyferth D, Plenio H, Rees Jr WS, Bu¨chner K. Silicon ceramics with a dash of boron. In: Bassindale AR, Gaspar PP, editors. Frontiers of organosilicon chemistry. Cambridge: The Royal Society of Chemistry; 1991. p. 15e27. [105] Blum YD, Laine RM. Catalytic process for making compounds having a non-lewis acid/base bond between a group IIIa metal and group VA nonmetal, US-4801439; 1989. [106] Blum YD, Laine RM. Process for making ceramic materials, US-5017529; 1991. [107] Su K, Remsen EE, Zank GA, Sneddon LG. Synthesis, characterization and ceramic conversion of borazine-modified hydridopolysilazanes. Polym Prep 1993;34:334e5. [108] Su K, Remsen EE, Zank GA, Sneddon LG. Synthesis, characterization, and ceramic conversion reactions of borazine-modified hydridopolysilazanes: new polymeric precursors to silicon nitride carbide boride (SiNCB) ceramic composites. Chem Mater 1993;5:547e56. [109] Wideman T, Su K, Remsen EE, Zank GA, Sneddon LG. Synthesis, characterization, and ceramic conversion reactions of borazine/silazane copolymers: new polymeric precursors to SiNCB ceramics. Chem Mater 1995;7:2203e12. [110] Wideman T, Su K, Remsen EE, Zank GA, Sneddon LG. New polymer precursors to SiNCB materials. Mat Res Soc Symp Proc 1996;410:185e90. [111] Wideman T, Cortez E, Remsen EE, Carroll PJ, Sneddon LG. Reactions of monofunctional boranes with hydridopolysilazane: synthesis, characterization, and ceramic conversion reactions of new processible precursors to SiNCB ceramic materials. Chem Mater 1997;9:2218e30.

1095

[112] Wideman T, Remsen EE, Cortez E, Chlanda VL, Sneddon LG. Amine-modified polyborazylenes: second-generation precursors to boron nitride. Chem Mater 1998;10:412e21. [113] Wideman T, Fazen PJ, Su K, Remsen EE, Zank GA, Sneddon LG. Second-generation polymeric precursors for BN and SiNCB ceramic materials. Appl Organomet Chem 1998;12:681e93. [114] Mu¨ller A, Gerstel P, Bunjes N, Berger F, Sigle W, Mu¨ller K, et al. Nanostructured SiC/BN/C ceramics derived from mixtures of B3N3H6 and jHSi(Me)C^C]. Int J Mat Res 2006;97:721e8. [115] Evans DA, Fu GC. The rhodium-catalyzed hydroboration of olefins: a mechanistic investigation. J Org Chem 1990;55: 2280e2. [116] Gervais C, Babonneau F, Maquet J, Bonhomme C, Massiot D, Framery E, et al. 15N cross-polarization using the inversionrecovery cross-polarization technique and 11B magic angle spinning NMR studies of reference compounds containing BeN bonds. Magn Reson Chem 1998;36:407e14. [117] Jansen M, Baldus H-P. Siliciumbornitridkeramik und Vorla¨uferverbindungen, Verfahren zu deren Herstellung sowie deren Verwendung. DE-4107108A1 1992. [118] Baldus H-P, Wagner O, Jansen M. Synthesis of advanced ceramics in the systems SieBeN and SieBeNeC employing novel precursor compounds. Mat Res Soc Symp Proc 1992;271:821e6. [119] Baldus H-P, Jansen M, Wagner O. New materials in the system Sie(N, C)eB and their characterization. Key Eng Mat 1994;89e91:75e80. [120] Baldus H-P, Jansen M. Moderne Hochleistungskeramikendamorphe anorganische Netzwerke aus molekularen Vorla¨ufern. Angew Chem 1997;109:338e54. [121] Baldus HP, Jansen M, Sporn D. Ceramics fibers for matrix composites in high-temperature engine applications. Science 1999;285:699e703. [122] Riedel R, Kienzle A, Petzow G, Bru¨ck M, Vaahs T. Borhaltige Polysilazane, Verfahren zu ihrer Herstellung, aus ihnen herstellbare Bor und Silicium enthaltende keramische Materialien, sowie deren Herstellung DE-4320783A1; 1994. [123] Kienzle A. Darstellung und Verarbeitung borhaltiger elementorganischer Vorstufen zur Herstellung keramischer Materialien in den Systemen SiCB und SiCBN, PhD Thesis, University of Stuttgart; 1994. [124] Riedel R, Kienzle A, Dressler W, Ruwisch L, Bill J, Aldinger F. A silicoboron carbonitride ceramic stable to 2.000  C. Nature 1996;382:796e8. [125] Riedel R, Bill J, Kienzle A. Boron-modified inorganic polymersdprecursors for the synthesis of multicomponent ceramics. Appl Organomet Chem 1996;10:241e56. [126] Riedel R, Ruwisch LM, An L, Raj R. amorphous silicoboron carbonitride ceramic with very high viscosity at temperatures above 1500  C. J Am Ceram Soc 1998;81:3341e4. [127] Jansen M, Ja¨schke T. Preparation, crystal structure, and spectroscopic characterization of [(H3C)3Si]NH(BCl2). Z Anorg Allg Chem 1999;625:1957e9. [128] Mu¨ller U, Weinmann M, Jansen M. Cl2MeSieNHeBCl2 and ClMe2SieNHeBCl2: novel processable single source precursors of amorphous Si/C/B/N ceramics. J Mater Chem 2008;18: 3671e9. [129] Jones PR, Myers JK. The hydroboration of methylchlorosilylalkenes. J Organomet Chem 1972;34:C9eC11.

1096

[130] Ruwisch LM, Du¨richen P, Riedel R. Synthesis of silyl substituted organoboranes by hydroboration of vinylsilanes. Polyhedron 2000;19:323e30. [131] Ruwisch LM. Synthese und Hochtemperaturverhalten borhaltiger Siliciumcarbonitride, PhD Thesis, Technical University of Darmstadt; 1998 [132] Riedel R, Ruwisch LM. Ceramics, McGraw-Hill yearbook of science & technology 1999. New York: McGraw Hill; 1998. 70e73. [133] Weinmann M, Schuhmacher J, Kummer H, Prinz S, Peng J, Seifert HJ, et al. Synthesis and thermal behavior of novel SieBeCeN ceramic precursors. Chem Mater 2000;12:623e32. [134] Aldinger F, Weinmann M, Bill J. Precursor-derived SieBeCeN ceramics. Pure Appl Chem 1998;70:439e48. [135] Weinmann M. High temperature stable ceramics from inorganic polymers in: precursor-derived ceramics. In: Bill J, Wakai F, Aldinger F, editors. Weinheim: Wiley-VCH; 1999. p. 83e92. [136] Weinmann M, Aldinger F. SieBeCeN precursors: synthesis and thermal stability. In: New properties from atomic level processingdproc. symposium on international joint project ceramics superplasticity, Tokyo; 1999. p. 16e23. [137] Weinmann M, Seifert HJ, Aldinger F. Boron-containing nonoxide ceramics from organometallic polymers: synthesis, thermolysis and the influence of boron on materials thermal stability. In: Davidson MG, Hughes AK, Marder TB, Wade K, editors. Contemporary boron chemistry. Cambridge: The Royal Society of Chemistry; 2000. p. 88e95. [138] Mu¨ller A, Gerstel P, Weinmann M, Bill J, Aldinger F. Correlation of boron content and high temperature stability in SieBeCeN ceramics. J Eur Ceram Soc 2000;20:2655e9. [139] Mu¨ller A, Gerstel P, Weinmann M, Bill J, Aldinger F. Correlation of boron content and high temperature stability in SieBeCeN ceramics II. J Eur Ceram Soc 2001;21:2171e7. [140] Weinmann M, Kamphowe TW, Fischer P, Aldinger F. Tris(hydridosilylethyl)boranes: highly reactive synthons for polymeric silicon compounds. J Organomet Chem 1999;592:115e27. [141] Weinmann M, Bill J, Aldinger F. Synthese borhaltiger Polysilazane durch katalytisch induzierte dehydrierende Kupplungsreaktionen von Hydridosilylethylboranen mit Ammoniak, ¨ berprima¨ren Aminen und polymeren Silazanen sowie deren U fu¨hrung in keramische Materialien, DE-19741459A1; 1999. [142] Weinmann M, Kamphowe TW, Schuhmacher J, Mu¨ller K, Aldinger F. Design of polymeric SieBeCeN ceramic precursors for the application in fiber-reinforced composite materials. Chem Mater 2000;12:2112e22. [143] Weinmann M, Kamphowe TW, Lee S-H, Aldinger F. Hochtemperaturstabile keramische Faserverbundwerkstoffe durch Infiltration und Thermolyse siliciumorganischer Precursoren. In: Wielage B, Leonhardt G, editors. Verbundwerkstoffe und Werkstoffverbunde. Weinheim: Wiley-VCH; 2001. p. 268e71. [144] Weinmann M, Kamphowe TW, Bill J, Aldinger F. Synthese borhaltiger Polysilazane durch Hydrosilylierungsreaktionen von Hydridosilylethylboranen mit vinylierten Polysilazanen sowie ¨ berfu¨hrung in keramische Materialien, DE19741460A1; deren U 1999. [145] See for example Generazio ER. Advanced Ceramic Matrix Composites: Design Approaches, Testing & Life Prediction Methods, Technomic Pub Co; 1995.

Handbook of Advanced Ceramics

[146] Weinmann M, Haug R, Bill J, Aldinger F, Schuhmacher J, Mu¨ller K. Boron-containing polysilylcarbodi-imides: a new class of molecular precursors for SieBeCeN ceramics. J Organomet Chem 1997;541:345e53. [147] Weinmann M, Haug R, Bill J, De Guire M, Aldinger F. Boronmodified polysilylcarbodi-imides as precursors for SieBeCeN ceramics: synthesis, plastic forming and high temperature behavior. Appl Organomet Chem 1998;12:725e34. [148] Schuhmacher J, Mu¨ller K, Weinmann M, Bill J, Aldinger F. Festko¨rper-NMR-spektroskopische Untersuchung der Keramisierung anorganischer Precursorpolymere zu SieBeCeN-Keramiken. In: Heinrich J, Ziegler G, Hermel W, Riedel H, editors. Proc. Werkstoffwoche 1998, B and VII, Keramik/Simulation. Weinheim: Wiley-VCH; 1999. p. 321e5. [149] Lechler R, Hausen H-D, Weidlein J. Reaktion von (CH3)3SieN] C]NeSi(CH3)3 mit Aluminium-, Gallium-und Indium-Trimethyl. J Organomet Chem 1989;359:1e12. [150] Weinmann M, Zern A, Ho¨rz M, Berger F, Mu¨ller K, Aldinger F. Precursor-derived SiC/Si3N4 nano/nano composites: synthesis, structure and phase evolution at high temperatures. J Met Nano Mater 2002;386e388:335e40. [151] Weinmann M, Ho¨rz M, Berger F, Mu¨ller A, Mu¨ller K, Aldinger F. Dehydrocoupling of tris(hydridosilylethyl)boranes and cyanamide: a novel access to boron-containing polysilylcarbodiimides. J Organomet Chem 2002;659:29e42. [152] Annual Report, Fraunhofer Institut fu¨r Silikatforschung, Wu¨rzburg Germany; 2004. [153] Weinmann M, Kroschel M, Ja¨schke T, Nuss J, Jansen M, Kolios G, et al. Towards continuous processes for the synthesis of precursors of amorphous Si/B/N/C ceramics. J Mater Chem 2008;18:1810e8. [154] Weinmann M, Jansen M. Siliconeboronecarbonenitrogen ceramics and precursor compounds, methods for the salt-free polymerization of RnHal3nSieXeBRmHal2m. WO 2009/ 124725 Al 2009. [155] Abe Y, Gunji T. Oligo- and polysiloxanes. Prog Polym Sci 2004;29:149e82. [156] Chojnowski J, Cypryk M, Kurjata J. Organic polysilanes interrupted by heteroatoms. Prog Polym Sci 2003;28:691e728. [157] Chojnowski J, Kurjata J, Rubinsztajn S. Poly(oxymultisilane)s by ring-opening polymerization. Fully methylated silicon analogues of oxirane and THF polymers. Makromol Chem Rapid Commun 1988;9:469e75. [158] Chi K. Carbon-containing monolithic glasses via the solegel process. Ceram Eng Sci Proc 1983;4:704e17. [159] Babonneau F, Thorne K, Mackenzie JD. Dimethyldiethoxysilane /tetraethoxysilane copolymers: precursors for the siliconecarboneoxygen system. Chem Mater 1989;1:554e8. [160] Soraru GD. Silicon oxycarbide glasses from gels. J Sol-Gel Sci Technol 1994;2:843e8. [161] Soraru GD, D’Andrea G, Campostrini R, Babonneau F, Mariotto G. Structural characterization and high-temperature behavior of silicon oxycarbide glasses prepared from solegel precursors containing SieH bonds. J Am Ceram Soc 1995;78:379e87. [162] Gervais C, Babonneau F, Dallabona N, Soraru GD. Solegelderived siliconeboron oxycarbide glasses containing mixed silicon oxycarbide (SiCxO4x) and boron oxycarbide (BCyO3y) units. J Am Ceram Soc 2001;84:2160e4.

Chapter | 11.1.10

Precursor-Derived Ceramics

[163] Ionescu E, Linck C, Fasel C, Mu¨ller M, Kleebe H-J, Riedel R. Polymer-derived SiOC/ZrO2 ceramic nanocomposites with excellent high-temperature stability. J Am Ceram Soc 2010;93: 241e50. [164] Ionescu E, Papendorf B, Kleebe H-J, Poli F, Mu¨ller K, Riedel R. Polymer-derived silicon oxycarbide/hafnia ceramic nanocomposites. Part I: phase and microstructure evolution during the ceramization process. J Am Ceram Soc 2010;93:1774e82. [165] Schwartz KB, Rowcliffe DJ, Blum YD, Laine RM. Thermal conversion of preceramic polysilazanes to Si3N4: characterization of pyrolysis products. Mat Res Soc Symp Proc 1986;73:407e12. [166] Soraru GD, Babonneau F, Mackenzie JD. Structural evolutions from polycarbosilane to SiC ceramic. J Mater Sci 1990;25: 3886e93. [167] Lavedrine A, Bahloul D, Goursat P, Chong Kwet Yive NS, Corriu RJ, Leclercq D, et al. Pyrolysis of polyvinylsilazane precursors to silicon carbonitride. J Eur Ceram Soc 1991;8:221e7. [168] Boury B, Corriu RJ, Douglas WE. Poly(carbosilane) precursors of silicon carbide: the effect of cross-linking on ceramic residue. Chem Mater 1991;3:487e9. [169] Schmidt WR, Interrante LV, Doremus RH, Trout TK, Marchetti PS, Maciel GE. Pyrolysis chemistry of an organometallic precursor to silicon carbide. Chem Mater 1991;3:257e67. [170] Chong Kwet Yive NS, Corriu RJ, Leclercq D, Mutin PH, Vioux A. Silicon carbonitride from polymeric precursors: thermal cross-linking and pyrolysis of oligosilazane model compounds. Chem Mater 1992;4:141e6. [171] Han HN, Lindquist DA, Haggerty JS, Seyferth D. Pyrolysis chemistry of poly(organosilazanes) to silicon ceramics. Chem Mater 1992;4:705e11. [172] Corriu RJ, Leclercq D, Mutin PH, Vioux A. Thermogravimetric mass spectrometric investigation of the thermal conversion of organosilican precursors into ceramics under argon and ammonia. 1. Poly(carbosilane). Chem Mater 1992;4:711e6. [173] Schmidt WR, Marchetti PS, Interrante LV, Hurley Jr WJ, Lewis RH, Doremus RH, et al. Ammonia-induced pyrolytic conversion of a vinylic polysilane to silicon nitride. Chem Mater 1992;4:937e47. [174] Chong Kwet Yive NS, Corriu RJ, Leclercq D, Mutin PH, Vioux A. Thermogravimetric analysis/mass spectrometry investigation of the thermal conversion of organosilicon precursors into ceramics under argon and ammonia. 2. Poly(silazanes). Chem Mater 1992;4:1263e71. [175] Brodie N, Majoral J-P, Disson J-P. An NMR study of the step by step pyrolysis of a polysilazane precursor of silicon nitride. Inorg Chem 1993;32:4646e9. [176] Mocaer D, Pailler R, Naslain R, Richard C, Pillot JP, Dunogue`s J, et al. SieCeN ceramics with a high microstructural stability elaborated from the pyrolysis of new polycarbosilazane precursors. J Mater Sci 1993;28:2615e31. [177] Laine RM, Babonneau F, Blowhowiak KY, Kennish RA, Rahn JA, Exarhos GJ, et al. The evolutionary process during pyrolytic transformation of poly(N-methylsilazane) from a preceramic polymer into an amorphous silicon nitride/carbon composite. J Am Ceram Soc 1995;78:137e45. [178] Babonneau F, Maquet J, Bonhomme C, Richter R, Roewer G, Bahloul D. 29Si and 13C NMR investigation of the polysilaneto-poly(carbosilane) conversion of poly(methylchlorosilanes)

1097

[179]

[180]

[181]

[182]

[183] [184] [185] [186]

[187] [188]

[189]

[190]

[191]

[192]

[193] [194]

[195] [196] [197]

[198] [199]

using cross-polarization and inversion recovery cross-polarization techniques. Chem Mater 1996;8:1415e28. Bahloul D, Pereira M, Ge´rardin C. Pyrolysis chemistry of polysilazane precursors to silicon carbonitride. J Mater Chem 1997;7:109e16. Ge´rardin C, Taulelle F, Bahloul D. Pyrolysis chemistry of polysilazane precursors to silicon carbonitride. J Mater Chem 1997;7:117e27. Liu Q, Wu H-J, Lewis R, Maciel GE, Interrante LV. Investigation of the pyrolytic conversion of poly(silylenemethylene) to silicon carbide. Chem Mater 1999;11:2038e48. Schuhmacher J, Weinmann M, Bill J, Aldinger F, Mu¨ller K. Solid state NMR studies of the preparation of SieCeN ceramics from polysilylcarbodiimide polymers. Chem Mater 1998;10:3913e22. Laine RM, Babonneau F. Preceramic polymer routes to silicon carbide. Chem Mater 1993;5:260e79. Greil P. Near net shape manufacturing of polymer derived ceramics. J Eur Ceram Soc 1998;18:1905e14. Greil P. Active-filler-controlled pyrolysis of preceramic polymers. J Am Ceram Soc 1995;78:835e48. Williams EA. NMR spectroscopy of organosilicon compounds. In: Patai S, Rappoport Z, editors. The chemistry of organic silicon compounds. London: J. Wiley & Sons; 1989. p. 511e55. Yajima S. Special heat-resisting materials from organometallic polymers. Am Ceram Soc Bull 1983;62:893e915. Yajima S, Hayashi J, Omori M, Okamura K. Development of a silicon carbide fibre with high tensile strength. Nature 1976;261:683e5. Yajima S, Hayashi J, Omori M, Okamura K. Synthesis of continuous Sic Fibers with high tensile strength. J Am Ceram Soc 1976;59:324e7. Yajima S, Iwai T, Yamamura T, Okamura K, Hasegawa Y. Synthesis of a polytitanocarbosilane and its conversion into inorganic compounds. J Mater Sci 1981;16:1349e55. Soraru GD, Babonneau F, Mackenzie JD. Structural concepts on new amorphous covalent solids. J Non-Cryst Solids 1988;106: 256e61. Polysilazanes with general structure [RSi(NH)1.5]n, e.g. commercially available poly(vinylsilsesquiazane) VT50, Hoechst AG, Germany. Popper P. New electrical ceramics and inorganic polymers. Brit Ceram Res Assn Spec Publ 1967;57:1e20. Seyferth D, Wiseman GH, Prud’homme C. Siliconenitrogen polymers and ceramics derived from reactions of dichlorosilane, H2SiCl2. Mater Sci Res 1984;17:263e9. Seyferth D, Wiseman GH. Polysilazane routes to silicon nitride. Am Chem Soc Polym, Div Polym Prep 1984;25:10e2. Vaahs T, Bru¨ck M, Bo¨cker WD. Polymer-derived silicon nitride and silicon carbonitride fibers. Adv Mater 1992;4:224e6. Bahloul D, Pereira M, Goursat P, Choong Kwet Yive NS, Corriu RJP. Preparation of silicon carbonitrides from an organosilicon polymer: I, thermal decomposition of the cross-linked polysilazane. J Am Ceram Soc 1993;76:1156e62. PCS is a commercially available product of Nippon Carbon Co. Ltd., Type S, Tokyo, Japan. PHPS is a commercially available product of Tonen Co., N-N110, Tokyo, Japan.

1098

[200] Iwamoto Y, Vo¨lger W, Kroke E, Riedel R, Saitou T, Matsunaga K. Crystallization behavior of amorphous silicon carbonitride ceramics derived from organometallic precursors. J Am Ceram Soc 2001;84:2170e8. [201] Bill J, Kamphowe TW, Mu¨ller A, Wichmann T, Zern A, Jalowiecki A, et al. Precursor-derived Sie(Be)CeN ceramics e thermolysis, amorphous state and crystallization. Appl Organomet Chem 2001;15:777e93. [202] Schuhmacher J. Festko¨rper-NMR-Untersuchungen zur Umwandlung von Polysilazanen und Polysilylcarbodiimiden in Sie(B)e CeN-Keramiken, PhD Thesis, University of Stuttgart; 2000. [203] Weinmann V. Molekulare Vorstufen fu¨r hochtemperaturstabile SieBeCeN-Keramiken, Habilitation Thesis, Universita¨t Stuttgart: Germany; 2003. [204] Schuhmacher J, Berger F, Weinmann M, Bill J, Aldinger F, Mu¨ller K. Solid sate NMR and FT IR studies of the preparation of SieBeCeN ceramics from boron-modified polysilazanes. Appl Organomet Chem 2001;15:809e19. [205] Dunham ML, Bailey DL, Mixer RY. New curing systems for silicon rubber. Ind Eng Chem 1957;49:1373e6. [206] Valles EM, Macosco Ch M. Properties of networks formed by end linking of poly(dimethylsiloxane). Macromolecules 1979;12: 673e9. [207] Heidingsfeldova M, Capka M. Rhodium complexes as catalysts for hydrosilylation crosslinking of silicone rubber. J Appl Polym Sci 1985;30:1837e46. [208] Grzelka A, Chojnowski J, Cypryk M, Fortuniak W, Hupfield PC, Taylor RG. Polysiloxanol condensation and disproportionation in the presence of a superacid. J Organomet Chem 2004;689:705e13. [209] Scheffler M, Bordia R, Travitzky N, Greil P. Development of a rapid crosslinking preceramic polymer system. J Eur Ceram Soc 2005;25:175e80. [210] Corriu RJP, Leclerq D, Mutin PH, Vioux A. Preparation and structure of silicon oxycarbide glasses derived from polysiloxane precursors. J Sol-Gel Sci Technol 1997;8:327e30. [211] Pantano CG, Singh AK, Zhang H. Silicon oxycarbide glasses. J Sol-Gel Sci Technol 1999;14:7e25. [212] Belot V, Corriu RJP, Leclerq D, Mutin PH, Vioux A. Thermal redistribution reactions in crosslinked polysiloxanes. J Polym Sci A: Polym Chem 1992;30:613e23. [213] Chomel AD, Dempsey P, Latournerie J, Hourlier-Bahloul D, Jayasooriya UA. Gel to glass transformation of methyltriethoxysilane: a silicon oxycarbide glass precursor investigated using vibrational spectroscopy. Chem Mater 2005;17:4468e73. [214] Bois L, Maquet J, Babonneau F, Mutin PH, Bahloul D. Structural characterization of solegel derived oxycarbide glasses. 1. Study of the pyrolysis process. Chem Mater 1994;6:796e802. [215] Kalfat R, Babonneau F, Gharbi N, Zarrouk H. 29 Si MAS NMR investigation of the pyrolysis process of cross-linked polysiloxanes prepared from polymethylhydrosiloxane. J Mater Chem 1996;6:1673e8. [216] Gualandris V, Hourlier-Bahloul D, Babonneau F. Structural investigation of the first stages of pyrolysis of SieCeO preceramic polymers containing SieH bonds. J Sol-Gel Sci Technol 1999;14:39e48. [217] Soraru GD, Pederiva L, Latournerie J, Raj R. Crystallization behavior of amorphous silicon carbonitride ceramics derived from organometallic precursors. J Am Ceram Soc 2002;85:2181e7.

Handbook of Advanced Ceramics

[218] Mutin PH. Control of the composition and structure of silicon oxycarbide and oxynitride glasses derived from polysiloxane precursors. J Sol-Gel Sci Technol 1999;14:27e38. [219] Liebau V, Hauser R, Riedel R. Amorphous SiBCO ceramics derived from novel polymeric precursors. C R Chim 2004;7:463e9. [220] Ngoumeni-Yappi R, Fasel C, Riedel R, Ischenko V, Pippel E, Woltersdorf J, et al. Tuning of the rheological properties and thermal behavior of boron-containing polysiloxanes. Chem Mater 2008;20:3601e8. [221] Harshe R, Balan C, Riedel R. Amorphous Si(Al)OC ceramic from polysiloxanes: bulk ceramic processing, crystallization behavior and applications. J Eur Ceram Soc 2004;24:3471e82. [222] Alonso B, Sanchez C. Structural investigation of polydimethylsiloxaneevanadate hybrid materials. J Mater Chem 2000;10:377e86. [223] Fukushima M, Yasuda E, Nakamura Y, Tanabe Y. Pyrolysis behavior of organiceinorganic hybrids with SieOeNb/SieOeTa oxygen bridged heterometallic bonds. J Ceram Soc Jpn 2003;111:857e9. [224] Fukushima M, Yasuda E, Nakamura Y, Teranishi Y, Nakamura K, Tanabe Y. Raman characterization of oxidation behavior of free carbon in silicon oxycarbide ceramics. J Ceram Soc Jpn 2004;112:612e4. [225] Dire´ S, Ceccato R, Babonneau F. Structural and microstructural evolution during pyrolysis of hybrid polydimethylsiloxanee titania nanocomposites. J Sol-Gel Sci Technol 2005;34:53e62. [226] Dire´ S, Ceccato R, Gialanella S, Babonneau F. Thermal evolution and crystallisation of polydimethylsiloxaneezirconia nanocomposites prepared by the solegel method. J Eur Ceram Soc 1999;19:2849e58. [227] Narisawa M, Yasuda H, Mori R, Mabuchi H, Oka K, Kim Y-W. Silicon carbide particle formation from carbon blackd polymethylsilsesquioxane mixtures with melt pressing. J Ceram Soc. Jpn 2008;116:121e5. [228] Ionescu E, Babonneau F, Gervais C. Polymer-to-ceramic transformation. In: Colombo P, Riedel R, Soraru GD, Kleebe H-J, editors. Polymer derived ceramics: from nanostructure to applications. Lancaster, PA, USA: DEStech Publications, Inc.; 2009. p. 108e26. [229] Brus J, Kolar F, Machovic V, Svitilova J. Structure of silicon oxycarbide glasses derived from poly(methylsiloxane) and poly [methyl(phenyl)siloxane] precursors. J Non-Cryst Solids 2001; 289:62e74. [230] Mackenzie KJD, Smith ME. Multinuclear solid-state NMR of inorganic materials. In: Pergamon Materials Series, vol. 6. Oxford: Pergamon-Elsevier; 2002. [231] Lippmaa E, Ma¨gi M, Samoson A, Engelhardt G, Grimmer A-R. Structural studies of silicates by solid-state high-resolution silicon-29 NMR. J Am Chem Soc 1980;102:4889e93. [232] Lippmaa E, Ma¨gi M, Samoson A, Engelhardt G. Investigation of the structure of zeolites by solid-state high-resolution silicon-29 NMR spectroscopy. J Am Chem Soc 1981;103:4992e6. [233] Hartman JS, Millard RL, Vance ER. A29Si magic angle spinning NMR and DTA study of thermal crystallization of sphene and zircon gels. J Mater Sci 1990;25:2785e90. [234] Andrianainarivelo M, Corriu R, Leclercq D, Mutin PH, Vioux A. Mixed oxides SiO2eZrO2 and SiO2eTiO2 by a non-hydrolytic solegel route. J Mater Chem 1996;6:1665e71.

Chapter | 11.1.10

Precursor-Derived Ceramics

[235] Kleebe H-J, Turquat C, Soraru GD. Structural characterization and high-temperature behavior of silicon oxycarbide glasses prepared from solegel precursors containing SieH bonds. J Am Ceram Soc 2001;84:1073e80. [236] Gregori G, Kleebe H-J, Blum YD, Babonneau F. Evolution of C-rich SiOC ceramicsePart II. Characterization by high lateral resolution techniques: electron energy-loss spectroscopy, highresolution TEM and energy-filtered TEM. Int J Mat Res 2006;97:710e20. [237] Evans AG. In: Cahn RW, Evans AG, McLean M, editors. Ceramic and ceramic composites as high-temperature structural materials: challenges and opportunities in high-temperature structural materials. London (: Chapman & Hall; 1996. [238] Gogotsi YG, Lavrenko VA. Corrosion of high-performance ceramics. Berlin, Heidelberg: Springer- Verlag; 1992. [239] Jorgensen PJ, Wadsworth ME, Cutler IB. Oxidation of silicon carbide. J Am Ceram Soc 1959;42:613e6. [240] Franz I, Langheinrich W. Conversion of silicon nitride into silicon dioxide through the influence of oxygen. Solid State Electron 1971;14:499e505. [241] Peuckert M, Vaahs T, Bru¨ck M. Ceramics from organometallic polymers. Adv Mater 1990;2:398e404. [242] Toreki W. Polymeric precursors to ceramicsda review. Polymer News 1991;16:1e14. [243] Riedel R. Advanced ceramics from inorganic polymers. In: Cahn RW, Haasen P, Kramer EJ, editors. Materials science and technology, chapter 11, advanced ceramics from inorganic polymers. Weinheim: VCH; 1996. p. 1e50. [244] Seifert HJ, Aldinger F. Applied phase studies. Z Metallkd 1996;87:841e53. [245] Seifert HJ, Aldinger F. Thermodynamic calculations in the SieBeCeNeO system. In: Bill J, Wakai F, Aldinger F, editors. Precursor-derived ceramics. Weinheim: Wiley-VCH; 1999. p. 165e74. [246] Liu AY, Cohen ML. Prediction of new low compressibility solids. Science 1989;245:841e2. [247] Kaner RB, Kouvetakis J, Warble CE, Sattler ML, Bartlett N. Boronecarbonenitrogen materials of graphite-like structure. Mater Res Bull 1987;22:399e404. [248] Kouvetakis J, Sasaki T, Shen C, Hagiwara R, Lerner M, M Krishnan K, et al. Novel aspects of graphite intercalation by fluorine and fluorides and new B/C, C/N and B/C/N materials based on the graphite network. Synth Metals 1989; 34:1e7. [249] Kawaguchi M, Kawashima T. Synthesis of a new graphite-like layered material of composition BC3N. J Chem Soc Chem Commun 1993;61:1133. [250] Kawaguchi M, Kawashima T, Nakajima T. Synthesis of a new graphite-like material of composition BCxN(x ¼ 3 and 7) as an electrode matrix. Denki Kagaku 1993;61:1403e8. [251] Maya L. Aminoborane polymers as precursors of CeNeB ceramic materials. J Am Ceram Soc 1988;71:1104e7. [252] Maya L. J. Semiconducting amorphous film containing carbon nitrogen and boron. Electrochem Soc 1988;135:1278e81. [253] Riedel R, Bill J, Passing G. A novel carbon material derived from pyridineeborane. Adv Mater 1991;3:551e2. [254] Bill J, Riedel R, Passing G. Amin-Borane als Precursoren fu¨r Borcarbidnitrid. Z Anorg Allg Chem 1992;610:83e90.

1099

[255] Bill J, Frieß M, Riedel R. Conversion of amine-boranes to boron carbide nitride. Eur J Solid State Inorg Chem 1992;29:195e212. [256] Du¨rr J, Schempp S, Lamparter H-P, Bill J, Steeb S, Aldinger F. X-ray and neutron small angle scattering with SieCeN ceramics using isotopic substitution. Solid State Ionics 1997;101e103: 1041e7. [257] Du¨rr J, Lamparter H-P, Bill J, Steeb S, Aldinger F. An X-ray and neutron scattering investigation of precursor derived Si24C43N33 ceramics. J Non-Cryst Solids 1998;232e234:155e61. [258] Schempp S, Du¨rr J, Lamparter H-P, Bill J, Aldinger F. Short range and medium range structure of amorphous SieCeN ceramics by X-ray and neutron diffraction. Z Naturf 1998;53a:127e33. [259] Jalowiecki A, Bill J, Aldinger F, Mayer J. Interface characterization of nanosized B-doped Si3N4/SiC ceramics. Composites Part A 1996;27:717e21. [260] Schuhmacher J, Mu¨ller K, Weinmann M, Bill J, Aldinger F. In: Heinrich J, Ziegler G, Hermel W, Riedel H, editors. Festko¨rperNMRspektroskopische Untersuchung der Keramisierung anorganischer Precursorpolymere zu Sie(B)eCeN-Keramiken, in Werkstoffwoche ‘98, Band VII: Symposium 9: Keramik/ Symposium 14: Simulation Keramik. Wiley-VCH; 1999 (in german). [261] Kasper B. Phasengleichgewichte im System BeCeNeSi, PhD Thesis, University of Stuttgart; 1996. [262] Weiss J, Lukas H-L, Lorenz J, Petzow G, Krieg H. Calculation of heterogeneous phase equilibria in oxideenitride systems: I. The quaternary system C/Si/N/O. Calphad 1981;5:125e40. [263] Seifert HJ, Peng J, Aldinger F. Phase equilibria and thermal analysis of SieCeN ceramics. J Alloys Compd 2001;320: 251e61. [264] Peng J, Seifert HJ, Aldinger F. Thermal analysis of SieCeN ceramics derived from polysilazanes, in “ceramicsdprocessing, reliability, tribology and wear”. In: Mu¨ller G, editor. Proc. EUROMAT ’99, vol. 12. Weinheim: Wiley-VCH; 1999. p. 120e6. [265] Seifert HJ, Lukas H-L, Aldinger F. Development of SieBeCeN ceramics supported by constitution and thermochemistry. Ber Bunsenges Phys Chem 1998;102:1309e13. [266] Varga T, Navrotsky A, Moats JL, Morcos RM, Poli F, Mu¨ller K, et al. Aminoborane polymers as precursors of CeNeB ceramic materials. J Am Ceram Soc 2007;90:3213e9. [267] Morcos RM, Navrotsky A, Varga T, Blum Y, Ahn D, Poli F, et al. Energetics of SixOyCz polymer-derived ceramics prepared under varying conditions. J Am Ceram Soc 2008;91:2969e74. [268] Morcos RM, Mera G, Navrotsky A, Varga T, Riedel R, Poli F, et al. Enthalpy of formation of carbon-rich polymer-derived amorphous SiCN ceramics. J Am Ceram Soc 2008;91:3349e54. [269] Peng J. Thermochemistry and constitution of precursor derived Sie(Be)CeN ceramics. PhD Thesis, University of Stuttgart; 2002. [270] Saha A, Raj R. Crystallization maps for SiCO amorphous ceramics. J Am Ceram Soc 2007;90:578e83. [271] Schiavon M, Gervais C, Babonneau F, Soraru GD. Crystallization behavior of novel silicon boron oxycarbide glasses. J Am Ceram Soc 2004;87:203e8. [272] Klonczynski A, Schneider G, Riedel R, Theissmann R. Influence of boron on the microstructure of polymer derived SiCO ceramics. Adv Eng Mater 2004;6:64e8.

1100

[273] Ionescu E, Papendorf B, Kleebe H-J, Riedel R. Polymer-derived silicon oxycarbide/hafnia ceramic nanocomposites. Part II: stability toward decomposition and microstructure evolution at T >> 1000  C. J Am Ceram Soc 2010;93:1783e9. [274] Schubert U, Hu¨sing N. Synthesis of inorganic materials. Weinheim: Wiley-VCH, ISBN 3-527-29550-X; 2000. [275] Bill J, Wakai F, Aldinger F. In: Precursor-derived ceramics: synthesis, structure and high-temperature mechanical properties. Weinheim: Wiley-VCH, ISBN 3-527-29814-2; 1999. [276] Bill J, Aldinger F. Progress in material synthesis. Z Metallkd 1996;87:827e40. [277] Greil P, Seibold M. Modelling of dimensional changes during polymer-ceramic conversion for bulk component fabrication. J Mater Sci 1992;27:1053e60. [278] Riedel R, Toma L, Fasel C, Miehe G. Polymer-derived mulliteeSiC-based nanocomposites. J Eur Ceram Soc 2009;29: 3079e90. [279] Riedel R, Pasing G, Scho¨nfelder H, Brook RJ. Synthesis of dense silicon-based ceramics at low temperatures. Nature 1992;355: 714e7. [280] Riedel R, Seher M, Mayer J, Szabo´ D-V. Polymer-derived Sibased bulk ceramics, part I: preparation, processing and properties. J Eur Ceram Soc 1995;15:703e15. [281] Seitz J, Bill J. Production of compact polysilazane-derived Si/C/ N-ceramics by plastic forming. J Mat Sci Lett 1996;15:391e3. [282] Haug R, Weinmann M, Bill J, Aldinger F. Plastic forming of preceramic polymers. J Eur Ceram Soc 1999;19:1e6. [283] Christ M, Thurn G, Weinmann M, Bill J, Aldinger F. High temperature mechanical properties of SieBeCeN precursorderived amorphous ceramics and the applicability of deformation models developed for metallic glasses. J Am Ceram Soc 2000;83:3025e32. [284] Christ M. Mechanische Hochtemperatureigenschaften von SieBeCeN-Precursorkeramiken. PhD Thesis, University of Stuttgart; 2001. [285] Zimmermann A, Bauer A, Christ M, Cai Y, Aldinger F. Hightemperature deformation of amorphous SieCeN and SieBeCeN ceramics derived from polymers. Acta Mater 2002;50:1187e96. [286] Primary creep 3 ¼ Asn$exp(Q/RT)tc where s is the applied stress, n the stress exponent, Q the activation energy, R the gas constant, T the absolute temperature, t the time, and c the time exponent. A is a constant. [287] Thurn G, Canel J, Bill J, Aldinger F. Compression creep behaviour of precursor-derived Si/C/N ceramics. J Eur Ceram Soc 1999;19:2317e23. [288] Thurn G, Aldinger F. Compression creep behaviour of precursor-derived ceramics. In: Bill J, Wakai F, Aldinger F, editors. Precursor-derived ceramics. Weinheim: Wiley-VCH; 1999. p. 237e45. [289] Scheffler FA, Torrey JD. Coatings. In: Colombo P, Riedel R, Soraru GD, Kleebe H-J, editors. Polymer derived ceramics. From nano-structure to applications. Lancaster, PA, USA: DEStech Publications; 2009. p. 358e68. [290] Marletta G, Toth A, Bertoti I, Minh Duc T, Sommer F, Ferencz K. Optical properties of ceramic-like layers obtained by low energy ion beam irradiation of polysiloxane films. Nucl Instrum Meth B 1998;141:684e92.

Handbook of Advanced Ceramics

[291] Nghiem QD, Jeon J-K, Hong L-Y, Kim D-P. Polymer derived SieCeBeN ceramics via hydroboration from borazine derivatives and trivinylcyclotrisilazane. J Organomet Chem 2003;688:27e35. [292] Torrey JD, Bordia RK. Processing of polymer-derived ceramic composite coatings on steel. J Am Ceram Soc 2008;91:41e5. [293] Torrey JD, Bordia RK. Mechanical properties of polymer-derived ceramic composite coatings on steel. J Eur Ceram Soc 2008;28: 253e7. [294] Goerke O, Feike E, Heine T, Trampert A, Schubert H. Ceramic coatings processed by spraying of siloxane precursors (polymerspraying). J Eur Ceram Soc 2004;24:2141e7. [295] Lei H, Francis LF, Gerberich WW, Scriven LE. Modeling stress and failure in shrinking coatings. Mat Res Soc Symp Proc 2001;653. Z.10.5.1. [296] Kozuka H, Takenaka S, Tokita H, Hirano T, Higashi Y, Hamatani T. Stress and cracks in gel-derived ceramic coatings and thick film formation. J Sol-Gel Sci Technol 2003;26:681e6. [297] Krauß H-J, Motz G. Laser pyrolysis of polysilazaneea new technique for the generation of ceramic-like coatings and structures. Key Eng Mater 2002;206e213:467e70. [298] Konetschny C. Hartstoffhaltige SieCeN-Keramiken aus Polysilazanen, PhD Thesis, Technical University of Darmstadt; 1999. [299] Pavese M, Fino P, Badini C, Ortona A, Marino G. HfB2/SiC as a protective coating for 2D Cf/SiC composites: effect of high temperature oxidation on mechanical properties. Surf Coat Technol 2008;202:2059e67. [300] Cromme P, Scheffler M, Greil P. Ceramic tapes from preceramic polymers. Adv Eng Mater 2002;4:873e7. [301] Kho J-G, Moon K-T, Nouet G, Ruterana P, Kim D-P. Boron-rich boron nitride (BN) films prepared by a single spin-coating process of a polymeric precursor. Thin Solid Films 2001;389:78e83. [302] Jou S, Sun C-T, Chen X. Silicon carbide films from polycarbosilane and their usage as buffer layers for diamond deposition. Diamond Relat Mater 2005;14:1688e94. [303] Termoss H, Toury B, Brioude A, Dazord J, Le Brusq J, Miele P. High purity boron nitride thin films prepared by the PDCs route. Surf Coat Technol 2007;201:7822e8. [304] Blum YD, Platz RM, Crawford E:J. Glass strengthening by polymer-derived ceramic coatings. J Am Ceram Soc 1990;73: 170e2. [305] Mucalo MR, Milestone NB, Vickridge IC, Swain MV. Preparation of ceramic coatings from pre-ceramic precursors. J Mater Sci 1994;29:4487e99. [306] Bakumov V, Gueinzius K, Hermann C, Schwarz M, Kroke E. Polysilazane-derived antibacterial silvereceramic nanocomposites. J Eur Ceram Soc 2007;27:3287e92. [307] Krenkel W, Naslain R, Schneider H. High temperature ceramic matrix composites. Weinheim: Wiley-VCH, ISBN 3-527-303200; 2001. and references cited therein. [308] Bill J, Heimann D. Polymer-derived ceramic coatings on C/CeSiC composites. J Eur Ceram Soc 1996;16:1115e20. [309] Heimann D. Oxidationsschutzschichten fu¨r kohlefaserversta¨rkte Verbundwerkstoffe durch Polymer-Pyrolyse, PhD Thesis, University of Stuttgart; 1997. [310] Jarosch K, Lasa HI. Novel riser simulator for methane reforming using high temperature membranes. Chem Eng Sci 1999;54: 1455e60.

Chapter | 11.1.10

Precursor-Derived Ceramics

[311] Kikuchi E, Nemoto Y, Kajikawa M, Uemiya S, Kojima T. Steam reforming of methane in membrane reactors: comparison of electroless-plating and CVD membranes and catalyst packing modes. Catal Today 2000;56:75e81. [312] Prabhu AK, Oyama ST. Highly hydrogen selective ceramic membranes: application to the transformation of greenhouse gases. J Memb Sci 2000;176:233e48. [313] Kurungot S, Yamaguchi T, Nakao S. Rh/g-Al2O3 catalytic layer integrated with solegel synthesized microporous silica membrane for compact membrane reactor applications. Cat Lett 2003;86:273e8. [314] Iwamoto Y. Precursors-derived ceramic membranes for high-temperature separation of hydrogen. J Ceram Soc Jpn 2007;115: 947e54. [315] Iwamoto Y. Membranes. In: Colombo P, Riedel R, Soraru GD, Kleebe H-J, editors. Polymer derived ceramics. From nanostructure to applications. Lancaster, PA, USA: DEStech Publications; 2009. p. 396. [316] Yoshino Y, Suzuki T, Nair BN, Taguchi H, Itoh N. Development of tubular substrates, silica based membranes and membrane modules for hydrogen separation at high temperature. J Memb Sci 2005;267:8e17. [317] Nagano T, Sato K, Saito T, Iwamoto Y. Gas permeation properties of amorphous SiC membranes synthesized from polycarbosilane without oxygen-curing process. J Ceram Soc Jpn 2006;114: 533e8. [318] Nagano T, Uno N, Saito T, Yamazaki S, Iwamoto Y. Gas permeance behavior at elevated temperature in mesoporous anodic oxidized alumina synthesized by pulse-sequential voltage method. Chem Eng Commun 2007;194:158e69. [319] Iwamoto Y, Sato K, Kato T, Inada T, Kubo Y. A hydrogenpermselective amorphous silica membrane derived from polysilazane. J Eur Ceram Soc 2005;25:257e64. [320] Shelekhin AB, Grosgogeat EJ, Hwang TS. Gas separation properties of a new polymer/inorganic composite membrane. J Memb Sci 1991;66:129e41. [321] Wach RA, Sugimoto M, Yoshikawa M. Formation of silicone carbide membrane by radiation curing of polycarbosilane and polyvinylsilane and its gas separation up to 250  C. J Am Ceram Soc 2007;90:275e8. [322] Suda H, Yamauchi H, Uchimaru Y, Fujiwara I, Haraya K. Structural evolution during conversion of polycarbosilane precursor into silicon carbide-based microporous membranes. J Ceram Soc Jpn 2006;114:539e44. [323] Kusakabe K, Li ZY, Maeda H, Morooka S. Preparation of supported composite membrane by pyrolysis of polycarbosilane for gas separation at high temperature. J Memb Sci 1995;103: 175e80. [324] Li Z, Kusakabe K, Morooka S. Preparation of thermostable amorphous SieCeO membrane and its application to gas separation at elevated temperature. J Memb Sci 1996;118:159e68. [325] Lee D, Zhang L, Oyama ST, Niu S, Saraf RF. Synthesis, characterization, and gas permeation properties of a hydrogen permeable silica membrane supported on porous alumina. J Memb Sci 2004;231:117e26.

1101

[326] Vo¨lger KW, Hauser R, Kroke E, Riedel R, Ikuhara YH, Iwamoto Y. Synthesis and characterization of novel non-oxide solegel derived mesoporous amorphous SieCeN membranes. J Ceram Soc Jpn 2006;114:567e70. [327] Hauser R, Nahar-Borchard S, Riedel R, Ikuhara YH, Iwamoto Y. Polymer-derived sibcn ceramic and their potential application for high temperature membranes. J Ceram Soc Jpn 2006;114:524e8. [328] Wallenberger FT. Advanced inorganic fibers: processes, structures, properties, applications. Dordrecht, Netherlands: Kluwer Academic Publishers, ISBN 0-412-60790-5; 2000. and references cited therein. [329] NicalonÔ and Hi-NicalonÔ are trademarks of Nippon Carbon Co., Tokyo, Japan. [330] Livage J, Sanchez C, Babonneau F. Molecular precursor routes to inorganic solids. In: Interrante LV, Hampden-Smith MJ, editors. Chemistry of advanced materials. New York: Wiley-VCH, ISBN 0-471-18590-6; 1998. p. 398e448. [331] Verbeck W. Formko¨rper aus homogenen Mischungen von Siliciumcarbid und Siliciumnitrid und Verfahren zu ihrer Herstellung, DE-2218960; 1973. [332] Nakaido Y, Otani Y, Kozakai N, Otani S. Silicon nitride carbide fiber from spinnable polymethylsilazanes. Chem Lett 1987:705e6. [333] LeGrow GE, Lim TF, Lipowitz J, Reaoch RS. Ceramics from hydridopolysilazane. Am Ceram Soc Bull 1987;66:363e7. [334] Jones RE, Rabe JA, Peterson AL. New ceramic fibre maintains physical properties at high temperatures. Canad Chem News 1992:9e14. [335] Delverdier O, Monthioux M, Mocaer D, Pailler R. Thermal behavior of polymer-derived ceramics. IV. Si/C/N/O fibers from an oxygen-cured polycarbosilazane. J Eur Ceram Soc 1994;14: 313e25. [336] Lu L, Feng CX, Song YC. Curing polysilazane fibres by exposure to boron trichloride. J Mater Sci Lett 1998;17:481e4. [337] Motz G, Hacker J, Ziegler G. Special modified silazanes for coatings, fibers and CMC’s. Ceram Eng Sci Proc 2000;21:307e14. [338] Bernard S, Weinmann M, Gerstel P, Miele P, Aldinger F. Boronmodified polysilazane as a novel single-source precursor for SiBCN ceramic fibers: synthesis, melt-spinning, curing and ceramic conversion. J Mater Chem 2005;15:289e99. [339] Bernard S, Weinmann M, Cornu D, Miele P, Aldinger F. Preparation of high-temperature stable SieBeCeN fibers from tailored single source polyborosilazanes. J Eur Ceram Soc 2005;25:251e6. [340] Bernard S, Duperrier S, Cornu D, Miele P, Weinmann M, Balan C, et al. Chemical tailoring of single-source molecular and polymeric precursors for the preparation of ceramic fibers. J Optoelectronic Adv Mater 2006;8:648e53. [341] Bernard S, Weinmann M, Cornu D, Miele P. A new generation of boron-based ceramic fibers: design, processing and properties of silicoboron carbonitride (SiBCN) fibers from boron-modified polyvinylsilazanes. Adv Sci Tech 2006;50:9e16. [342] Lee S-H, Weinmann M, Aldinger F. Processing and properties of C/SieBeCeN fiber-reinforced ceramic matrix composites prepared by precursor impregnation and pyrolysis. Acta Mater 2008;56:1529e38.