Fiber-Rein forced Composites: C and Sic Fibers

Ceramic Engineering and Science Proceedings

© 2001 The American Ceramic Society

Mrityunjay Singh & Todd Jessen

Composite Processing

Ceramic Engineering and Science Proceedings

© 2001 The American Ceramic Society

Mrityunjay Singh & Todd Jessen

COST EFFECTIVE PROCESSING OF CMC COMPOSITES BY MELT INFILTRATION (LSI-PROCESS)

Walter Krenkel German Aerospace Center (DLR) Pfaffenwaldring 38-40 70569 Stuttgart Germany

ABSTRACT High performance ceramics are still largely produced via powder routes. The main

disadvantage of these monolithic materials is their brittle failure behaviour and the thus dissatisfactory damage tolerance of ceramic components. The most favourable way to improve the fracture toughness of ceramics is the reinforcement with thermally stable continuous fibres. However, long manufacturing times, multiple reinfiltration steps and expensive raw materials lead to high material costs of these fibre reinforced ceramic matrix composites (CMC) which have prevented their breakthrough to terrestrial applications up until now. To overcome these restrictions and widen the applicability of CMCs - in particular, to enter in fields of mechanical engineering - a novel cost efficient manufacturing route has been developed by DLR.

The process is based on the infiltration of a reactive fluid phase into porous carbon fibre preforms. Molten silicon is used as the reactive fluid, which replaces the initial pore volume of the preform and reacts subsequently with the carbon matrix to form silicon carbide. A one-shot infiltration is sufficient for the densification of the matrix for this liquid silicon infiltration (LSI) process. Therefore, short processing times can be achieved which lead, in addition to the use of commercially available uncoated carbon fibres and cheap raw materials like phenolics and granules of silicon, to the lowest manufacturing costs of all CMC materials. This paper deals with the fabrication of so-called C/C-Sic composites and with design and cost aspects for the manufacture of C/C-Sic components.

INTRODUCTION Different processing techniques are currently in use for the development and prototype

production of continuous fibre reinforced ceramic matrix composites, each showing its own benefits in terms of material’s properties and variability in processing. The most common fabrication method of fibre reinforced ceramics is chemical vapour infiltration (CVI). Isothermal CVI processing, however, means long manufacturing periods that can last from weeks to months, due to low deposition rates and multiple rework cycles. Gradient CVI processing reduces the infiltration time considerably, but manufacturing is restricted to rather simple and standardized components like plates or tubes. Alternatively, organometallic polymers like polycarbosilanes and polysilazanes can be used as precursors to form the ceramic matrix during a pyrolysis step. The main drawback of this liquid polymer infiltration process (LPI), however, is the necessity of at least three reinfiltration steps to densify the composite and reach sufficient strength levels. This

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Ceramic Engineering and Science Proceedings

© 2001 The American Ceramic Society

Mrityunjay Singh & Todd Jessen

prolongs the processing time considerably and, as a consequence of the precursor's high prices, results in high CMC costs.

To overcome these inherent restrictions, DLR developed a novel manufacture route for CMC components based on the infiltration of molten silicon into carbonkarbon preforms. Taking advantage of the experiences with the manufacturing of reaction-bonded non-reinforced SiSiC ceramics one expected to reach the following goals:

Almost dense matrix of carbon, silicon and Sic. Dimension-stable manufacture of large and thin CMC components. Fast and simple processing.

First attempts to infiltrate carbodcarbon composites by liquid silicon have been conducted for more than twenty years [ 1,2,3]. As a result of these basic investigations, the carbon fibres have to be coated prior to the infiltration of silicon in order to reduce the degree of fibre degradation. Also, highly graphitized carbon fibres like high modulus fibres (HM) have been recommended as reinforcement for the fibre preform which are more stable in contact with silicon than only carbonized fibres. Both requirements are in contrast with a cost-efficient processing of CMCs and only poor fracture toughnesses have been achieved by the infiltration of commercial carbodcarbon preforms. Therefore, DLR developed new C/C-composites with adapted microstructures which allow the use of uncoated and cheap HT- as well as IM- fibres.

PREFORM FABRICATION The fibre preform fabrication starts with the manufacture of carbon fibre reinforced plastic

composites (CFRP) with matrices of high carbon content. Normally. commercially available resins like phenolics or other aromatic polymers (e.g. XP-60) are used to fabricate laminates by common CFRP techniques like resin transfer moulding (RTM), autoclave or hot pressing and filament winding. After curing, the composites are additionally postcured for the complete polymerization of the matrix. Typical values of a bidirectionally reinforced CFRP component with fibre contents of 60 Vol. % are densities of 1,49 glcm3 and an open porosity of less than 1 %. Subsequently, the CFRP composites are pyrolysed under inert atmosphere (N2) at temperatures of between 900 "C and 1650 "C to convert the polymer matrix to amorphous carbon. In principle, polymer non- reinforced precursors shrink isotropically, by which the extent of volume contraction is dependent upon the mass loss of the carbonized polymer. For example the precursor XP-60 shows in all three spacial axes a more or less equal change in length of approx. 26 % at 900 OC, whereby a volumetric contraction of ca. 60 % within the matrix during pyrolysis occurs. The average mass losses from XP-60 at 900 OC and at 1550 "C lie by 36 % and 39 %, respectively. Thus, at pyrolysis temperatures of 900 "C, the conversion of the precursor to carbon is over 90 YO complete [4].

If the neat resin (NR) has a mass ratio of:

M =*+(e) NR

a density ratio of:

R=l+(?) NN

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and a thickness ratio (under the assumption of isotropic behaviour):

D=l+(?), =g

Mass Density Volume Thickness

(?)),IR (:) (?)NR (?)NR D NR

, -36 YO 0.64 +61.4% 1.614 -60.35 Yo 0.396 -26.5 % 0.735

(3)

then the matrix constants of XP-60 in Table I result:

Whilst mass, thickness and volume decrease, the density of the precursor increases by approx. 61 %. In a non-reinforced state, these high volume changes result in an inhomogeneous body interspersed with cracks.

Two dimensional reinforcements of woven fabrics are preferred for the architecture of the fibre preforms. Under the assumption that dimensional change is prevented by the thermally stable fibres in plane, a macroscopical change in geometry for two dimensional laminates occurs primarily in the thickness, i.e. perpendicular to the fibres. This change in geometry is strongly dependent on interfacial bond strength between the fibres and the matrix and can not be directly derived from the dimensional changes of the individual components.

For the extreme case of high fibre/matrix bond (FMB) strengths, the laminate shrinks theoretically proportionally to the matrix contraction. The change of thickness arising through pyrolysis can then be derived as:

D

#fcm

thickness ratio of the precursor during pyrolysis fibre volume fraction in the CFRP state

For the case of high FMBs the shrinkage in thickness reaches a maximum. The resulting open porosity during pyrolysis can then be calculated according the equation (5) :

For very low fibre/matrix bondings the thickness of the CFRP laminate remains unaffected and the open porosity results in:

e'c/c = (l-V)'(l-@fCFRP) (6)

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The open porosities of the C/C-preform represent the volume which is accessible to the liquid silicon. Figure 1 shows for technically relevant fibre volume contents between 50 % and 70 % the corresponding porosities which amount between 30 % (low FMB) and 11% (high FMB).

30

28

s D 24 i

P 22 f 2o ,B 16

I 8

14

12 J t I I I I I

Fig. 1 correlation between the open porosity e'cIc and the fibre volume content of 2D-CFRP laminates after pyrolysis at 900 O C

High fibre/matrix bondings lead to C/C microstructures which are characterised by translaminar cracks and segments of fibre bundles. These distinct C/C segments are more or less dense and therefore hardly accessible to silicon. The crack pattern, however, acts as a communicating tube system with high capillary forces for liquids of low viscosity. Figure 2 shows the morphology of such a C/C perform which is preferably used for a subsequent silicon infiltration.

Fig. 2 Segmented microstructure of a C/C laminate due to high fibrelmatrix bond strengths (fibre orientation is * 45' / 0' / 90" /* 45' )

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MELT INFILTRATION OF SILICON INTO THE C/C PREFORM Elementary silicon possess several characteristics in contrast to most other metals: Anomaly with phase transition (change in density of approx. 8 %) Extremely low coefficient of expansion (4.1 x 1/K) Highest enthalpy of liquefication of all metals (50 kJ/mol)

The anomaly of silicon means that a volume expansion during solidification of the molten silicon (ca. 1420 "C) occurs and the density of the solid (2,34 g/cm3) is lower than the melt (2,53 g/cm3). For this reason there is a danger from high residual silicon contents, particularly during cooling, that due to the released enthalpy and the volume increase the C/C-SIC laminate can be destroyed.

In order to successfully siliconize, an exact silicon dosage is necessary to ensure a complete filling of the C/C material on one hand, and to avoid, on the other hand, the danger of component destruction through excessive silicon dosage.

The stoichiometrical conversion of molten silicon with solid carbon occurs following the

Density p g/cm3

Molar mass M g/mol Molar volume V cm3/mol

equimolar reaction: 1 mol Si,+l mol Cs CS 1 mol Sics

Silicon Graphite Amorphous carbon Sic 2,53 2,26 1,84 3,22

28,09 12,o 1 12,Ol 40,lO 11,ll 5,3 1 6,53 12,45

(142OOC)

(7)

That means, for the ideal case, one mole of silicon (28,l g) reacts with one mole of carbon (12,O g) to form one mole of silicon carbide (40,l g). The formation of S ic is, with a reaction enthalpy of 4 8 kJ/mol, highly exothermic, so that by the conversion of graphite a violent thermal development with a sharp rise in temperature of several hundred degrees is observed.

The aim of the LSI process is to convert the whole or at least a greater majority of the molten silicon through the reaction with carbon to Sic. The reaction kinetics for the conversion of carbon to Sic is predominantly dependent on the microstructure, or rather the microporosity, than the degree of graphitization. A fast reaction can occur only under the condition of accessible pores and large surface area.

The quantative conversion of carbon and silicon to silicon carbide in a closed system leads theoretically to a reduction in volume, as the molar volume of silicon carbide is less than the sum of the respective molar volumes of the reactants, Si and C (Table 11).

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The conversion of gaphitic carbon to silicon carbide results in a theoretical volume increase of 135 %. For amorphous carbon, with a density of 1,84 g/cm', the theoretical volumetric expansion is still 91 %. Through the larger molar volume of silicon carbide in contrast to the reactant carbon, there is a principal tendency, with increasing reaction times, that cracks and pores close completely.

The infiltration of liquid silicon at temperatures above 1420 "C rapidly fills the cracks within some minutes. The dense matrix within the fibre bundles shields off the liquid silicon. Therefore, a layer of silicon carbide results around the bundles and only a small amount of load carrying carbon fibres are converted and damaged, Figure 3 shows S i c crystals with different sizes, filling the former cracks and forming a ceramic coating around the C/C-segments.

Fig. 3 Layers of S i c crystals with different sizes create a protection coating around the carbon fibres (magnitude 750 X)

As an example, the silicon-uptake of 35 % during a two hours siliconization result in a matrix composition of about 62 % in mass of Sic, 26 % of carbon and 12 Yo of residual silicon. The corresponding values for the phase composition of the resulting C/C-Sic composite are summarized in Figure 4.

The constitution of the matrix and the material's microstructure mainly depend on the choice of matrix precursor and fibre structure, the processing parameters and the fibre/matrix interface. One main advantage of the LSI process lies in the possibility to vary these parameters in a wide range leading to different qualities of C/C-Sic materials.

448

Mass Volume

Fig. 4 Composition of a C/C-SIC standard laminate

Typically, C/C-Sic composites are made of bidirectionally woven fabrics which are stacked together or wound to the desired thickness. An orthotropic behaviour results for the technical as well as thermophysical properties, which show considerable differences between the in-plane and thickness direction.

With respect to the microstructure, composition and characteristics C/C-Sic composites differ from all other structural materials, in particular from monolithic ceramics, and represent a separate class of material. Three C/C-SIC qualities have been standardized which fit to the requirements of different applications (Table 111).

Table 111. Mechanical and thermophysical properties of 2D C/C-Sic composites at room temperature

Property Unit XB XT XD Density lo' kg/m3 1 9 1,9 2,3 Open porosity % 3.5 3 s 1 ,o Interlaminar shear strength MPa 28 33

Flexural strength MPa 160 300 80

Tensile strength MPa 80 190 30

Strain to failure % 0,15 0,35 0,04

Young's modulus GPa 60 60 100

Coefficient ofthermal expansion 1/K 100 OC 11 - 1 - 1 1,s (Reference temperature 25OC) 100"CI 2,5 2 S 4,5

1500°C 11 2,5 2,2 3 3 1500 C I 6 3 7 7 3

Thermal conductivity WlmK 200°C 11 1 8 3 22,6 3 3 , l 2OO0CI 9,O 10,3 18,2

1650 "C 1) 17,O 20,s 23, l 165O"C.l 7,5 8 3 12,4

Specific heat J/kgK 25°C 750 690 720

1400°C 1550 1540 1450

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Due to their low coefficient of thermal expansion, their high thermal conductivity and moderate modulus, C/C-Sic materials show an excellent thermal shock stability. They retain their strength level at elevated temperatures, similar to carbodcarbon materials. Moreover, their high temperature strength is superior to the level at room temperature, i.e. the higher the temperature, the higher the strength (Figure 5) . The maximum temperature of CIC-Sic composites under stationary conditions is limited to 1500- 1700 "C. However, considerably higher temperatures of up to 3000 OC can be applied in cases, where the required lifethe amounts to only a few minutes, e.g. for space or military applications.

100

311u T 2 6 5

ILSS Flcrur.1 Modulus

Fig. 5

DESIGNING WITH C/C-SIC COMPOSITES Designing with C/C-SIC materials principally follows the same rules and requirements as exist

for other composite materials. However, due to the high temperatures during their manufacture (at least 1420 "C) and their heterogeneous microstructure, some specific characteristics must be taken into account for a suitable design of C/C-Sic components. In general, the matrix shrinkage during the pyrolysis step can lead to a significant amount of residual stresses within the component, especially for closed geometries like tubes. If the components can shrink unrestrained, the resulting stress level is much lower but considerable dimensional changes may occur.

High temperature properties of C/C-Sic composites

Generally, volumetric changes during the manufacture of C/C-Sic are determined by the preform architecture and by the volume fraction of fibres. As the carbon fibres possess different characteristics to the matrix, the shrinkage of the composite is direction dependent. Short fibre reinforced C/C-SIC with a random orientation show a more or less isotropic shrinkage in thickness, length and breath in case of e.g. flat plates. In contrast, an anisotropic behaviour is observed if the reinforcing phase is of a bidirectional nature (e.g. fabrics). In this case, the pyrolysis step is accompanied by a macroscopical dimensional change, which due to the shrinkage impediment of the fibres only occurs transverse to the fabrics.

During a 900 "C pyrolysis the laminate thickness shrinks by approx. 4.4 %. The pyrolysis step at 1650 OC under vacuum results in an additional volumetric contraction of 1.74 %. Here, the main geometrical change also occurs in the direction of the laminate thickness. Minor changes (-0.25 %)

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occur in breadth and length. During siliconization, mass increases by 44 % without any essential change in shape and geometry (Table IV).

In total, the anisotropic characteristics of the laminates result in dimensional changes over the whole process of less than -7 %. This shrinkage occurs mainly in one direction, namely perpendicular to the fibre orientation.

Table IV. Dimensional and mass changes occurring within 10 mm 2D-laminates (HTA fibres, XP-60) during the LSI process in relationship to the previous material state

In case of an angled plate the irreversible change in thickness (Ad / d ) ,c leads to a change of the angle p during pyrolysis. According the equation:

this spring forward effect can be calculated for orthotropic composites (Figure 6).

For example, the mean shrinkage during pyrolysis of -4.5 % reduces a rectangular angle by 4.3" to 85.7". Vice versa, to get an angle of 90" in the C/C-Sic composite the CFRP component must show an angle of 94.05'.

45 1

' I ' i

Fig. 6 Spring forward effect of orthotropic reinforced C/C-Sic laminates due to irreversible change in thickness direction

Due to the fact that all dimensional change within CIC-Sic components is complete after the pyrolysis step, the manufacture of complex components is possible through modular construction. Simple building units such as plates, tubes and profiles can be successfully joined within the siliconizing step of the LSI process (in-situ) using a carbonaceous paste with the optional addition of either carbon felt or fabric as the joining material to form complex components [6 ] .

Porous C/C components are prepared and fixed together and molten silicon is caused to flow between the surfaces and react with the carbon material to convert it to S ic and bond the surfaces together. Interlocking the C/C parts prior to siliconizing increases the joining area and leads to a stable assembly which needs no further supports during the reactive bonding in the furnace. In-situ joining is desirable as it eliminates expensive and complicated machining as well as the need for additional metallic bolts or ceramic adhesives. The joining strength level lies nearly in the same range as the shear stresses which can be transferred in the C/C-Sic composite and are therefore sufficiently high for most applications [7].

COST ASPECTS

Generally, the main cost drivers for the manufacture of CMC composites are: Long processing times High fibre costs High machining efforts

Low manufacturing costs are one basic requirement for an industrial use of CMC materials.

In comparison to other CMC manufacturing processes the LSI process offers a big potential to reduce the costs by decreasing the influence of all of these parameters. The total processing time to fabricate C/C-Sic components lasts about 200 hours, including CFRP manufacture, machining and all high temperature furnace steps (Figure 7). An economic study based an a precommercial production of C/C-Sic plates showed, that the most time consuming step for manufacturing one C/C-Sic part represents the pyrolysis, which take roughly 62 '70 of the total processing time. The other substantial share of processing time comes from the siliconizing step, which needs about one quarter of the total time, whereas the green body fabrication of the CFW component (9 %) and the machining step (2 %) play only minor roles.

452

CFRP manufacture

9% I

Siliconizing f 27%

Machining 2%

pyrolysis f 62%

Fig.7 Shares of the processing time for the manufacturing of bidirectionally reinforced C/C-SIC plates

Substantial time savings can be obtained by the optimisation of the state-of-the-art pyrolysis process. Other main reduction potentials lie in a net-shape manufacture of the CFRP body which avoids waste portions and reduces costly machining steps. The fibre costs play no dominant role in the LSI process as comparatively cheap and uncoated carbon fibres of the HT-type can be used One additional possibility to reduce the preform costs is the use of short fibres in cases where the residual strength levels of the C/C-Sic materials are sufficient for the respective application.

Figure 8 shows exemplary the manufacturing for a series production of internally ventilated brake disks of C/C-Sic composites [8]. First, two half shells of CFRP are manufactured in near net-shape by applying the hot-press technique. The subsequently pyrolized parts are joined together and then infiltrated with silicon at temperatures beyond 1420 "C

f \

Net-Shape Manufactured Ventilated CIC-Sic Brake Disk

CarbonlCarbon -Half-Shells

Fig. 8 Patented net -shape manufacture of ventilated, short fibre reinforced C/C-Sic brake disks by using hot-press technique and in-situ joining method

A cost calculation on this process showed, that the costs for raw materials (short fibres, precursor, etc.) and energy again play minor roles. The most critical portion arises from the grinding of the friction surfaces which is still very time consuming. Therefore, automated machining processes and non-destructive evaluation (NDE) methods for a reliable quality

45 3

assurance are the aim of future research to fiuther reduce the actual CMC material costs of approx. 500 Eurokg to a much lower level.

Also, even if the manufacturing costs will be reduced considerably in the near future, CMC materials mostly remain structural materials for niche applications only. To replace traditional materials like metals or polymers in new constructions, a life cycle analysis (LCA) which considers the entire costs, accumulated over the products whole life cycle, must be taken into account as basis for the material’s choice. Only with a such consequent design to life cycle cost (LCC) approach one can take full benefit of the excellent lightweight potential and wear stability, of CMC materials, allowing for example new designs of thermally stable calibration plates for coordinate measuring machines or lifetime braking systems for automotive vehicles.

CONCLUSIONS A novel method has been developed to infiltrate carbodcarbon parts with liquid silicon and to

convert the matrix mainly to silicon carbide. As no fibre coating and no reinfiltration steps are necessary, the LSI process offers the potential of cost effective manufacture of CMC due to low raw material costs and comparatively short manufacturing times. The main geometrical change occur during the pyrolysis at 900 “C and shape stability is achieved after this step. Despite the anomaly of silicon and the distinct differences in molar volume of the reactants, no volumetric change is observed through siliconization.

Simple as well as complex building units can be successfully joined within the siliconizing step of the LSI process (in-situ) using a carbonaceous paste with the optional addition of either carbon felt or fabric as the joining material to form complex components. Porous C/C components are prepared and fixed together and molten silicon is caused to flow between the surfaces and react with the carbon material to convert it to SIC and bond the surfaces together. In-situ joining is desirable as it eliminates expensive and complicated machining as well as the need for additional metallic bolts or ceramic adhesives. Ceramic matrix composites produced via the LSI process are structural materials which retain their strength at elevated temperatures. These characteristics make them useful materials for a variety of aerospace as well as non-aerospace applications particularly under conditions of severe wear or erosion.

REFERENCES

Compo~ites”,4‘~ London Conference on Carbon and Graphite, pp 23 1-235,1974.

Composites”, Ceramic Bulletin, Vol. 54, No. 12, 1975.

‘C.C. Evans, A.C. Parmee, R.W. Rainbow, “Silicon Treatment of Carbon Fibre-Carbon

’W.B. Hillig, R.L. Mehan, C.R. Morelock, V.I. DeCarlo, W. Laskow, ”SilicodSilicon Carbide

3R. Gadow, ”Die Silizierung von Kohlenstoff’, Doctoral Thesis, University of Karlsruhe, 1986. 4W. Krenkel, N. Liitzenburger, “Near Net Shape Manufacture of CMC Components”

Proceedings of the 12Ih Int. Conference on Composite Materials (ICCM- It),Paris, July 5-7, 1999. ’W. Krenkel, “Development of a Cost Efficient Process for the Manufacture of CMC

Components”, Doctoral Thesis, University of Stuttgart, DLR- Forschungsbericht 2000-04, 2000. 6W. Krenkel, T. Henke, “Modular Design of CMC Structures by Reaction Bonding of Sic”,

ASM International Conference Materials Solutions, Cincinnati, USA, November 1-4,1999. ’W. Krenkel, “High Performance Ceramics for Structural Application”, International

Confirence on Composite Science and Technology ICCST/3, Durban, South Africa, January 1 1 - 13 2000.

‘W. Krenkel, R. Kochendarfer, ‘‘ Method of Manufacturing a Friction Element”, US-Patent 6, 086, 814, July 11, 2000.

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