Heinrich J.G., Aldinger F. (Eds.) Ceramic Materials and Components for Engines

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AlN COMPOUNDING

A1N ceramics are most commonly manufactured via

axial and cold isostatic pressing. The expected

advantage of injection molding

based on the high

efficiency, process stability and materials economy

resulting in low specific cost for mass production of an

advanced ceramic component with complex net shape

and performing surface quality obtained by

this

CIM

processing

[4].

For injection molding of ceramic components, the

knowledge of the rheological behaviour

[5]

of the

compounds is one of the most important requirements

to obtain components without defects. The flow

behaviour depends on the grain size distribution, grain

fraction mix, grain morphology and the binder

formulation. Low flow resistance due to optimized

viscosity leads to a good form filling behaviour,

warping and allows netshape tolerances. It leads

furthermore to reliable mechanical properties, low

mold abrasion and supports the optimization of

machine parameters (p,

v). Depending on the

rheological behaviour of the compound, a highly

performing and stable processing can be achieved.

A special ceramic feedstock together with a well

defined and optimized rheological behaviour has been

developed [3]. The optimized feedstock (78

wt.%

AlN,

22%

BASF binder and additives) recipe is made by co-

milling or mixing sinter additive oxides (CaO,

and Ce203) for superior corrosion resistance properties

together with commercial AlN powder with

monomodal grain size distribution. After mixing with

organic binders and plastifiers in a heatable double-

sigma-blade header, the compounding is performed

by continous extrusion and rotation cutting to pellets.

RHEOLOGY AND INJECTION

MOLDING

The CIM process (Fig.1) is related to the polymer and

metal injection molding but has some special features.

Due to the high powder filler content of the compound

(about 80

wt.%),

a reliable compoundation

homogenisation is very important to achieve defect fiee

green body components with high density and

performing mechanical properties. The abrasive

properties of the ceramic powder require a wear

resistant design of the injection molding machine,

especially the plastification unit and the injection

nozzle. Highly wear resistant and polished steel is also

required for the mold. Before sintering, the temporary

organic binder has to be removed with a well

controlled pyrolysis or chemical degradation.

Otherwise the component will be destroyed during the

sintering process. Especially for the CIM and MIM

process the BASF catalytical debinding method is

applied.

remove the binder, mineral or organic acid

cracks the polymer binder into gaseous monomers,

which are combusted in an afterburning process.

This

debinding method

very fast and gentle to the green

body. The sintering process works at very high

temperatures up to 1900°C and in the case of A1N

under inert gas atmosphere, Ar or N

with optional

addition of NH3.

CanFOt

cracks

pores

I----

Figure

CIM

process flow chart

DEBINDED AND SINTERED AlN

COMPONENT

The AlN green bodies were catalyticaly debinded with

the BASF

IFKB

method

hours, 125"C,

NZ-

atmosphere and organic acid

(

vol.% oxalic acid,

vol.% acetic acid,

vol.% formic acid) for a smooth

debinding of the non-oxide AlN powder) and sintered

in a graphite kiln at 1850°C with three hours dwell in

atmosphere (1 bar). The oxide dope additives of the

compound lead to a superior densification due to a

reduction of the oxygen concentration on the AlN

powder surface and accelerates the densification during

the sinter process

[6].

The liquid phase coating also

improves the corrosion resistance due to the superior

chemical resistivity than the undoped AlN. The liquid

sintering process is shown in Fig.

AlN is soluted in

the solution range A by the the liquid phase and

transported by diffusion to the range of enrichment B

and deposited.

This

process leads to the densification

of the material

[7].

Due to the complex sintering

behaviour of AlN an improved sintering process was

developed. The sintering curve is shown in Fig.3.

The SEM-micrographs of the sintered body with Ce02

sintering additive (Fig.

and

show a typical ceramic

AlN sinter texture with some defects which are caused

by debinding problems. The specimen with this

572

additive develops the best texture. Fig.

shows the

typically liquid sintering phase, which coates the AlN

powder and leads to a densification of the material. The

pore volume is

The samples with CaO and

YO3

additives show inferior results.

First long time corrosion tests

(30

days in concentrated

burner condensate) showed an explicite improvement

in corrosion resistance

Further tests under working

conditions will be done in the near future.

Figure

SEM micrograph of a typical sintered CIM

specimen with debinding caused defects

AlN specimen from CIM processing

OPTICAL CHARACTERIZATION

WITH POLARIZED LIGHT

MICROSCOPY

Figure

Liquid phase sintering mechanism

of A1N (liquid phase: black; AlN grain:

grey); A: range of solution;

range of

enrichment

Due to its crystal structure, AlN shows optical

birefrigeance. The polarized light microscopy is a

method to make birefrigeance visible and measureable.

For

the examination of thin sections

the samples, the

thickness is reduced to

[8,

91.

The observed interferece colour depends

the

chemistry and concentration of the sintering additive.

The higher the birefrigeance, the more oxygen

diffusion into the AlN structure takes place. The best

thermal conductivity has the pure AlN. With increasing

oxygen impurities the thermal conductivity decreases

but the corrosion resistance increases.

r!!!!!!!!!

The result of our examinations is that the value

*.Oh

,.Oh

2.0h

,,Oh

.Lh

%Oh

'Lh

ILh

*Oh ,.Oh

ro..n

birefrigeance gives information about the quality

thermal conductive components (Fig.

and

8).

The

Figure 3:

Sinter program for A1N components higher the birefrigeance, the lower the thermal

conductivity.

mgM1

Figure

SEM micrograph of a sintered AlN

Figure

AlN sample with 1 wt.% CaO

Interference colours: orange

order.

573

Figure

A1N sample with

wt.% CeQ

Interference colours: Yellow

order

Figure

AlN sample with

wt.% additives.

Interference colours: Blue

2"d

order

The CeG additive shows a moderate ogygen

concentration in the sinter phase with a interference

colour of yellow first order. CaO with lower

concentration shows higher interference colours,

YO3

shows blue second order and the lowest thermal

conductivity

COMPONENT DESIGN

The shape of the sintered component is checked by

using a coordinate measurement machine allowing to

verify the dimensional stability after debinding and

sintering procedure. Due to the advantageous

thermophysical properties of AlN, the heat exchanger

elements can be designed more compact

smaller

than conventional models (Fig.9 and

11).

A segmented

ribbed design (Fig.10 and

12)

enables to achieve an

efficient convective flow of the hot burner exhaust gas

around the exchanger. The design of a pressed thyristor

cooler (Fig.

11)

is less effective for the use in domestic

heat exchangers. Because of the completely different

temperature gradients, the flow rate on the water side is

inferior to the ribbed design. The production cost for

this design is significantly higher because the required

amount of ceramic powder material is increased at

about

50%.

Accurate process monitoring via simultaneious process

control (SPC) of the flow behaviour during injection

and form filling results in

increase of the quality and

netshape size

the sintered ceramic component with

low texture induced damages and allows to decrease

the production cost together with higher lifetime of the

mold. This results allow in a first step a small scale

manufacturing of heat exchanger components.

Figure

Conventional steel/copper domestic heat

exchanger (courtesy of Bosch-Junkers, Germany)

Figure

10:

Form

filling study for the new designed ribbed

AlN heat exchanger component, demonstrated with a fan

rotor as reference shape

Figure

11:

AlN thyristor heat exchanger, acc.

ANCeram, Bayreuth, Germany

[2]

$rib

pas

side

Figure

12:

Schematic illustration of a ribbed A1N

heat exchanger component

CONCLUSION

Due to its advantageous thermophysical properties,

AlN has been recently established. The material is

known in a wide range of applications in electronics,

computer applications and optical components. A new

application is a high performance hot gas

water heat

exchanger with superior performance for domestic

thermal value heating systems up to 15 kW capacity,

which can substitute conventional steel/copper heat

exchangers. AlN shows superior thermal conductivity

and an improved long term corrosion behaviour. A first

compact design for a ribbed heat exchanger was

developed. The CIM manufacturing of AlN heat

exchangers was shown to be competitive for a

production in near net shape, high surface quality and

for

large lot sizes. The CIM process must be optimized

for ceramic requirements in regard to molding

parameters and mold design.

The quality of CIM formed ceramic components

depends on the rheological properties of the feedstock

and the manufacturing parameters like flow resistance,

form filling behaviour and deforming

ejection.

Netshape tolerances, reproducable high surface quality

and minimized mold abrasion are the result of

optimized processing parameters. The chemistry of the

sintering additives is very important for the

components quality. CeQ seems to be the best additive

for optimized AlN components. Due to the minmized

oxygen exchange with the A1N structure, the expected

thermal conductivity and the corrosion resistance is

superior. Polarized light microsopy confirms this fact.

REFERENCES

[

11 Werdecker, W., Aldinger, F., Aluminum

nitride- an alternative ceramic substrate for

high power applications in microcircuits,

IEEE

Trans. CornDon.. Hvbrids and Manuf.

Technoloey,

[4] 1984,

pp. 399-404

[2] ANCeram, Product Information, ANCeram

GmbH

Co., Germany, 1996

[3] Fischer,

R.,

Gadow,

R.,

Schafer, G.,

Fabrication of AlN Heat exchangers by

ceramic injection molding, Ceramiq

engineering and science Droceedines

[4]

Gadow,

R.,

SuBmuth, G., SpritzgieBen in der

technischen Keramik, Technische Keramik,

1988,

Fritz, H.-G.,

Einfuhrung

die Rheologie und

Rheometrie der Kunststofle,

Technische

Akademie Esslingen, Esslingen, 1996.

[6] German

M. (1996): Sintering Theory And

Practice.-

Wiley

Sons, inc.

[7] Hundere, A. M. (1995): Sintering Of

Aluminium Nitride (AlN) Ceramics.- Institutt

for uorganisk Kjemi Norges Tekniske

Hargskole, Universitetet i Trondheim.

Fischer,

R.,

Franz, E.-D. und Telle,

(1992):

Gefigeuntersuchungen an SpritzguRke-

ramiken mit der Methode der Durchlicht-

Polarisations-Mikroskopie.

Sprechsaal

fir

Betriebsmanagement

Technologie

International Magazine Traditio-nal

and Advanced Ceramics ,Glass and Materials

Technologie, 125. (1992)

;

797-804.

[9] Wahlstrom,

(1979): Optical

Crystallography.- John Wily and Sons, New

York.

(1999), pp.595-602

[4]

pp.206-2

10,

[5]

[8]

575

This Page Intentionally Left Blank

NEW OPPORTUNITY FOR BIMODAL MICROSTRUCTURE CONTROL

IN SILICON NITRIDE

Hai-Doo Kim, Byung-Dong Han, Dong-Soo Park, Dong-Ho Han*, and 11-So0 Kim*

Ceramic Materials Group, Korea Institute

Machinery and Materials,

66 Sangnam-Dong, Changwon, Kyungnam, 641-010, Korea

Dept.

Advanced Materials

Eng.,

Dong-eui Univ.,

San 24, Kaya-Dong, Pusanjin-Ku, Pusan, 614-714, Korea

ABSTRACT

Two step sintering process comprising of carbothermal

reduction process of the compacts of Si3N4, Y2O3, A1203,

and C mixture and the following gas pressure sintering

process was investigated in order to produce the bimodal

microstructure in silicon nitride ceramics.

Carbothermal reduction treatment of powder compacts

of silicon nitride, sintering additives and carbon

mixture

takes the oxygen out of Y-Si-Al-0-N glass, which leads to

the precipitation of crystalline oxynitride phase such

apatite due to the increase in

N/O

ratio in the glass.

Precipitation of apatite phase from Y-Si-A1-0-N glass

leads to the decrease in glass content.

Precipitation of apatite phase and disappearance due to its

melting during sintering process result in the retardation of

densification process up to 1750

“c

followed by the rapid

increase in density due to the recovery

liquid content,

which provides the situation to maximize the difference

between the growing grains and the matrix

grains,

which

leads to the distinct bimodal microstructure.

Since grain growth and densification processes are

thermally activated processes, holding the densification

process to a high temperature followed by the reactivation

of sintering process promotes the development of bimodal

microstructure.

analogous ‘transient solid phase sintering’ induced

by the carbothermal reduction treatment of powder

compacts of silicon nitride, sintering additives, and carbon

mixture may provide a new opportunity to control the

microstructure into bimodal way.

INTRODUCTION

One of the methods to improve the strength and

hture

toughness of silicon nitride is to control the microstructure

into bimodal way, in which the large elongated

grains

in a

fine matrix act

reinforcing grains, mostly, by the crack

bridging mechanism(1-2). In order to have an extreme

bimodal microstructure,

fine

matrix

grains

are prerequisite,

which can be obtained from fine starting a-Si3N4

particles(3-5). Intensive milling is an effective method for

reducing the particle size. However, it might give

rise

to an

increased oxygen content due to the increase in surface

silica content, which affects both phase development and

microstructures(6). It is, therefore, necessary to reduce the

oxygen content while keeping the milled a-Si3N4 particles

fine

possible. Possible ways to reduce the oxygen

content after an intensive milling are chemical treatment

such

HF rinse(7) or carbothermal reduction treatment(&

14).

Most of the works on carbothermal reduction process are

concerned with the reduction of SiOz layer on Si3N4

particles(8-14). However, when the carbon is added to

Si3N4 powder mixture with sintering additives, the result

would be different from the above case, in that the carbon

may take oxygen not only from Si02 but also fiom oxide

glass or oxynitride glass generated from Si3N4(+Si02) and

the sintering additives because the carbothermal reduction

temperature employed in most cases(8-14) is 1450°C or

higher, which is well above the eutectic

temperature( 1370

)

in the system Y203-A1203-Si02( 15)

and, at 1450C,

to p-Si3N4 phase transformation is

possible(6, 16).

When

the carbothermal reduction treatment of the

powder compacts consisting of silicon nitride, sintering

additives and carbon takes the oxygen out of the glass

formed,

this

changes the N/O ratio in oxynitride glassy

phase during sintering compared to the general case

without carbothermal reduction treatment. Since the

oxynitride

glasses

with different

N/O

ratio show different

properties such

viscosity, glass transition temperature,

etc.(l7) and might result in a different precipitation

behaviour due to the limited solubility of nitrogen in

oxynitride glass(l8),

this

would leads to a different

sintering behaviour and

thus, different microstructural

development.

this

work, the carbothermal reduction process of the

compacts of Si3N4, Y2O3, Al2O3, and C mixtures and the

following gas pressure sintering process were investigated.

Special attention was paid to the second phases precipitated

after carbothermal reduction process, the phase

development with increasing temperature and the

microstructural development.

EXPERIMENTAL PROCEDURE

batches of

wt?!

a-Si3N4

(E-10,

Ube

Industries Ltd.,

Tokyo,

Japan),

wt%

Y203 (Fine, H. C. Starck Co., Berlin,

Germany), and 1

wt%

A1203

(AKP-30,

Sumitomo

Chemical Co.,

Osaka,

Japan) with and without 0.6

wt%

carbon(Activated carbon, Aldrich Chemicals Co.,

577

Milwaukee,

WI)

were prepared by planetary ball milling

for 20 h using Si3N4 balls of

in diameter and

methanol in a Si3N4 jar. Hereafter, these batches

are

named

C-0 and C-0.6. Slurries were dried on a hot plate while

stirring, sieved to -100+150 mesh, pressed into compacts of

in diameter by CIP at 250MPa

The compacts were carbothermally processed in a tube

furnace at 1450°C for 10 h in a flowing N2 atmosphere(0.2

literlmin). In order to examine the phase development, the

compacts after carbothermal reduction treatment were

sintered at 1550 to 1850°C for

0.5

h under 1 MPa N2

pressure by 100°C interval. Compacts were

also

sintered at

1850°C for 3 h and 6 h under 2 MPaN2 pressure to see the

densification and microstructural development.

After the planetary ball milling, particle size distribution

and specific surface area were measured by laser

diffraction method (LS-130, Coulter Co., Hialeah, FL) and

BET method (ASAP-201

Micromeritics, Norcross, GA),

respectively. Oxygen and carbon contents before and after

carbothermal reduction treatment were measured using

Oxygen Analyzer (TC-136, LECO Co., St. Joseph, MI) and

Carbon Analyzer (CS-344, LECO Co.). Density was

measured by Archimedes method.

Phase

development was

investigated using X-ray dihtometer (Rigaku Co.,

Tokyo, Japan) on the cross section of the specimen. CuKa

radiation, 35kV-25mA condition and O.O67O/s scan speed

were employed for

this

investigation. The

ratio was

calculated using Gazzara's equation.( 19). The cross section

of the specimen was polished to

with diamond paste

and plasma-etched (96

CF4

02) for

microstructural analysis using SEM (JSM-5800, Jeol,

Tokyo, Japan).

RESULTS

AND DISCUSSION

calculation considering that oxygen content in a-Si3N4

powder

1.16

and that in the sintering additives is

1.74

Oxygen content of the same batch after 4 h

planetiuy ball mill

3.5

wt%

according to the

measurement using the oxygen analyzer. The oxygen

content of 2.9

3.5

seems,

therefore, to be a reference

value because the oxygen contents of the batches which

were prepared without any special treatments such

prolonged milling or an oxidation may fall into

this

range.

Oxygen content of C-0 batch after 20 h planetary ball

mill

4.0

wt%

and that of C-0.6 batch is 4.2

wt%,

while

those

after

carbothermal reduction treatment

are

4.0

wt%

and 3.6

wt%,

respectively. The remaining carbon content

after carbothermal reduction is in 0.1-0.14

wt%

which is

comparable to the carbon content of a-Si3N4 starting

powder(O.1

wt%).

Fig. 1 shows the

XRD

patterns of C-0.6 with increasing

temperature. Just after the carbothermal reduction

treatment, a-Si3N4

major peaks and p-Si3N4 and apatite

minor phases. Appearance of apatite phase suggests that

oxygen-deficient glass resulting from the carbothermal

reduction treatment leads to the Y-Si-Al-0-N glassy phase

with high N/O ratio. Increase in N10 ratio of the oxynitride

glass

gives rise to the precipitation of apatite phase. With

increasing temperature,

phase transformation occurs

with apatite second phase until 1650°C. At 1750"C,

phase transformation is complete and apatite phase

disappeared. The melting point of apatite phase is known

1750"C(20) or 1700-1800"C(21), which suggests that

the disappearance of apatite phase is due to melting. The

precipitation of apatite phase is expected from the pseudo-

quaternary phase equilibrium diagram for the system Si-Y-

0-N(22), in that the composition point shifted downwards

due to the oxygen loss while the yttrium content is nearly

constant.

Density increases with increasing temperature and time

for the specimens of C-0 and C-Odis shown in Fig. 2. C-0

specimens without carboth-al reduction

treatment

shows

the typical densification

curve

which the rapid density

increase

a relatively low temperature followed by a

slower densification by the diffusion-controlled or reaction-

controlled mechanisms(23). C-0.6 specimens show the

The particle size distribution after 20 h planetary ball

milling shows very narrow particle size distribution

(D~:0.38~, D50:0.22p, Dlo:O.13p).

wt%

y203

and 1

wt%

A203

without milling is 2-9

Oxygen Content of a batch of 93

wt%

l85O0C

hh.A

1750°C

1650'

I.,.I.,.,.,.I.I

2Theta (28)

Fig. 1.

XRD

peaks for the specimen C-0.6 with

sintering temperature

Fig. 2. Relative densities of the specimens with increasing

sintering temperature

578

Fig. 3.

SEM

micrographs for the specimen C-0 and C-0.6 sintered at 1850“Cfor (a) 0.5h and (b) 6h

delayed densification until 1750°C which is the melting

point of apatite phase, followed by the rapid increase in

densification due to the increase in liquid content. Note that

the density of C-0.6 specimen at 1750°C is still &5%TD

while that of C-0 is nearly 90%TD. C-0.6 specimens have

the apatite phase

second phase after carbothermal

reduction. The precipitation of crystalline phase such

apatite from glassy phase leads to the decrease in liquid

content. This leads to the retardation of sintering process,

especially, for early stage of liquid phase sintering. Just

after the melting of ap atite at 1750

“c,

the amount of liquid

increases, which triggers the densification process at a

relatively high temperature such

1750

“c

Since the grain

growth and densification processes are thermally activated

process, the occurrence of initial stage sintering at a

relatively high temperature such

1750°C compared to a

general case at 1450°C may give rise to a significant

difference.

Fig. 3 shows the

SEM

micrographs for C-0 and C-0.6

specimens sintered at 1850°C for

0.5

h and 6 h. After

sintering for 0.5 h, C-0 specimens showed dense

microstructures with fine grains which correspond to 97%

TD, while (2-0.6 specimen showed some pores with

80%

TD(Fig.3-a). It

noticeable that large elongated grains

are

already formed even at

80%

(C-0.6,

0.5

h), which can

act

the seeds for the growth of large elongated grains.

With further increase in sintering time to 6 h, width of the

large elongated grains is getting larger and larger while the

matrix grains remain fine(Fig.3-b). These microstructures

are due to the formation of apatite phase that acts

transient phase. The precipitation of apatite phase retards

the densification process at low temperature, and with

increasing temperature, just after the melting of apatite

transient phase, it triggers the densification process with

increased amount of liquid phase.

This

suggests that

suppressing the densification process to a certain high

temperature followed by the reactivation of sintering

process seems to provide a beneficial effect on the

development of a bimodal microstructure.

CONCLUSIONS

1. Carbothermal reduction treatment

powder compacts

of silicon nitride, sintering additives and carbon mixture

takes the oxygen out of Y-Si-Al-0-N glass.

Loss

of oxygen in Y-Si-Al-0-N glass leads to the

precipitation of crystalline oxynitride phases such

apatite

due to the increase in the

N/O

ratio.

Precipitation

apatite phase and its disappearance due

to melting during sintering process result in the retardation

of densification process up to 1750“C, followed by the

rapid densification due to the increase in liquid content.

4. Occurrence of rapid densification process at a high

temperature such

1750

“C

maximizes the size difference

between the growing grains and the matrix grains, and

this

leads to a distinct bimodal microstructure.

579

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580

DESIGN OF SiCN

PRECURSORS FOR VARIOUS APPLICATIONS

Gunter Motz*, Jurgen Hacker and Gunter Ziegler

University

Bayreuth, Institute for Materials Research,

D-95440

Bayreuth, Germany

ABSTRACT

Main applications for preceramic polymers are

ceramic fibres, matrix composites and coatings. The

requirements for the polymer are determined by the

intended application and differ from

a cross-linkable

liquid to meltable and curable or unmeltable but soluble

solids.

For

applying these precursors on a larger scale,

they must be processable by conventional polymer-

processing techniques. Moreover, the starting material

should be reasonably cheap and readily available, and

the synthesis should be uncomplicated.

Therefore, novel tailored Si-C-N precursors were

synthesised for various applications only by using the

same two chlorosilanes but by different reaction types.

Applying the liquid phase infiltration (LP1)-process for

producing CMCs coammonolysis of CH3C12SiH and

CH3C12SiCH=CH2 leads to

liquid cross-linkable

precursor with

high ceramic yield

(75

after

pyrolysis at 1400 "C. Hydrosilylation of the chloro-

silanes

the first step and subsequent ammonolysis

result

meltable solid

110 "C) with a very good

stretchability because of specific ethylene-bridges. This

polymer was also produced in

pilot plant and

continuously meltspun to fibres with small diameters

20 pm). Subsequent curing and pyrolysis in a furnace

1000

"C (N2) leads to non-oxidic SiCN fibres.

The liquid SiCN precursor mentioned above was

modified by an organometallic compound (e.g.

Ti(NR2)4) and resulted in unmeltable but soluble solids.

These polymers are suitable

for

applications

inorganic polymer- and at higher temperatures

ceramic-like-coatings (CLC). With corresponding

precursor solutions different substrates (metals,

ceramics, glass, plastics) with complex-shaped geometry

can be coated by using simple dip- and spray-coating

techniques. In all cases adhesion to the substrate is very

strong. Selected results from the research areas

mentioned above are presented

order to illustrate that

precursor must be tailor-made for successful use

different applications.

INTRODUCTION

The main application areas for preceramic polymers

are ceramic fibres, matrices

composites and coatings.

The requirements for the polymer are given by the

appropriate application and differ from

cross-linkable

liquid to meltable and curable

unmeltable but soluble

solids.

For the broader distribution of these precursors,

they should be processable by conventional polymer

processing techniques. Furthermore, the starting material

should be readily available and relatively cheap and the

synthesis should be uncomplicated.

Polysilazanes are well suited to fulfil the require-

ments

preceramic polymers. Both educts, chloro-

silanes and ammonia, are very cheap products

the

chemistry industry.

most cases only

one

two

step reaction is necessary to get the desired polysilazane

high yield. The degree

cross-linking (state of

aggregation) and reactivity of the polymer is adjustable

with different functional groups.

In this work we report on three tailored new

polysilazanes and their application

matrix in CMCs

and graphite, ceramic fibres as well as ceramic-like-

coatings.

EXPERIMENTAL PROCEDURE

All operations for synthesis and processing of the

precursors were carried out in an inert gas atmosphere

because of sensitivity to air and moisture.

The liquid silazane HPS was synthesised by

coammonolysis of dichloromethylsilane and dichloro-

methylvinylsilane

toluene

described elsewhere

[I].

Because of its low viscosity and high ceramic yield

(Tab.

this precursor is suitable to infiltrate porous

substrates like graphite, and for processing CMCs by the

polymer infiltration route

121.

Hydrosilylation and subsequent ammonolysis of the

chlorosilanes yields quantitatively the ethylene-bridged

colourless, brittle and meltable solid (ABSE). The

melting point strongly depends on the processing

conditions. This precursor was continuously meltspun to

endless fibres with diameters of about

pilot

plant (Fraunhofer ISC, Wiirzburg, Germany) by drawing

speeds up to

500

dmin. Subsequent curing of the green

fibres by electron beam irradiation and subsequent

pyrolysis

an innert gas furnace at

1000

"C leads to

non-oxide ceramic SiCN fibres with very good oxidation

resistance at temperatures up

1500

"C.

Both the synthesis of the HPS and the ABSE

precursor can be transferred easily to

pilot plant (batch

size 40

1).

The reaction between transition metal compounds of

the type M(NR2)4 (here M

Ti,

alkyl) and the

nitrogen function (NH-group) of the silazane backbone

(here HPS precursor) is the basis for the synthesis of the

unmeltable but soluble TIP precursor. The very moisture

sensitive solid can easily be dissolved

solvents like

cyclohexane, octane, toluene

well as xylene.

581