Structure and Properties of Composites

1 Fibrous Composite Materials Anthony Kelly University of Surrey, Guildford, Surrey, U.K.

List of Symbols and Abbreviations 1.1 Composites in Materials Selection 1.2 Economic Background 1.3 Environmental Concerns 1.3.1 Recycling 1.4 Limits to the Mechanical Properties of Fibres 1.5 Control of Cracking in Brittle Laminates 1.6 Novel Forming Methods 1.6.1 Reconstructive Processing 1.6.2 Reactive Processing 1.6.3 In Situ Processing 1.6.4 Metal Matrix Composites and Spray Methods 1.7 Fibre Geometries 1.8 Failure under Multiaxial Stress 1.9 References

Materials Science and Technology Copyright © WILEY-VCH Verlag GmbH & Co KGaA. Allrightsreserved.

2 3 5 8 8 9 11 12 13 14 15 16 17 18 23

1 Fibrous Composite Materials

List of Symbols and Abbreviations b E F,G,H L9M9N V

thickness of 90 lamellae elastic modulus

y

T

work of fracture transverse ply failure stress angle Poisson ratio shear stress (frictional) shear stress

BMC CFRP CTE gdp GNP grp PPTA SMC LCP

bulk moulding compounds carbon fibre reinforced plastic coefficient of thermal expansion gross domestic product gross national product glass reinforced plastics polyparaphenylene terephthalamide sheet moulding compounds liquid-crystalline polymer

£

e V

o

constants area (volume) fraction

1.1 Composites in Materials Selection

1.1 Composites in Materials Selection Composite materials are combinations of materials put together to achieve a particular function. The combinations may be materials of the same class, e.g., two metals, or of a different class, e.g., a glass and a plastic. The generic types of composite are indicated in Fig. 1-1. The advantages to the engineer fall under three headings. First, composites enable a unique combination of properties to be achieved. For example, tungsten wire is very stiff (405 GPa) but very dense (19.3 Mg m~ 3 ). A combination of graphite fibre in epoxy resin is nearly as stiff (306 GPa) but with a density of only 1.5 Mgm" 3 . The carbon fibre itself is much stiffer than the tungsten; values up to 1000GPa at a density of 2.6 M g m " 3 are possible. Besides the attainment of unique combinations of properties, another advantage is that the property can be made to vary con-

tinuously over a range. If only a small number of simple materials are available the variation of a given property among those materials is digitised to the values the individual materials may show. Figure 1-2 shows the variation of coefficient of expansion with volume fraction for a composite consisting of aluminium containing silicon carbide particles. It is seen that matching the coefficient of thermal expansion with that of other materials is easily obtained. The similar non-digitisation of properties is important in, for example, acoustic wave devices and in many other matching situations, e.g., in prosthetic devices. A third important property is that composites can sometimes attain a value of a given physical property not attainable by either of the two components alone. Figure 1-3 shows thermal conductivity of various materials and it is quite clear there that composites can attain a lower thermal conductivity than that of others. This may be regarded as a rather special example since the combination is with air or vac-

METALS Fe,AI,Ni,Cu,Ti,Pbetc Strong,Stiff,Dense,Tough

POLYMERS

CERAMICS C.AI^.SiC.BA C, Hard, High Melting, Stiff, intermediate density

(CH 2 ) n ,(cONH (CH 2 ) 6 ) n etc

[-NH-0-CO-], easily formed, soft but fibres stiff and strong. Low density

GLASS

RUBBER elastomers high extensions

TOUGHENED. POLYSTYRENE

Figure 1-1. Composites are combinations of two or more materials.

inorganic polymeric metallic

1 Fibrous Composite Materials

Thermal Expansion Matching

20-i Brass 18-

J0%

Bronze

^ 1 2 %

Copper

^ 1 5 %

Stainless Steel ^ ^ 1 7 % 16-

^^ 25%

Monel 14-

\ ^

Nickel

X

Iron, Mild Steel, Beryllium

12-

29% 34% ,38%

Carbon Steel

?,

10-

Figure 1-2. Variation of coefficient of thermal expansion with volume fraction

0

10

20

40

30

of silicon carbide particles in aluminium. Ordinate values x 10 "^K" 1 .

Volume % Silicon Carbide in Aluminium Matrix

PSZ ZrO2

Mullite SiO2(q) ice

Si

TiO2

RBSC BeO

AIN SiC (C)BN Diamond C fibre

CERAMICS SiO 2 GLASS (B/Si/O) p-imide CORK

-•

; p-ester -epoxy-

MICROT^ERM

POLYMERS

Cu/p-imide

FIREBRICK

BALSA

glass / p°xy

e

CFRP ( J _ ) CFRP(II) epoxy epoxy COMPOSITES

STILL AIR

SiC/BeO

Tl Li Fe Pb

Na

Ni

CFRP(II)

Cu Au Ag

METALS THERMAL CONDUCTIVITY (Close to R.T.) I 102

10 1

1 ^ ^ I ^ I J

I 10

I 100

Figure 1-3. Thermal conductivity at room temperature of a variety of materials.

1000

1.2 Economic Background

uum as one of the components. However, it indicates the generally important message, which stimulates the imagination, that often empty space can be an important component of the properties of a solid. The second and third advantages of a composite from the point of view of the variation of thermomechanical properties may stimulate the suggestion of new applications. However, the real technological importance of composites is the putative advantage of the fibrous composite over other materials by producing very high strength, very high stiffness per unit weight. This is important in all modes of transportation and in moving bodies. The advantages of such materials though are only being slowly realised in engineering and this is because of the background economic situation which we now describe. I believe it is important to emphasise that this background economic situation is different today from that which has occurred in the previous history of mankind.

1.2 Economic Background Within all developed countries there is now and there will continue to be consistent and mounting public, professional, statutory and commercial pressure to reduce the negative impact of all mining and manufacturing processes and of consumer products on the quality of life. In a word, do more with less. At the same time there is demand for improved living standards, great comfort, ease of transportation and increased leisure. These pressures produce enormous opportunities for the use of performance materials (of which composite materials are the most important subclass) because these are stiffer, stronger and tougher per unit weight than conventional materials, re-

quire less energy to produce them and can, if appropriate measures are taken, be safely recycled. Their surfaces may be engineered, they can be made available as films or fibres or foams. They produce lighter, smaller, more efficient and comfortable machines and structures of all kinds. The tremendous opportunities for the use of performance materials arise against another background, which is one of decreasing importance of all materials as a proportion of gross domestic product (gdp) in all developed countries when this importance is measured as the weight of material consumed per capita as a proportion of the GNP per capita. This point is illustrated by the graphs in Fig. 1-4 for the U.S.A. The curve for the GNP per capita is expressed in both dollars for the year in question and corrected for inflation to 1985 prices. The trend is clear. Steel is the major constructional material. Tin is a material widely used in the electronics industry (30% of tin consumed in the US is in the form of solder), packaging and in the production of bearings and coatings. Rubber is a material very widely consumed in industry and in domestic situations. Synthetic rubber makes up more than 70% of the US consumption of rubber. The consumption of these three materials per head is decreasing while the GNP or gdp per head is rising. The total material input per unit of gdp for these materials is thus decreasing. On the other hand from the data in Fig. 1-5 for polyurethanes and for nylon and for polypropylene it may be said that the rate of growth of consumption can exceed the increase of gdp per head, which indicates expanding use and expanding market and hence the possibility of very profitable commerce. Data for instance for advanced composites show a greater increase in consumption, very much larger than that of the gdp per capita. On Fig. 1-5

1 Fibrous Composite Materials

GNP/Capita($x103)

Consumption/Capita

25

1500

1200

900

600

10

300

1950

1955

1960

1965

1970

1975

1980

1985

1990

1995

Year Figure 1-4. U.S.A. consumption of various commodities 1950-1988.

GNP/Capita ($x103)

Consumption/Capita

10

1960

1965

1970

1975

1980 Year

Figure 1-5. U.S.A. consumption of various plastics.

1985

1990

1995

1.2 Economic Background

composite sales by a major composite company for the years 1985 to 1990 are added as the single dark line, scaled to sales in 1985, and with sales in succeeding years deflated using the GNP deflator for the U.S.A. For all materials then, with the exception of selected performance materials the contribution of materials to the gdp/capita is going down. The reasons for the decrease are (in the case of any particular material, e.g., synthetic rubber in the U.S.A.) waste minimisation, substitution, improvement of properties (do more with less) and static or only slowly growing markets for the artifacts in which the material is used. This last point is very important. World motor vehicle production shows a steady increase up to the early 1970s and thereafter a much slower rate of increase no larger than the growth of the world GDP. The collapse of total demand for merchant shipping is well known; contrast this with the continued increased production of domestic refrigerators. The historic demand for aircraft passenger seats, which I take as a crude measure of the total volume of materials to be used in a big added value engineering industry, from 1975 to 1990 and the prediction by Boeing (who holds more than 50% of the world market) up to 2005 shows that the total demand grew rapidly in the 1980s but; from 1990 on, the total demand of some 140 000 new seats per year is approximately static. The conclusions from these data are that, after nearly 50 years without a major destructive war the world market for machines and structures employing engineering materials generally, and which composite materials are penetrating, or hoping to penetrate, such as automobiles, aircraft, transportation equipment and housing, etc., are mature in their concept (virtually all passenger cars have four wheels) and

the numbers produced are growing only as the world gdp. The major exceptions to this trend that I detect are in artefacts concerned with communication of information (and entertainment) of all kinds, household appliances and many types of medical equipment and prosthetic devices. The consequence of these quasi-static markets worldwide provides both an opportunity for composite materials suppliers but also involves heavy risk. The opportunity clearly arises because the users of composite materials, e.g., automobile makers, see two prime advantages in introducing them, (a) doing more with less, and meeting environmental standards, and lower fuel consumption; (b) obtaining competitive advantage through improved performance, appearance, economy and comfort, and in the case of medical devices producing properties hitherto unattainable. The producers (again, for example, car makers) are, however, under intense worldwide competition with one another. Some may not survive. While struggling to do so, most do not perceive themselves as able to make the necessary capital investments to develop the full capability of performance materials. At the same time they are not willing to take the risk of not replacing their heavy and expensive capital equipment which uses (improved) conventional materials. Against this background, if I am right, the opportunities for performance materials are strategic. This is because the material quality of life ultimately depends on the availability of performance materials a very obvious fact but one which tends to get overlooked in the heat of the present debate - witness the added value in an artificial heart valve. Before proceeding further, I must emphasise that there are very few, if any, unique performance materials solutions to

8

1 Fibrous Composite Materials

a given problem. This is because, with important reservations in detail, the ultimate mechanical and some other common physical properties of metals, plastics and ceramics are nowadays shown to be very similar. This fact coupled with the idea of using a composite composed of two or many more materials increases the range of materials available to a designer as we saw in Sec. 1.1. The idea of a composite also carries with it the notion of fashioning the surface of any article and introducing additional materials on the surface so as to reduce wear, dispel electricity, improve colour, lustre, attractiveness, etc. It invites the idea of a multi-material solution to many design problems rather than choosing a single material. It follows that the preferred commercial solution will depend on history, processability and quality much more than on inherent properties of a particular material.

1.3 Environmental Concerns Two very important factors limiting market growth of any material are: (1) The customer's requirement to achieve more in terms of product performance with less material - do more with less. (2) (a) (b) (c) (d)

Waste minimisation by use of cleanest raw materials; recycling within the factory; extracting energy from waste; using waste to produce valuable products;

while at the same time reducing airborne emissions, waterborne emissions, solid residue and implicit environmental impact such as energy, feedstock, transport and aesthetic effects.

1.3.1 Recycling There are many legislative moves, national and local, to establish comprehensive programmes to encourage recycling and waste reduction. The mounting concern to be able to recycle the materials of a product may lead to a form of "leasing" of the material. Recycling occurs at three stages (i) in the factory; (ii) in a prompt industrial form, i.e., when identified by a fabricator who knows its provenance and its constitution, and (iii) when the artifact is junked. The last form of recycling involves collection either of a single product, e.g., newspaper or a platinum crucible, or more generally, where the material has been incorporated into an artifact, it involves identification, separation from other materials and transportation to a factory. In many cases the concept of artifacts of which the component materials may be recycled conflicts with two important design considerations influencing the use of composite materials. These are: (a) Part integration-elimination of joints, fixings, fittings. Large parts of complicated shape are made efficiently but become very difficult to dismantle. (b) Concepts of multimaterial solutions to engineering problems. These produce components composed of very different materials which are highly effective in use but must be separated for effective recycling. However, plastics scrap can be sorted and some impurities removed. The secondary market for clean regrind of thermoplastics is well established and clean regrind can command prices of 70% of that of virgin quality. All chemical companies are investigating the possibility of finding applications of this secondary material. The recy-

1.4 Limits to the Mechanical Properties of Fibres

cling of plastic fuel tanks of cars, for example, leads to little deterioration of property. However, there is, in general, a loss of property and a cascade principle must be followed, i.e., use of the material in progressively less demanding applications, e.g., bottle, wood substitutes, fuel, in that order. Recycling of organic materials cannot continue indefinitely and beyond a certain level there is no point from the ecological standpoint of expending energy in purifying contaminated plastics scrap and the final repose must be the incinerator or the rubbish dump. Cured thermosetting plastics scrap is usually treated as household refuse. Incineration does not raise fresh problems for sheet moulding compounds (SMC), bulk moulding compounds (BMC) or glass reinforced plastics (grp) polyester compounds. The ground material can be recycled as a filler where the glass fibres retain their length and used to reinforce other plastics or cement, for example. Recycling procedures for advanced composites are particularly worthwhile. For advanced carbon fibre reinforced thermosets fully one-third of purchased material may be lost during manufacture. Fabricators can clearly visualise cost savings if this material is recycled and without this are greatly encouraged to look for alternative materials. There are opportunities in my opinion for sophisticated high value added methods for identification of plastics, cleaning, dismantling of structures (e.g., by use of stress corrosion effects), chemical methods of separating fibre and resin in advanced composites and of disassociating multimaterials generally, including separation of metals and non-metals. Where performance materials are used in medical or agricultural applications or in packaging

of expensive items, then the question of disposal may offer opportunity for biodegradable polymers. The automobile industry, particularly in Germany, is taking recycling very seriously in collaboration with the major chemical companies. Studies embrace dismantling techniques, sorting techniques, cleaning and size reduction. The long-term targets are closed cycles of material with recycling quotas of 30% aimed at. This figure is comparable to the percentage achieved with metals.

1.4 Limits to the Mechanical Properties of Fibres The latest theoretical estimates of the possible stiffness and strength of fibres attainable at room temperature can be compared with experiment. Figure 1-6 compares experimental values (top) with theoretical values for the elastic modulus divided by the density at room temperature for a number of materials. The theory for the elastic modulus of a perfect specimen represents the latest values reviewed by Macmillan (1988) and Macmillan and Kelly (1992) and the density has been taken as the defect-free X-ray density at room temperature. The experimental values are the largest known to the present author and the values of the density are experimental ones. The values of density in both cases have been converted to specific gravity, taking the value for water to be exactly 1. It can be seen that there is still room for improvement in the stiffness of carbon fibres. The absence of an accurate theoretical value for thin layer platelets of metals, e.g., Cu:Ni, is noteworthy - in fact this high modulus is to my

10

1 Fibrous Composite Materials Kevlar 149 Cu:Ni 2 nm layers

4-3

-S13N4

SiC Ccore) C T-300

-Nylon

EXPERIMENT (fibrous materials)

FP PET|

Cu

fSiO 2 20

AI2O37 saffimax 40

30

50

60

(Wcore)

Polyethylene

•rSiC Nicalon

PBT

I P120 -fGPa)

150

70 80 90 100

300

200

400

500

THEORY AIN(0001)

Al

(100)

7 Silica

A,f

NaCU

S

T AI 2 O 3

30

')l

' Polyethyjene

7Si 3 ni 4

rB4C

I

Si 7

TiB 2



Silica

T NaCI

dm

5

6

7

8 9 10

15

20

30

40

50

60 70 80

(GPa)

Figure 1-7. Experimental (room temperature) and theoretical values of the tensile strength divided by the experimental specific gravity at room temperature. For symbols see Fig. 1-6.

1.5 Control of Cracking in Brittle Laminates

11

1.5 Control of Cracking in Brittle Laminates

needed so that the following factors are controlled or their effects known to be eliminated:

Cracking in laminates designed to resist biaxial in-plane stress occurs in the plies transverse to the major principal stress when the laminate is strained (in uniaxial straining) to a strain comparable with the transverse failure strain of a single ply, see e.g., Bailey et al. (1979). It is very important in assessing the load bearing capability in the direction of the major principal stress and occurs in

(1) variations of microstructure of prepreg; (2) temperatures of cure and temperature below which the matrix does not flow and is effectively solid; (3) effects due to differences of thermal expansion coefficient between the matrix and the fibres and between plies of different orientation; (4) variation of transverse cracking strain of a single ply.

(1) uniaxial straining, (2) fatigue conditions under cyclic loading, (3) under temperature change and thermal cycling.

When all of these effects are sufficiently well controlled and/or eliminated then there is very good evidence that, all other things being held constant, the transverse ply cracking strain (stress) of a cross ply laminate is increased if the plies are made thinner, and this effect is important for thickness of plies of around 0.5 mm or less. There are two main explanations. The first is that the failure strain (stress) of brittle bodies is governed by stochastic processes so that thinner plies will be found, on average, to be stronger than thicker plies. The second explanation depends on the energetics of cracks in partially cracked media, where one component, longitudinal plies (or fibres, say), can sustain the full applied instantaneous load when the other transverse ply (or matrix) fails. The first explanation is not easily amenable to quantitative test or rather its proponents perhaps prefer not to subject it to such rigorous analysis. The second has been moderately well developed. What experiments there are satisfying conditions (1) to (4) above indubitably favour the second type of explanation (Kelly, 1988). The quantitative evaluation of the energy approach has been most fully developed by McCartney (1992). Subject to the assumption of a sliding interface between

It also occurs under bending and torsion and other more complicated stress states. The phenomenon has been most explored with thermosetting resin matrices, all of which shrink on curing and few of which are cured at the temperature of test. In addition, many experiments are carried out which use prepreg obtained from the manufacturer or other supplier which is used to build up plies of the laminate and cured by the experimenter who later investigates first ply and other failures. There is very good experimental evidence that nominally similar prepregs have very different microstructures, degrees of cure and even may be partially crystallised! Such effects are well recognised with the resins of higher temperature capability such as the various imides. However, in my opinion, their effects are not sufficiently well recognised with epoxies and polyesters. It may also be noted in passing that glass matrix composites show similar effects depending on the method of fabrication. In carrying out rational analysis and, more importantly, when predicting transverse cracking strains, therefore, care is

12

1 Fibrous Composite Materials

plies with that sliding governed by a frictional shear stress T, a lower limit to the transverse ply failure strain is given, on the simplest assumptions, by ZL c T

(1-1)

where £ c is the elastic modulus of the laminates, V1 and V2 the area (volume) fractions of the 90° and 0° lamella respectively, E± and E2 their moduli, a = Ex VJ(E2 V2) and b the thickness of the 90° lamella(e). The unknowns in this equation are % the shear stress, y the work of fracture and, most important, there is, in general, a lack of knowledge of the origin of the value of s, i.e., the effect of interlaminar stresses, due usually to differences of CTE between the various plies and between fibre and matrix in a given ply; sT above. Recently Pryce and Smith (1992) have gone some way to extracting experimentally the value of T and of the origin from which the strain is to be measured. This has not yet been done for a laminate but has been accomplished for a uniaxial fibre reinforced system. Two sets of specimens are prepared and one of these loaded continuously to failure to yield tangent values of modulus and of Poisson's ratio. Another set are loaded, unloaded and reloaded to progressively higher levels. For each loading and unloading cycle the residual strain (permanent set) is determined from the resultant hysteresis loop. As the specimen is stressed the matrix damage is in the form of an array of cracks spanning the width and thickness of the specimen; the density of cracks increases with increasing applied strain. In Pryce and Smith's experiments, in Nicalon fibre (SiC plus additives) in calcium aluminosilicate glass ceramic cracking commences at a strain of 0.08% and the stress-strain

curve becomes linear again at strains greater than 0.4%. Matrix cracking has saturated at this stage, which is confirmed by crack counting. On unloading from relatively low maximum stresses and reloading several times to the same stress a closed hysteresis loop was obtained, i.e., the residual strain reached a constant value. The crack density also remained constant during the loop. These two quantities were measured and also the maximum stress attained during that particular hysteresis loop. From these measurements values of T for sliding between fibres and matrix (not between plies of the laminate) and of the residual stress (strain) can be found, by assuming a constant sliding friction stress between fibres and matrix and calculating the stress-strain curve over the region where the crack density is increasing. An example is shown in Fig. 1-8. The method gives great promise for a quantitative evaluation of the parameter x in Eq. (1-1) by application of the same method to laminates. Perhaps for clarity one should emphasise finally that the value of T and of the thermal strains may be found if the stressstrain curve is analysed quantitatively over the range of strain where crack density is changing. For a given crack density in the transverse ply the longitudinal modulus, Poisson ratio and thermal expansion coefficient can, of course, be calculated (as McCartney and others have shown) for a given crack density without the need to know T (or the thermal strain).

1.6 Novel Forming Methods The recognition of the composite principle as a means of providing very strong, stiff materials of light weight and high tern-

1.6 Novel Forming Methods

13

OOU"

300«T 250CL

w co 200to T5 150 J Shear strength = 8 MPa
80%) provided specific volume fractions of definite size ratios are used, see e.g., Ritter (1971). The use of a polymer to fashion material to shape and its subsequent removal is the basis of much processing of ceramics. A good example is found in a recent paper by Wright etal. (1990), who have made thin springs from ceramics, e.g., A12O3 and YBa 2 Cu 3 O :c . The suspension consists of some 86-88% (by weight) of ceramic with a mixture of various phthalates with stearic acid and polystyrene and sometimes other additives. This is subject to high shear mixing. The softening point may be controlled by the composition as well as the volatility of the diluent so as to effect removal of the

binder. The physics needed concerns not only the particle sizes of the powder of ceramic but the influence of additives on softening point and volatility. Attention to the range of particle sizes is necessary so as to avoid shrinkage (maximum volume fraction) and to maintain the ability to extrude (minimum volume fraction). 1.6.2 Reactive Processing

It is the difficulty of plastically deforming or even of inducing change of shape of the hard high melting point ceramic materials which otherwise have many desirable properties that leads, I believe, to what I call reactive processing, in which one material is totally transformed into another of the required properties and sometimes shaped at the same time. A process which has attracted a good deal of attention is illustrated in Figs. 1-9 a and l-9b - see Urquhart (1991) - which is a novel technique for fabricating ceramics or composites with a ceramic matrix by the directed oxidation of a molten metal either to produce the ceramic alone or through a mass of filler material if a ceramic matrix composite is required. The figure illustrates the outward growth of a ceramic-metal reaction product from a molten metal exposed to a vapour phase oxidant. Additives will be made to the metal to assist oxidation, e.g., Si, Ge, Sn and P plus a little Mg in aluminium. A wide range of composites comprising both particulate and fibrous filler materials have been demonstrated, e.g., A12O3/A1, A1N/A1, TiN/Ti and ZrN/Zr. The putative attractive features of the technology, in comparison with conventional ceramic technology, include the ease of making a ceramic matrix composite directly at low temperature, the elimination of shrinkage, the potential to make large ceramic bodies and projected low

1.6 Novel Forming Methods

15

Filler With Reaction Product Matrix

Refractory Container

(a)

Refractory Container

(b)

cost. An obvious problem is the final elimination of the metal. Of course the reaction need not be with a material extraneous to the composite (in the above case, oxygen). Tani and Wada (1990 a) describe a process which they call internal synthesis for preparing (particulate) ceramic composites such as SiC reinforced with internally synthesized TiB 2 . The process gets over the difficulty that TiB2 is not usually available in small enough particles sizes and furthermore reacts with water. In one method SiC powder mixed with TiN and amorphous boron was hot pressed at above 2000 °C in an argon atmosphere. The molar content of B is more than twice that of TiN. During hot processing the reaction TiN + 2B -> TiB 2 + Y2N2 occurred at between 1100°C and 1700°C. The excess boron aids the densification (sintering) of the SiC so that the process provides hot pressed SiC with small TiB2 particles dispersed within it. The necessity to eliminate a gas is, of course, a disadvantage, if hot pressing has to be used and a proposed process shows promise in which the TiB 2 is synthesised by a pressureless sintering reaction involving SiC, TiO 2 , B4C and car-

Figure 1-9. (a) Schematic of formation of a ceramic-metal reaction product from a molten metal and vapour phase oxidant. (b) Schematic of (a) occurring through a mass of filler material (possibly fibres).

bon. Again excess B4C and C are needed as sintering aids for the SiC (Tani and Wada, 1990b). Of course reaction injection moulding to make reinforced plastics is a very wellknown process. 1.6.3 In Situ Processing Most types of strengthening and stiffening methods used in metal and ceramic processing consist of the introduction of an extraneous phase, for example to block dislocation motion. The same is true of polymers par excellence, where a filler stiffens and, of course, in a composite, reinforces as well. In situ processing to produce the reinforcement directly within the matrix has been accomplished in the metal and ceramic case by directional solidification of eutectics, where the temperature gradient at the solid liquid interface and the rate of growth are both closely controlled so as to produce a rod or lamellar eutectic of the desired orientation. The process is well described by McLean (1983). The disadvantages of this elegant conception is that the volume fraction of reinforcing phase is controlled by the equilibrium phase dia-

16

1 Fibrous Composite Materials

gram and although the addition of other elements may vary this a little, a large change is impossible. Other problems needing to be attacked are that the processing is very sensitive to the presence of trace elements whereas a viable commercial method should not require super purity of the ingredients. The other problem - one where a better understanding of the physical limitation of the process would be useful - is that the current rates of production possible with eutectic alloys is very slow relative to directionally solidified and single crystal superalloys. In situ processing also naturally suggests itself as a means of introducing reinforcement directly into a polymer. Takanyagi etal. (1980) describe the production of a thermoplastic (nylon 6 or 66) containing fibres of polyparaphenylene terephthalamide (PPTA) directly, essentially by polymerising the nylon and then extruding the sulphuric acid solution of the reaction product into a large amount of water and ethanol. The molecular mass and the aspect ratio of the rigid rod component (similar to some form of Kevlar) has of course a marked effect on the mechanical properties. As with metallic eutectics, processing speeds would need to be greatly increased if the method were to show real commercial possibilities. Prevorsek and Mills (1989) discuss other possibilities and point out why melt processing is not possible in this case - see also Paulokowsky et al. (1991). The problems to be solved are related of course to processing of rigid rod polymers which form liquid crystals - lyotropic in this case. The high speed processing of such crystals to produce high strength fibres such as Dupont's or Akzo's aramid fibres is well known. This question also naturally arises of the use of liquid crystals to assist the process-

ing of more conventional polymers. These are the thermotropic liquid crystals which are used to produce moulded parts with a nice surface finish. Tupperware made of Amoco's Xydar is a good example and others are Vectra LCP of Celanese. So-called structural liquid crystal polymers are thermoplastic polyesters or polyamides of aromatic acids, amines, or hydroxyl compounds. They retain some order in the molten state. On solidification they tend to be more highly ordered than other polymers and so have superior properties. Sometimes these can be deliberately increased by incorporation of rigid rod aromatic molecules. The advantage is that they show much lower viscosities than ordinary thermoplastics at similar molecular weights, though with rather high processing temperatures: 340-370 °C. Of course such polymers can be filled with up to 50% of glass or talc or other filler. LCPs compete with the higher temperature resins for use in electrical components, where they must remain stable at soldering temperature. There are many interesting problems in the rheology of such materials. See for example Doraiswamy and Metzner (1986). 1.6.4 Metal Matrix Composites and Spray Methods It is the ease of casting processes that seems to be the key to the commercial viability of metal composites. The advantages in specific strength and stiffness of the material incorporating ceramic particles or fibres over the unreinforced metal are well known. The overriding considerations to achieve commercial success are safety, speed (hence low cost) and subsequent processability of the composite by conventional means. Continuous fibres are too expensive, whiskers are toxic and so the

1.7 Fibre Geometries

reinforcement preferred is a flake form. Consideration must then be given to the size and shape and the method of introduction. Particles about 10-50 jim in diameter and with aspect ratio 1 -10 are considered. The possible methods of introduction are powder blending, mechanical alloying, melt stirring, compocasting and spray codeposition. The third and fifth of these are the ones chosen by the major aluminium producers. In melt stirring, the particles (up to 20% volume fraction) are stirred into the molten material under vacuum by a high shear mixer (at ~ 600 °C!); the problems are clearly avoidance of oxidation and the presence of water. Cospray involves clever design of the atomiser of the aluminium alloys into drops - mainly between solidus and liquidus when the substrate is struck - and 15 jum diameter particles of SiC (say) are introduced and the aluminium droplets are maintained at 100 jim size for safety reasons. A nice set of physical problems arise concerning the process of extracting heat from the aluminium - can one alter the latent heat of melting? - and again the geometry of the relative particle shapes and sizes to obtain the various packings. The material builds up on the collector almost without voids, although a subsequent extrusion is necessary in order to eliminate all porosity. Without the introduction of a ceramic filler or reinforcer, the method has great advantage for conventional powder processing since the powder (spray) is produced and deposited with very little porosity in one operation. Most processes are developments of the so-called Osprey method invented about 1977. The atomisation produces particles with a density in the spray of that of a heavy gas (a few grams per litre) moving very rapidly at 50100ms" 1 , i.e., at gas kinetic velocity. The process physics is fascinating, involving

17

measurement of the particle size, speeds, number per unit volume, state (i.e., solid, liquid, mushy) and the proportions of these and problems of course of the extraction of heat and the effects of impact.

1.7 Fibre Geometries Fibre composites require application in shapes other than thin sheets stressed in their plane so as to overcome poor damage tolerance and delamination resistance. The need to consider three-dimensional arrays of fibres arises and interest is renewed in the physics of, and computer control of weaving and braiding processes. Some surprisingly simple geometrical arrays do not appear to have been fully explored in the literature open to composite specialists. Questions arise as to possible maximum density of packing of fibre bundles of various cross sections. If straight fibres are arranged along three orthogonal directions, the maximum packing fraction attainable is 0.75 (or 3A) for fibres of square cross-section of 3 TI/16 = 0.589 for fibres of circular cross-section. Such a structure yields easily in shear parallel to any of the (orthogonal) axes and in any direction on these planes. It is well known to structural engineers that the rigid frame most economical of material to resist a given set of forces is tetrahedral (Michell, 1904). An array of uniform fibres aligned along the direction normal to the faces of a regular tetrahedron has a surprisingly high packing fraction of 0.5 (or Vi) provided fibres of triangular section are provided. Such an array resists all tensile stresses and most shears except for those in highly specific directions on specific planes. The directions turn out to be the edges of the tetrahedron defining the directions of the fibres. Fibres arranged solely along these directions form

18

1 Fibrous Composite Materials

a dodecahedral arrangement with the fibres running normal to the faces of a rhombic dodecahedron. This dodecahedral arrangement has fibres running in six directions in space and among these six directions no more than three lie in any one plane. An array in these directions is such that any general strain can be resisted by the need to extend at least one set of fibres and there is no easy direction of shear or of extension. This would be called in weaving parlance a 6-D arrangement. Questions then arise such as: (i) What is the densest array of fibres possible in this 6-D arrangement? The answer is 1/3. (ii) What is the cross section of fibre bundle necessary to achieve this packing fraction? Answer: rhombic. (iii) What is the simplest technique of weaving or braiding that can attain this structure? (iv) Can this technique be followed by a machine? Studies of this type are important because it may not be, for instance, necessary to resist a general strain by elongation of one set of fibres. A lack of resistance to shear in a very small number of clearly defined directions may not be of great importance. Then it would follow that perhaps the tetrahedral arrangement of fibres, where a packing fraction of lA in a 4-D arrangement can be attained, would be suitable in a particular application. For all these arrangements the tensile and shear moduli must be estimated without the complication of fibre-fibre interaction. In all of these considerations there is a developing need to be able to display in two dimensions not only the directions in which the fibres are to run but also the cross-sectional form of the fibre. Of course,

a stereogram will represent in two dimensions directions arranged in three dimensions in space, and this representation is what is called angle true. Parkhouse and I find that the stereogram is of course fine for directions but is not good for finding the gaps between the arrays of fibres running in the various directions.

1.8 Failure under Multiaxial Stress As high performance composites go into service in critical situations, there is being recognised an increasing need for soundly based criteria for failure under multiaxial stresses. Physically satisfying theories are of two types only. A maximum stress criterion where it is envisaged that there is a certain upper limit to the force (identified as producing either a shear or a tension) that a solid can withstand. The second is based on the idea of a maximum failure strain, whereby it is envisaged that the relative displacement of the component parts of a solid cannot exceed a certain amount if the cohesion of the solid is to remain. In order to clarify the present situation it may be appropriate to review quickly the classic case of yield criteria in metals. The Tresca yield criterion states that yield will occur (not fracture) when the magnitude of the maximum shearing stress has a value (

babil

n -

' J^.—-^

-2 -

m = 100

-4 / / 0.5

10% £

1% 1

10 i

i

1.0 Stress

1.5

2.0

Figure 2-2. Plots of Eq. (2-8) for various values of the Weibull modulus (from van der Zwaag, 1989).

32

2 Fibers and Whiskers

Table 2-1. Weibull shape parameter for several reinforcing fibers and whiskers. Fiber

Nicalon SiC

SCS-6(CVD-SiC) Carbon Type I Carbon Type II Boron (W core)b Boron (W core)c Kevlar49 d Fiber FP SiC Whisker a

Heat cleaned;

b

Gage length (mm)

Weibull modulus, m

References

2 10 10-220 10-175 3 40 0.5-50 0.5-50 40 40 — — -

8 5.5 2.2-3.6 2.9-3.4 10-16 5.0-8.1 3.8-8.0 ~7 -19 ~9 6.5 1.77

Warren and Andersson, 1980 Andersson and Warren, 1984 Bunsell et al., 1988 Wu and Netravali, 1992 Le Petitcorps et al., 1988 Hitchon and Phillips, 1979 Hitchon and Phillips, 1979 Le Petitcorps et al., 1988 Le Petitcorps et al., 1988 Wagner et al., 1984 Nunes, 1983 Petrovic et al., 1985

140 urn, uncoated;

c

140 urn , thick BdC coating;

d

fiber yarn.

Table 2-2. Values of the parameters of modified Weibull distribution for several fibers (Phani, 1987, 1988 a, b). Fiber

E-Glass S-994 Glass Nicalon SiC Fiber FP (A12O3) Carbon Type I Carbon Type II

Gage length (mm)

(GPa)

(GPa)

(GPa)

(GPa)

15-120 0.25-60 10 254 50 50

6.0 6.07 16.0 3.0 9.9 10.15

0.62 0.55 0 0 1.0 1.0

0.55 0.28 2.0 0.2 2.4 2.3

2.07 4.48 4.5 1.2 5.0 5.0

1.13 1.29 2.88 0.05 0.41 1.41

5.30 5.98 4.66 6.19 12.18 13.07

where au and 99%a-Al 2 O 3

20

3.95

1.38

379

7.0

0.35

200

PRD-166*

DuPont

80%aAl 2 O 3 -20% ZrO 2

20

4.20

2.07

379

9.0

—

200

Sumitomo

Textron/ Sumitomo

85% Al 2 O 3 -15%SiO 2

-17

3.25

1.50

200

4.0

-

500,1000

Mitsui Mining

99.5% a-Al2O3

10

3.60

-1.80

-280

7.0

—

Nextel 312

3M

62% A1 2 O 3 -14% B 2 O 3 -24% SiO2

11

2.70

-1.72

152

3.0

1.2

130,400,780

Nextel 440

3M

70% Al 2 O 3 -2% B 2 O 3 -28%SiO2

11

3.10

1.4-2.1

193-241

4.5

1.2

130,400,780

Nextel 480

3M

70% Al 2 O 3 -2% B 2 O 3 -28% SiO2

11

3.05

-1.90

220

4.5

-

130,440,780

Saphikon

Saphikon

a-Al2O3 (Sapphire)

150

3.96

2

480

-

-

NA

ICI

95% Al 2 O 3 -5% SiO2

3

2.80

1.00

100

—

—

NA

Almax

* DuPont has recently discontinued the production of this fiber.

ro

Alumima and Aluminai-Based Fibe

Saffil (Short Fiber)

1000

CJ>

62

2 Fibers and Whiskers

is conducted in two steps. First is a low temperature treatment (400-500 °C) to dry, remove the fugitive species, and decompose the salt to aluminum oxide. During this stage, the hydrated species lose water and form polymeric networks of Al-O-Al and ultimately colloidal species. As a result of the removal of water and additives, the fibers become very porous. The second heat treatment is performed at temperatures above 1000° to remove the porosity and induce crystallization. The final structure and hence properties of the fiber are critically determined by the conditions, e.g., time, temperature, and atmosphere, of these post-spinning heat treatments. During the high temperature firing the amorphous alumina crystallizes by going through a series of transition phases until it assumes its most stable form which is oc-alumina. The sequence of transitional phases that may occur in alumina is shown in Fig. 2-23 along with associated changes in some properties. It must be noted, however, that in pure alumina it is extremely difficult to control the conditions such that the above progression through the phases can be observed. Such a controlled progression can only be achieved, and some of the transition phases stabilized, through the addition of a second phase such as silica. From a phase stability point of view it is desirable to produce fibers of pure oc-alumina. It is also desirable for the fiber to have low porosity and small grain size that does not grow significantly with temperature. Unfortunately, in pure alumina the

conditions required for low porosity and small, stable grain size are mutually exclusive. Grain boundaries in pure alumina are extremely mobile which leads to extensive grain growth during the high temperature firing required for the removal of porosity or in high temperature service applications. Grain growth can be inhibited by the addition of a second phase such as acidic oxides of Si, P, B and Zr. The price to pay, however, is the fall in fiber stiffness and maximum use temperature due to reduced crystallinity. Alumina fibers currently produced are mostly of a, y or 8 types, a-alumina fibers incorporate a small amount of MgO as a grain growth inhibitor, a-alumina, which is the densest form of A12O3, has a hexagonal structure with oxygen anions occupying the regular close-packed hexagonal lattice sites and the small aluminum cations positioned at 2/3 of the octahedral interstices. It is thermodynamically very stable and thus has a nonreactive surface. 7-Alumina has the structure of spinel, MgAl 2 O 4 , (Ervin, 1952) and occurs as very small crystallites (about 0.01 jim). Fibers of y-alumina contain silica as the stabilizing second phase. Their structure consists of small y-alumina crystallites embedded homogeneously in a matrix of amorphous silica. y-Alumina has a lower density and lower modulus of elasticity than oc-alumina, and its surface is active. 8-Alumina falls between a and y phases; its crystallite size is larger than that of y-alumina and reaches about 500 A. Compared to a-alumina, both y and 5 phases are less prone to grain

T|/y -

Small

- Crystallite Size -

Large

High

-Pore Volume—

Low

High

-Tensile Strength-

Low

Figure 2-23. The sequence of transitional phases of alumina and associated changes in some properties.

2.7 Alumina and Alumina-Based Fibers

63

2000

1500 J

i

K

1000 _

55

I

o

s

Figure 2-24. The variation of strength versus temperature for sapphire loaded along the C-axis (from Shahinian, 1971). 400

800

1200

1600

2000

Temperature ( C )

growth during high temperature exposure. However, the amorphous silica phase in these fibers softens at temperatures around 1100 °C and leads to a drop in strength and modulus. 2.7.2 Sapphire Fibers

Sapphire filaments are single crystalline a-alumina fibers. They were first produced in 1967 by Tyco. The production method involves drawing the filament from molten alumina contained in a molybdenum crucible. The crystallographic orientation and shape of the fibers is controlled by the geometry and orientation of the sapphire seed used to draw the filament. Sapphire filaments with c-axis perpendicular to the fiber axis have been shown to have excellent tensile creep resistance (Tressler and Barber, 1974). The orientation of the sapphire is critical to its thermomechanical stability. Small misorientations may cause a significant drop in strength at temperatures above ~ 1000 °C due to occurrence of uninterrupted slip along the crystallographic planes. Sapphire fibers have received increasing attention as reinforcements for ceramic

matrix composites because of their inherent oxidation resistance and excellent thermomechanical stability. The early sapphire fibers had a diameter of 225 jim, Young's modulus of 275 GPa and tensile strength of 2.1 GPa. New sapphire fibers with a diameter of 140-150 jim and Young's modulus of ~ 470 GPa are currently commercially available from Saphikon. The strength of sapphire filaments as a function of temperature (Fig. 2-24) exhibits three distinct regions with a minimum occurring around 600 °C (Shahinian, 1971). In the lower temperature range of 25-800 °C, it has been proposed that strength is controlled by stress corrosion at surface flaws (Charles and Shawn, 1962). 2.7.3 Polycrystalline a-Alumina Fibers

The first commercially produced alumina fiber was Fiber FP develop in 1974 by researchers at E. I. DuPont de Nemours & Co. (Dhingra, 1980). Fiber FP is 99% a-alumina with an initial grain size of 0.5 |im and a diameter of 20 jim. The crystallinity of the fiber gives it a rough surface as shown in Fig. 2-25. It is prepared by extruding a slurry of oc-alu-

64

2 Fibers and Whiskers

! 2/

Figure 2-25. The surface morphology of alumina fiber FP (from Dhingra, 1980).

mina particles in an aqueous solution of an aluminum precursor such as aluminum oxychloride. The a-alumina particles are added to seed crystallization during the high temperature firing. Fiber FP is also available with a SiO 2 coating which is applied to facilitate the handling of the fibers and to reduce its strong bonding to metal matrices. SiO2 has been shown to improve the room temperature strength of the fiber (1900 MPa for sized vs. 1380 MPa for unsized) by removing the surface roughness present due to the alumina grain boundaries. However, the softening of SiO2 at high temperatures can be detrimental to the elevated temperature strength of the fiber. Being almost pure oc-alumina the temperature capability of Fiber FP is rather limited as grain growth occurs at temperatures above ~ 1000 °C, leading to considerable fall in the strength. The presence of MgO as grain growth inhibitor in the fiber also leads to plastic deformation and soft-

ening around 1000 °C and thus drop in strength. Another limitation of Fiber FP is its low elongation at break of ~ 0.35%. The strength of Fiber FP is very sensitive to flaws and demonstrates a Weibull distribution with respect to gauge length with a Weibull modulus of 6.5 (Nunes, 1983). In an effort to solve the above problems with the intention to extend the application of the fiber to high temperature ceramic composites, DuPont has introduced a new alumina fiber under the developmental designation PRD-166 (Romine, 1987). PRD-166 contains 15 to 20% yttria partially stabilized zirconia grains in a polycrystalline matrix of a-alumina matrix with a grain size of 0.5 |im. Like Fiber FP, it has a diameter of 20 jim but exhibits an almost 100% improvement in its failure strain and ~ 50% improvement in its strength over Fiber FP. The transformation of zirconia from a tetragonal to a monoclinic phase may be providing this improvement (see also Chaps. 7 and 8 in

2.7 Alumina and Alumina-Based Fibers

Vol. 11). The presence of 0.1 jim zirconia grains in the alumina microstructure also inhibits grain growth thus giving the fiber a better high temperature strength retention capability. PRD-166 fibers may also receive a silica coating to remove the surface roughness and to enhance wettability by metals. It has been noted that the silica coating increases the room temperature strength of the fiber by about 25% but reduces the high temperature strength by 15% at 1000 °C and 25% at 1200 °C (Pysher et al., 1989). A new a-alumina fiber has recently been introduced by Mitsui Mining Company under the tradename Almax (Horiki et al., 1978). Almax is 99.5% a-alumina in polycrystalline form. It has a diameter of only 10 |im which renders it suitable for textile preforming. It is available in untwisted, woven and braided forms as well as chopped fibers of various lengths. Despite its a-alumina structure, a maximum use temperature of 1400 °C has been cited by the manufacturer. Another high temperature alumina fiber has recently been developed by Interatom GmbH in Germany. This fiber has no organic additives and is thus expected to show good thermomechanical stability. 2.7.4 Alumina-Based Fibers

Sumika alumina fiber, which is commercially available from Sumitomo, contains 85% A12O3 and 15% SiO2 and is produced from an aqueous solution of an organo-aluminum precursor (Horikiri et al., 1978; Abe et al., 1982); the processing steps are shown in Fig. 2-26. The presence of 15% SiO2 completely stabilizes alumina in the y-form up to a temperature of 1127 °C in both inert and oxidizing environments. The microstructure of Sumika fiber consists of very small crystallites of y-alumina

65

\ Polymerization

Preparation of spinning mix

t Spinning mix

Dry-spinnig

Precursor fiber (polyaluminoxane and alkyl silicate)

\ Calcination

Alumina fiber

(85wt%AI2Q3-

Figure 2-26. The processing steps for the fabrication of Sumika alumina fiber (from Abe et al., 1982).

(about 0.01 jim) embedded uniformly in a matrix of amorphous silica. The softening of this amorphous silica phase above 800 °C leads to a gradual fall in strength and a more pronounced fall in the modulus up to 1150°C (Lesniewski et al., 1990). Above 1150°C the fiber retains little strength. Transformation of the alumina/ silica structure to mullite (3 A12O3 • 2SiO2) starts above 1127 °C and becomes very vigorous around 1400 °C. The variations of srength and elastic modulus with temperature are shown in Figs. 2-27 and 2-28. A series of alumina-based fibers containing A12O3, B 2 O 3 and SiO2 are also produced by 3 M Company under the tradename Nextel (Snowman and Johnson, 1985; Johnson et al., 1987). Nextel fibers with a diameter of 11 |nm are prepared by

66

2 Fibers and Whiskers 1.2 1.0

Q.

£ 0.6 OAir • Argon

0.2 0.0

200

Each point represents the average of 30 tests.

400

600

800

1000

1200

1400

Temperature (C)

Figure 2-27. Variations of the strength of Sumika alumina fiber with temperature (from Lesniewski, Aubin and Bunsell, 1990). 200

small crystallite size of mullite results in good strength and fiber flexibility. Nextel 480 has been reported to stand ceramic composite processing temperatures of up to 1300°C and service temperatures of up to 1200 °C when used with suitable matrix materials (Snowman and Johnson, 1985). The room temperature strength of Nextel 480 is reportedly higher than those of FP or PRD-166 fibers partly due to its smaller diameter (Pysher et al., 1989). All three fibers, however, exhibited appreciable strength and modulus reductions at temperatures above 800 °C and none retained any significant strength at 1300°C.

150

2.7.5 Polycrystalline Alumina Staples 100 LU O Air • Argon

50

0

200

Each point represents the average of 30 tests.

400

600

800

1000

1200

1400

Temperature (C)

Figure 2-28. Variation of the elastic modulus of Sumika alumina fiber with temperature (from Lesniewski, Aubin and Bunsell, 1990).

a sol-gel process and are available in both unsized condition and with an appropriate sizing for textile handling. Fibers have an oval shape with the major diameter almost twice the minor diameter. Nextel 312 consists of 62% A12O3, 14% B 2 O 3 and 24% SiO 2 . Its structure is composed of crystallites of aluminum borate and some mullite in an amorphous silicate phase. This makes it useful up to a temperature of -1000 °C. In Nextel 440 and 480 the concentration of boria is increasingly reduced and alumina increased. The reduction of boria and increase of alumina promotes the formation of mullite and raises the elastic modulus but also the density. Nextel 480 is virtually all mullite with a 3 to 2 mole ratio of alumina to silica. The

Polycrystalline alumina fibers are also available in short staple form. An example is Saffil fibers produced by ICI Ltd. by aluminum salt decomposition (Birchall et al., 1985). Saffil fibers have a mean diameter of about 3.5 jim and a composition of about 95% alumina and 5% silica. Upon heating the as-spun fibers, r|-alumina appears as the first crystalline phase. Further heating at higher temperatures leads to the conversion of r|- to y- and finally 5-alumina with a grain size of about 500 A. The fibers may also contain some alpha alumina. Saffil fiber was originally produced in the form of random mat nonwoven blanket as a thermal insulator. The staple form, which can be as long as several centimeters, later found use as reinforcement for metals and ceramics.

2.8 Silicon Carbide and Silicon Carbide-Based Fibers 2.8.1 Introduction

Silicon carbide (SiC) is an excellent candidate for reinforcement of ceramics and

67

2.8 Silicon Carbide and Silicon Carbide-Based Fibers 0 -1 -2

Passive

_ ^

-3 -4 . -5 .

Active

Figure 2-29. The active-passive transition in the oxidation behavior of SiC (from Antill and Warburton, 1971).

-6 For reaction SiC (s) + 2SiO ^ tr 3SiO(g) + CO (9) -7

i

1100

i

1150

i

i

1200

1250

1300

1350

1400

Temperature (°C)

some metals because of its high strength and stiffness, good thermomechanical stability, and lower density and coefficient of thermal expansion than refractory oxides. Like some metals and ceramics, SiC has an interesting oxidation behavior (Costello and Tressler, 1986; Narushima et al, 1989). In the presence of sufficient levels of oxygen, as in air, and below a temperature of ~ 1200°C, SiC undergoes "passive" oxidation in which a stable SiO2 film forms according to the following reaction: 3 O 9 - + 2 S i O 7 + 2 C O (g)

Virgin SiC Fiber Heat Treated for 54Hrs at 1200°C Passive Oxidation, POa = 10~ 3 atm

(2-15)

The SiO2 film acts as a diffusion barrier and slows down further oxidation of SiC markedly. At very low oxygen levels or high temperatures, however, SiC exhibits "active" oxidation by the following reaction: SiC + O 2 -> SiO(g) + CO

fibers exposed to passive and active heat treatments. The SiO2 surface layer formation under conditions of passive oxidation gives fibers a smooth appearance while active oxidation results in a rough, pitted surface.

(a) Virgin SiC Fiber Heat Treated for 180 Hrs at 1200°C Active Oxidation, PQ2 = 10 ~ 8 atm

(2-16) (g)

In this case a stable oxide layer does not develop, and the oxidation rate increases with increasing partial pressure of oxygen until it finally drops when the oxygen partial pressure becomes large enough so that an active/passive transition takes place. Conditions for the active/passive oxidation of SiC are shown in Fig. 2-29. Figure 2-30 shows the surface characteristics of SiC

(b) Figure 2-30. The surface appearance of Nicalon SiC fiber oxidized under (a) passive and (b) active conditions (from Sands and Parvizi-Majidi, 1992).

68

2 Fibers and Whiskers

scale and their properties. The processing, microstructure and thermomechanical stability of some of these fibers are discussed below.

SiC is not easily sintered. Therefore, its production in fiber form through the conventional inorganic processes is very difficult. SiC fibers are at present produced either by chemical vapor deposition (CVD) or from a polymer precursor. CVD-derived SiC monifilaments grown onto a tungsten or carbon substrate have been in existence since the mid 1960s. These fibers exhibit good thermomechanical stability due to their stoichiometry p-SiC structure, but have a large diameter which limits their effectiveness for toughness enhancement in ceramic matrix composites and precludes textile preforming and net-shape manufacturing of complex geometries. The polymer precursor route is not only potentially cheaper than CVD method but is currently the only method for the production of small diameter SiC fibers. Table 2-12 lists SiC fibers currently available on a commercial or experimental

2.8.2 CVD-Derived SiC Filaments Chemical vapor deposited SiC monofilaments were first fabricated in the mid 1960s by General Technologies Corporation using the processing technology developed for boron fibers. The production of SiC fibers, however, was less costly than that of boron fibers due to the cheaper starting materials and greater deposition rates. In 1972, AVCO Corporation started the commercial production of a CVD-SiC fiber with a tensile strength over 3 GPa (DeBolt and Krukonis, 1973; DeBolt et al, 1973). These early fibers used a tungsten wire, about 12.5 jim in diameter, as the substrate material. It was found, however,

Table 2-12. SiC-based fibers and their properties. Designation

Producer

Composition

(wt.%) p-SiC Textron p-SiC Textron Nippon Si-C-O Carbon a Nicalon (HVR) Nippon Si-C-O Carbon a Nicalon (LVR) Nippon Si-C-O Carbon a Carburandum Carburandum oc-SiC SiC p-SiC Dow Corning Dow Corning SiC X9-6371 HPZ Dow Corning 57% Si-28% N 10% C-4% O UBE Ind. Tyranno Si-Ti-C-O Lox-M UBE Ind.b 38% Si-2% Ti50%C-10%O SCS-6 SCS-9 Nicalon (CG)

Marketed in U.S.A by Dow Corning;

b

Diameter Density Tensile Young's strength modulus

Coef. Elontherm. gation expan. ( x K T ^ C T 1 ) (%)

(g/cm3)

(GPa)

(GPa)

143 78 12-15

3.00 2.80 2.55

-3.8 -2.9 2.97

-400 -330 193

4.2 — 4.0

_ — 1.5

12-15

2.32

2.93

186

—

1.4

15

-2.50

2.93

186

—

—

30-50

3.1

1.38

415

—

—

8-10

2.9

2.28

393

-3.5

—

11

2.4

2.24

206

3.0

1.0

8.10 8-10

2.35 -2.50

3.38 —

193 —

3.10 —

— —

(urn)

Marketed in U.S.A. by Textron Speciality Materials.

2.8 Silicon Carbide and Silicon Carbide-Based Fibers

that the tungsten core limited the thermal stability of the fiber as tungsten reacted with SiC to form W 2 C and W 5 Si 3 . The reaction layer developed during filament production thickened considerably when the fiber was thermally treated at temperatures above 1000 °C thus leading to significant strength reductions (DeBolt et al., 1974; Crane and Krukonis, 1975; Lindley and Jones, 1975). As a result, the tungsten substrate was later replaced by carbon which not only provides better thermodynamic stability with SiC but is also cheaper and lighter (Q = 19 g/cm3 for tungsten and 1.8 g/cm3 for carbon) than tungsten. CVD-derived SiC fibers are deposited from a mixture of silane gasses and hydrogen. The process involves drawing the resistively heated W or C wire substrate through a cylindrical reaction chamber. The gas mixture is then introduced into the chamber and allowed to react and deposit SiC onto the surface of the substrate which is maintained at a temperature above 1200 °C. SiC fibers are currently produced at a temperature of ~ 1350 °C. The rate of reaction and the composition and structure of the resulting fiber are critically dependent on the composition, pressure and supply rate of the precursor gases (Martineau et al., 1984; Weiss and Diefendorf, 1973) as well as the deposition temperature. Higher deposition temperatures increase the rate of production but lead to a coarse, and hence weak, crystalline structure. Low temperatures are associated with slow deposition rates and an amorphous fiber structure which has a high initial strength but loses the strength due to crystallization at high temperatures. A variety of silane gases including methyltrichlorosilane (CH3SiCl3), methyldichlorosilane (Cl2CH3SiH), ethyltrichlorosilane (C2H3SiCl3), or tetrachlorosilane (SiCl4) have been used as reactant gases for the deposition of SiC. Methyl-

69

dichlorosilane provides the fastest production rate and requires a lower deposition temperature; the fibers produced, however, have a low tensile strength due to the rough fiber surface (DeBolt and Krukonis, 1973). Methyltrichlorosilane or the more widely available ethyltrichlorosilane, on the other hand, yield fibers with high strength and smooth surface but at prohibitively low growth rates. A compromise between good mechanical properties and high deposition rates is achieved by using a mixture of silane gases. Present CVD-SiC fibers utilize a 3 3 37 ^m carbon substrate which has been given a pyrolytic carbon coating about 1.5 jim thick. The pyrolytic carbon layer provides good thermal conductivity and serves as a buffer to reduce processing stresses which result from the thermal expansion mismatch between the carbon core and the SiC mantle (Nutt and Wawner, 1985). At the end of SiC deposition, the fiber surface is given a carbonaceous coating by varying the processing conditions and composition of the reactant gases. Carbon coating results in a significant increase in fiber strength due to the healing of surface flaws (Cornie et al., 1981). It also provides the weakly bonded interface critical to the toughness of ceramic matrix composites. However, carbon suffers from rapid oxidation. Furthermore, it is not suitable for many metal matrices where chemical reactions may take place at the fiber/metal interface. For these reasons, the carbon coating often incorporates, or is overcoated with, SiC. A family of SiC monofilaments are currently produced by Textron Corporation under the designation "SCS". These filaments have a C-SiC coating. Several SCS fibers with different coating thicknesses and C/Si ratios in the coating are available. SCS-2 and SCS-8 fibers were developed primarily

70

2 Fibers and Whiskers

for use in aluminum matrices while SCS-6 fibers with a thicker coating were offered for the more reactive titanium matrices (DeBolt et al., 1982). These fibers have a diameter of about 140jnm which is too large for toughening of ceramic matrices but provides a convenient model system for the interfacial studies in SiC reinforced CMCs. Textron Corporation has recently introduced a new fiber, SCS-9, with a diameter of only about 70 jam for use primarily in ceramic matrix composites. Depending on the processing conditions, fibers with a wide range of compositions from stoichiometric SiC to SiC with free C or Si can be obtained. Free carbon improves resistance to self-abrasion. SiC has two polymorphs: a-SiC with a hexagonal wurtzite structure and P-SiC with a cubic zinc blende structure. SiC fibers produced by CVD are primarily of P-SiC type with a high density of stacking faults and microtwins. Wawner and Nutt (Nutt and Wawner, 1985; Wawner et al., 1983) studied the microstructure of CVD-derived SiC fibers by transmission electron microscopy and found that columnar grains of p-SiC grow in the direction of deposition with their closed-packed {111} planes perpendicular to the direction of deposition. An interesting feature of the structure of these fibers is a distinct transition in the grain structure occurring at a diameter of ~80 |im as shown in Fig. 2-31. Columnar P-SiC crystallites in the inner SiC ring are about 40-50 nm in diameter while those in the outer sheath are ~90-100 nm. This change in the microstructure has been attributed to the design of the reaction chamber and the method of processing. The most widely used member of the SCS series of fibers is SCS-6 fiber. A detailed study of the microstructure and composition of SCS-6 fiber has been carried out by Ning, Pirouz and co-workers (1990,

y^-^M:

Figure 2-31. The transition in the grain structure of CVD-fabricated SiC fiber (from Wawner, Teng, and Nutt, 1983).

1991) using transmission electron microscopy and scanning Auger microscopy. They found the fiber to consist of many concentric layers as shown in Fig. 2-32. Starting from the fiber center, the core (carbon substrate) structure consists of randomly oriented blocks of turbostatic carbon (TC) about 1-5 nm in size. The next layer is the pyrolytic carbon coating (~1.5 j^m thick) which also has a turbostratic structure. However, the TC blocks in this layer are in the range of 30-50 nm and are preferentially oriented with their c-axis parallel to the radial direction. The SiC mantle with a total thickness of ~ 50 |im consists of four distinct layers. The transition from one layer to the next is marked by a sharp change in the grain size and degree of preferential orientation. The sharpest transition occurs between SiC-3 and SiC-4 layers. The first three layers, i.e.,

2.8 Silicon Carbide and Silicon Carbide-Based Fibers

71

outermost coating total thickness 3 jam

inner coating 1.5 (im

Figure 2-32. The ring structure of SCS-6 SiC fiber (from Ning and Pirous, 1991).

SiC-1, SiC-2, and SiC-3 layers, consist of elongated (3-SiC crystallites with the long axis parallel to the crystal directions. Moving away from the center, the crystallites become increasingly larger and aligned with the radial direction. The crystallite length is — 5-15nm immediately adjacent to the pyrolytic carbon layer and of the order of microns within the SiC-3 layer. The enlargening of the grains is accompanied by increased faulting on the {111} planes. The transition from the SiC-3 region to the SiC-4 region is marked by an almost doubling in the size of the elongated p-SiC grains, which are now heavily faulted, and a sharp increase in the alignment of the directions of SiC parallel to the radial direction. A significant finding is that the SiC-4 layer is almost perfectly stoichiometric silicon carbide while the SiC-1, SiC-2 and SiC-3 regions have 10-20% excess carbon. The presence of boundaries between the first three layers of SiC is attributed to a step-wise decrease in the Si/C ratio outwards from the fiber center. The outermost carbonaceous coating on the SCS-6 fibers exhibits three distinct sublayers (Ning and Pirouz, 1991; Ning et al, 1990; Pirouz et al, 1989): the outermost sublayer is carbon-rich, containing SiC crystallites that grow in density and decrease in size in a radial direction inwards

from the fiber surface; the thin middle sublayer is silicon-free carbon; and the third, innermost sublayer is very similar in structure to the first layer. In all sublayers, carbon has a turbostratic structure with the weakly bonded graphitic planes parallel to the fiber surface. The debonding of the turbostratic graphite layers in the coating provides the weak interface desirable in ceramic matrix composites. Debonding occurs in the middle, silicon-free sublayer or at its interface with the innermost sublayer since SiC crystallites in the other two sublayer strengthen the carbon (Ning et al, 1990). The occurrence of debonding in or near the middle sublayer is desirable since debonding between the SiC sheath and the innermost coating sublayer would result in the complete removal of the coating and, therefore, a loss in fiber strength. During exposure to oxygen at high temperatures, SiC crystallites in the outer coating sublayer oxidize to SiO 2 , thus providing some protection against further oxidation of the carbon coating. DiCarlo (1989) noted that for short times of at least 30 min, the carbon content is not significantly affected by exposure to 100% oxygen up to 1200 °C. He further noted that the carbon layer is virtually unaffected by argon treatment up to 1400 °C. However, a pretreatment with argon above 900 °C followed by a 100% oxygen treatment results in accel-

72

2 Fibers and Whiskers

erated carbon attack compared to direct oxygen treatment. It has been suggested that argon pretreatment above 900 °C causes the outer coating sublayer to become enriched in carbon, thus eliminating the protective effect of SiC particles.

Dichlorodimethvlsilane Dechlorination with Na (to NaCI)

2.8.3 Polymer-Derived SiC-Based Fibers In 1975, Yajima and co-workers introduced a new process for making SiC-based fibers from a spun polymer precursor (Yajima et al., 1976 a, b, 1978; Hasegawa et al., 1980; Hasegawa and Okamura, 1983; Yajima, 1985). Following some improvement of the initial process, polymer-derived SiC multifilaments were produced on a pilot scale in 1978. In 1981, Nippon Carbon Company initiated the commercial production of these fibers in continuous form of 500 filaments/yarn under the tradename "Nicalon" The main stages in the preparation of Nicalon SiC fiber are shown in Fig. 2-33. Dimethylpolysilane (DMPS) is obtained by dechlorination of the starting material, dichloromethylsilane, with molten sodium at about 130°C under a nitrogen atmosphere. DMPS is polymerized to polycarbosilane (PCS) by heating at 450-470 °C in an autoclave under argon atmosphere. The molecular weight of PCS rises markedly with treatment temperature within the range of 450 to 470 °C. The molecular weight of PCS has a significant effect on the properties of the resultant SiC fibers. At present, an average molecular weight of about 1500 is used. Following a vacuum distillation treatment at up to 280 °C to remove the undesirable low-molecularweight PCS components, PCS is melt spun in the presence of N 2 gas at about 350 °C to form continuous polycarbosilane fibers. To improve the melt-spinnability of the PCS polymer, thus reducing the fiber di-

Dimethylpolysilane Polymerization at 470°C in autoclave (in argon)

•j-Si—(p4 Polycarbosilane Melt spinning at 350°C (in nitrogen) Polyca rbosilane fibers

Curing at 190°C in air or RT in ozone Polycarbo silane fibers made infu sible through cross-linki ng by oxygen Pyrolysis-heating to 1300°C or less in hydrogen or vacuum (100°C/hr) Si C fiber with arnorphous or microc rystalline p-SiC structure

Figure 2-33. The steps in the processing of Nicalon SiC (from Andersson and Warren, 1984).

ameter, a few percent of phenylchlorosilane is added to the starting material. This promotes a chain-like molecular structure in polycarbosilane which results in almost 30% reduction in fiber diameter. The spun PCS fibers are then cured in air for about 30 min at a temperature up to 200 °C. This induces cross-linking of PCS polymer chains by oxygen which is necessary to prevent melting of fibers during the subsequent high temperature heat treat-

73

2.8 Silicon Carbide and Silicon Carbide-Based Fibers

ment. The cured fibers are finally converted to SiC by heating at 1200-1300 °C in N 2 gas or vacuum. The pyrolysis temperature has a profound effect on the microstructure and properties of the SiC fiber (Yajima, 1985). Maximum tensile strength is achieved around 1200-1250 °C, while the Young's modulus remains virtually independent of the heat treatment temperature as shown in Figs. 2-34 and 2-35, respectively. X-ray powder diffraction patterns of SiC fibers pyrolyzed at temperatures from 900 to 1500°C indicate that fibers have an amorphous structure up to 1000 °C (Fig. 2-36). Beyond 1000 °C, crystallization of SiC to P-SiC takes place. Crystallization and grain growth become

900° 1000°

a "E & fi

1100°

5

1200°

I

1300°

1500° 20

30

40

50

60

70

80

26

Figure 2-36. X-ray diffraction patterns of Nicalon SiC fibers pyrolyzed at different temperatures (from Yajima, 1985).

400

600

800

1000

1200

1400

1600

1800

Temperature (°C)

Figure 2-34. The variation of the strength of Nicalon SiC fiber with the pyrolysis temperature (from Yajima, 1985). 250

400

600

800

1000

1200

1400

1600

1800

Temperature (°C)

Figure 2-35. The variation of the Young's modulus of Nicalon SiC fiber with the pyrolysis temperature (from Yajima, 1985).

significant beyond 1300 °C and are accompanied by a drop in fiber strength. Nicalon fibers thus produced have a non-stoichiometric composition consisting of SiC, SiO2 and free C. The presence of SiO2 accounts for the lower Young's modulus of Nicalon compared to the CVDderived SiC fibers. The oxygen is introduced mainly during the curing treatment. The initial laboratory grade fibers had a composition of 64% (by weight) SiC, 15% C and 21% SiO2 (Yajima et al., 1979). Nippon Carbon Company has since produced several generations of Nicalon fibers with progressively lower oxygen levels and slightly varying microstructures. The family of Nicalon fibers currently produced commercially includes the ceramic grade (CG) for high temperature mechanical applications, the high volume resistivity (HVR) grade (low dielectric constant fiber), and the low volume resistivity (LVR) grade (higher conductivity fiber). A low-oxygen-

74

2 Fibers and Whiskers

content Nicalon fiber with improved thermal stability has been developed recently but is not yet commercially available (Okamura et al., 1988, 1989); this fiber is discussed in Sec. 2.8.4. The nominal compositions of several generations of the ceramic grade Nicalon fiber and the experimental low oxygen fiber are shown in Table 2-13. The standard grade Nicalon fibers produced earlier had a virtually amorphous structure while the ceramic grade fibers consist of very fine crystallites of f$-SiC, about 1.7-2.5 nm in size, in an amorphous matrix (Simon and Bunsell, 1983, 1984; Okamura, 1987; Sawyer et al., 1985, 1986). The continuous amorphous phase is a silicon oxycarbide, SiCaOb (a + b = 4) and controls the modulus and fracture properties of the fiber (Lipowitz, 1991; Sawyer etal., 1987). The P-SiC crystallite size is dependent not only on the heat treatment temperature but also on the average molecular weight of the polycarbosilane Precursor (Hasegawa and Okamura, 1983; Hasegawa, 1989,1990). Stronger and stiffer fibers are obtained by using larger molecular weight PCS. Laffon etal. (1989), and Maniette and Oberlin (1989) conducted TEM studies of the fiber microstructure which have indicated that the excess carbon is in the form of nanocrystallites with a turbostratic structure. Using X-ray scattering, Lipowitz et al. (1990) found fine porosity in the range of ~ 5-20 vol.% in polymer-derived S i - C - O (Nicalon) and S i - C - N - O fibers. The pores were on the Table 2-13. Compositions of Nicalon fibers3 (Takeda etal., 1991). Fiber Ceramic grade Low oxygen a

Atomic ratio.

Si

C

O

H

1.00 1.00

1.29 1.31

0.38 0.02

0.10 0.05

order of a few nanometers in diameter. The average pore size increased and volume fraction of porosity decreased with increasing pyrolysis temperature. Porosity leads to reduced density and elastic modulus compared to a fully dense fiber. Scanning Auger depth profiling of the fiber surface has revealed that fibers possess a carbonrich surface layer of ~ 60 nm thickness (Lipowitz, 1991). The diameter of Nicalon fibers varies in the range of 10 to 20 |im with most fibers having a diameter of ~ 15 jam. Yajima (1985) showed that both fiber strength and Young's modulus decrease rapidly with increasing fiber diameter. The strength averages around 2.9 GPa, the Young's modulus around 190 GPa and density around 2.55 g/cm3. The density is lower than that of stoichiometric (3-SiC (3.16 g/cm3) due to the presence of SiO2 and C in the structure. Fiber strength varies markedly with the gage length due to flaw distribution. Testing both pilot and commercial fibers with a gage length in the range of 2-300 jim, Andersson and Warren (1984) found that fiber strength could, to a good approximation, be given by the following relationship fu

= 8.12-(lnL/7.8)

(2-17)

where L is the gage length. A similar relationship was also found by Simon and Bunsell (1982) for commercial fibers but at somewhat lower strength levels. Nicalon fibers exhibit significant thermal degradation at temperatures above ~1200°C. The thermal stability and strength loss of Nicalon fiber in various environments have been investigated by many researchers (Mah et al., 1984; Clark et al., 1985, 1986, 1987; Pysher et al., 1989; Bibbo et al., 1991). The loss of strength is more severe in non-oxidizing environments such as argon, nitrogen or vacuum where crystallization and grain growth of

2.8 Silicon Carbide and Silicon Carbide-Based Fibers

p-SiC as well as the development of a porous structure have been reported (Clark etal, 1985, 1986, 1987). Thermodynamic instability of S i - C - O system in nonoxidizing environments at temperatures above 1200 °C is due to the following carbothermic reduction reactions (Lipowitz, 1991; Luthra, 1986). 3C + SiO2 C + SiO2

SiO + CO

2C + SiO

SiC

(2-18)

The oxygen in the fiber reacts with excess carbon causing the evolution of CO and some SiO gases. This intrinsic oxidation process is accompanied by a weight loss and the development of porosity in the fiber. It has been suggested that free carbon particles in the fiber promote its thermomechanical stability by hindering creep and grain growth (Catoire etal., 1987). Once the carbon is removed, grain growth of the P-SiC crystallites occurs as observed in elevated temperature aging of fibers in argon 1200°C (Bibbo etal, 1991). Thermodynamic calculations by Luthra (1986) and experimental studies by Mah et al. (1984) have indicated that CO gas is the predominant decomposition product of the intrinsic oxidation of Nicalon fiber at temperatures above 1200°C. Bibbo etal. (1991) studied heat treated ceramic grade Nicalon fibers at 1300°C under various partial pressures of CO gas and noted that fibers retained 75% of their strength when heat treated in CO as compared to only 25% strength retention in argon. The loss in strength is much less severe and fibers exhibit a weight gain when aging takes place in an oxidizing environment such as air. This has been attributed to the formation of a stable SiO2 skin which serves as a diffusion barrier for SiO and CO gases,

75

thus slowing down the oxidation and consumption of free carbon (Mah et al, 1984; Clark et al, 1985, 1986, 1987). Mah et al. (1984) noted that strength retention was higher for a 1300-1400 °C aging than a 1200 °C aging in air for 2 h. This was linked to a better formation of the SiO2 shell at the higher temperature range. Clark et al. (1987) studied the oxidation behavior of Nicalon fibers after 12 h exposure to 1400 °C under a O 2 /He atmosphere with varying oxygen levels. They recorded a weight loss accompanied by P-SiC grain growth and a significant drop in fiber strength when heat treatment was carried out at a low oxygen level of 8 ppm. This was indicative of active oxidation of the fiber. At oxygen levels greater than 100 ppm, however, fibers exhibited a weight gain, due to the protective SiO2 layer formation, and a less severe strength reduction. The thickness of the silica layer increased with the oxygen level. The results indicated that there exists an optimum SiO2 thickness below which diffusion control is insufficient and above which residual stresses caused by the thermal expansion mismatch between the fiber and silica become large enough to bring about cracking of the SiO2 shell. The benefit of higher temperatures in producing a thicker, more protective silica coating is somewhat offset by the crystallization of silica to cristobalite at temperatures above 1400°C (Clark etal, 1985, 1986, 1987). Upong cooling to room temperature cristobalite undergoes a phase transformation to ot-cristobalite at about 250 °C. This phase change is accompanied by a volume change which leads to the cracking of SiO2 shell. The need for a fiber with better thermal stability than Nicalon resulted in the introduction of a S i - T i - C - O fiber by Ube Industries, Ltd. under the tradename "Tyranno" (Yamamura etal, 1988). The

76

2 Fibers and Whiskers

synthesis and pyrolysis of polytitanocarbosilane) have been described by Yajima et al. (1981) and by Okamura (1987). Tyranno fibers are produced by the melt spinning and pyrolysis of polytitanocarbosilane in a process analogous to the production of Nicalon fiber. Polytitanocarbosilane is formed by the cross-linking of polycarbosilane with a titanium compound such as titanium tetraisopropoxide and is easily spun into fiber form. The addition of 1.5 to 4.0 wt.% Ti was made to improve the thermal stability of the fiber. Tyranno fiber is reported to have a predominantly amorphous structure up to 1300°C (Yamamura et al., 1988). It has a strain to failure of about 1.4 to 1.7% which is considerably greater than that of Nicalon fiber. Thermal stability studies of the Tyranno fiber in a glass-ceramic matrix, however, have shown it to be less stable and more reactive than Nicalon fiber (Brennan, 1991). The reason for this was determined to be the significantly higher oxygen level of the Tyranno fiber (~20at.%). Fischbach et al. (1988) investigated the microstructure and mechanical properties of Tyranno fibers at room temperature in the as-fabricated condition and after heat treatment at 1300°C and 1350 °C under a nitrogen atmosphere. They reported significant reductions in both strength and modulus of fibers after heat treatment, accompanied by an almost doubling of the mean crystallite size. 2.8.4 Recent Developments in SiC-Based Fibers

The inability of both Nicalon and CVDderived SiC fibers to withstand continuous exposure to the high temperatures required of ceramic composites applications has sent off a wave of research activities in pursuit of a high temperature SiC fiber. None of these efforts, however, have yet

produced a new commercial fiber. Among these efforts is the recent development of a low oxygen content Nicalon fiber (the HiNicalon fiber). It was described above that the thermal instability of the commercial Nicalon fiber is due to its high oxygen concentration of about 10 wt.% which is introduced during the oxidative cross linking step of the fiber production (Clark, 1985). By replacing the oxidation curing with electron beam radiation curing, Okamura etal. (1988, 1989) have succeeded in preparing fibers with a low oxygen content of about 0.4%. The low oxygen fiber reportedly has a high strength of 2.8 GPa and an elastic modulus of 270 GPa which is considerably higher than that of the ceramic grade fiber (Takeda etal., 1992). Takeda etal. (1991, 1992) compared the thermal stability of the low oxygen content fibers with the oxidation-cured ceramic grade fibers. Fibers were exposed to temperatures in the range of 1500 to 2000 °C in an argon atmosphere for periods of 1 h and 10 h. In contrast to the dramatic compositional changes and severe crystallization of the ceramic grade fiber, the low oxygen fiber virtually retained its chemical composition and exhibited only a ten-fold increase in its crystallite size up to 2000 °C. The beneficial effect of a low oxygen content has also been demonstrated by Toreki et al. (1992) who prepared continuous silicon carbide fibers with an oxygen content of less than 2 wt.% by dry spinning of polycarbosilane solution and subsequent pyrolysis of the polymer fiber. Ube Industries Ltd. has also developed a new S i - T i - C - O fiber with lower oxygen content (~13 wt.%) and therefore more refractoriness than Tyranno fiber. The new fiber is marketed under the tradename Lox-M. It has a diameter in the range of 8.5-11 jam and is available in various tow sizes from 400 to 1600 filaments per tow.

2.9 Silicon Nitride-Based Fibers

The microstructure of the fiber borders on the amorphous, consisting of extremely fine-grained (3-SiC (1600°C) of meltspun organosilicon polymers. The fiber is over 95% (3-SiC with a crystallite size of 30 to 40 nm. Also among the new developments is the fabrication of a sintered oc-SiC fiber by Frechette and co-workers (1991) at Carborundum Company. The fiber is produced in a continuous process by sintering spun filaments of submicron SiC power in a polymeric binder. Fibers with diameter in the range of 25-100 jim and an average strength of about 1 GPa have been produced. oc-SiC fibers have shown almost no loss of strength after 4 h of exposure to a temperature of 1550°C in air.

2.9 Silicon Nitride-Based Fibers The polymer precursor route offers considerable flexibility in terms of the fiber compositions and microstructures that can be produced (Lipowitz, 1991; Okamura, 1987; Okamura et al., 1988). In addition to silicon carbide fibers described above, continuous silicon nitride (Okamura et al., 1987), silicon carbonitride (Legrow et al., 1987; Penn et al., 1982), and silicon oxynitride (Okamura, 1987) fibers have been prepared by the synthesis and pyrolysis of organosilicon polymers. However, at the time of this writing, none of these fibers have been commercialized. Silicon carbonitride fibers with a typical composition of 60.0 wt.% Si, 32.6 wt.% N, 2.3 wt.% C, and 2.2 wt.% O are produced

71

by Dow Corning Corp. from hydridopolysilazane (HPZ) (Lipowitz, 1991; Legrow et al., 1987). HPZ polymer is synthesized by the reaction of trichlorosilane with hexamethyldisilazane and is melt-spun into fiber form. The preceramic fiber is cured in an oxygen-free argon gas saturated with trichlorosilane vapor. The cured fibers with chlorosilane cross linking are then pyrolyzed at temperatures up to 1200 °C under flowing high purity nitrogen to silicon carbonitride. Silicon carbonitride fibers thus produced have a diameter ranging from 10-12 Jim, a tensile strength of ~ 2.8 GPa, and an elastic modulus of ~ 180 GPa. The fiber microstructure consists of amorphous silicon carbonitride with about 5 wt.% nanocrystalline carbon with a turbostratic structure. The fiber is not fully dense, but contains about 20 vol.% fine porosity. Fibers show a surface layer of ~0.5 nm thickness that is richer in oxygen and poorer in nitrogen compared to the uniform S i - C - N - O composition of the interior. It has been reported that HPZ fibers tested in air at temperatures up to 1400 °C behaved similarly to Nicalon and Tyranno fibers in that strength and modulus dropped sharply above about 1200 °C (Lipowitz, 1991). Okamura et al. (1987) have prepared continuous silicon carbide (Si3N4) and silicon oxynitride (Si-N-O) fibers from polycarbosilanes. The melt-spun polycarbosilane fibers were given an electron beam radiation curing treatment for the preparation of silicon nitride fibers and an oxygen curing treatment for silicon oxynitride fibers. The cross-linked preceramic fibers were then heat treated in NH 3 gas at temperatures up to 1400 °C to obtain silicon nitride and silicon oxynitride fibers with a diameter of 11-13 jam. X-ray diffraction patterns of fibers nitrided at different temperatures revealed that both fibers were

78

2 Fibers and Whiskers

amorphous up to 1300°C, but crystallized to oc-Si3N4 at higher temperatures. For the silicon nitride fiber, the strength and modulus reached their maxima, ~ 1.0-1.5 GPa for strength and -100-130 GPa for modulus, at 1300°C, beyond which temperature they both dropped sharply with strength reaching almost zero at 1400 °C due to the crystallization of oc-Si3N4. The strength and modulus of the silicon oxynitride fiber depend on the level of oxygen introduced during the curing process. Fibers cured at 160 °C (O/Si = 0.36) showed the highest strength, modulus, and thermal stability. These fibers had a higher strength and modulus than the silicon nitride fibers.

2.10 Whiskers 2.10.1 Introduction

Whiskers are single crystal reinforcements with an aspect ratio of usually over 20. The high degree of structural perfection of whiskers which stems from their single crystalline morphology and small diameter (usually submicron) bestows on them a very high Young's modulus and a strength that approaches the theoretical strength of the material calculated by Orowan (1949) to be (2-19) where E is the Young's modulus, y is the surface energy per unit area, and a0 is the equilibrium spacing of the atomic planes. Carbon whiskers have the highest specific strength and specific Young's modulus among all existing reinforcements. The absence of crystalline defects is also responsible for the better creep resistance and, therefore, higher temperature capability of the whiskers compared to their fiber counterparts. A variety of ceramic whiskers are

currently produced for use primarily in ceramic and metal matrices. Some of these whiskers and their properties are shown in Table 2-14. The addition of whiskers to a monolithic ceramic raises the toughness through a variety of mechanisms which include crack deflection, whiskers bridging, and whisker pullout. Three to four-fold increases in the toughness of alumina have been reported by the incorporation of SiC whiskers (Wei and Becher, 1985; Tiegs and Becher, 1987). Although whiskers are usually quite short, they possess large aspect ratios, usually over 20, due to their small diameters and are thus efficient reinforcements. Whisker diameter is also a critical parameter in terms of its strength; increasing the diameter raises the probability of the presence of defects, hence lowering the strength. Being single crystals, most whiskers have a non-circular cross section. The shape of the cross section depends on the crystal structure and the growth orientation of the whisker. Thus, it can be triangular, hexagonal, rhombohedral, or another shape. For a non-circular cross section, whisker diameter is defined as the square root of the cross-sectional surface area. The shape of the whisker affects the total whisker/matrix interfacial area and hence the composite properties. Whiskers with submicron diameters may pose a potential health hazard similar to that of asbestos (Bogoroch and Luck, 1988). Asbestos fibers with a diameter less than 1 jim and aspect ratio greater than five are known to cause irritation of the eyes, skin, and pulmonary system and may even lead to pulmonary fibrosis and cancer. Precautions must, therefore, be taken during handling of dry whiskers or machining of whisker reinforced composites.

79

2.10 Whiskers

Table 2-14. Whiskers and their properties. Material

Designation

Phase

Manuf.

Diameter Density

(mm) SiC SiC

Silar SC-9

sew

a 95% p-5% a

SiC

Tokama

P

SiC SiC

VLS Versite

P

Si 3 N 4

SNW

>97%a

Si3N4 C

Versite —

— —

SiO2

XPVI

10-6oC-l

Young's Tensile Length modulus strength )

(GPa)

(GPa)

(mm)

3.2 3.21

4.7 4.4-4.7

689 483

6.89 20.68

10-80 10-40

3.19

—

400-700

3-14

50-200

— —

— —

578 -560

8.40 -11

0.2-0.5

—

—

—

—

-5,000 up to 3,000 50-300

— >l) value.

,

0

0.4

0.8

1.2 1.6 2.0 2.4 Shear s t r a i n , %

2.8

3.2

Figure 3-17. Shear stress-strain curve obtained from a tension test of a [±45] 2S graphite/epoxy specimen. p.

2

Load span - L

Support span

Note that as an alternative, the lamina shear modulus may be directly determined from the relation G12 = ^4Ex

100-

Figure 3-18. Four-point loading.

(3-15)

where Ex and vxy are the longitudinal modulus and major Poisson's ratio, respectively, of the {±45} 2S laminate. An example of shear response is shown in Fig. 3-17. 3.6.4

Flexure Test Methods Support span

The four-point flexure test is shown in Fig. 3-18 and the three-point flexure test is shown in Fig. 3-19. The quarter-point loading has been utilized primarily for high modulus materials such as graphite/epoxy and boron/epoxy composites. These test methods are used for determining flexural strength and modulus. The test is not recommended for generating design data. However, it does provide a simple test for

Figure 3-19. Three-point loading.

quality control. Flexure tests for high modulus, continuous filament composites are generally limited to unidirectional materials with fibers oriented at either 0° or 90° to the beam axis. Specific requirements for radius of load noses and supports are given in ASTM

3.6 Test Methods

D790-71. Span-to-depth ratios, L/d, depend on the ratio of tensile strength parallel to the beam axis to interlaminar shear strength. For strength ratios less than 8 to 1, an L/d ratio of 16 is recommended. This is typical of fiberglass composites. For high modulus materials such as graphite/epoxy and boron/epoxy a value of L/d = 32 is recommended for 0° unidirectional composites. For 90° unidirectional properties of these materials, L/d = 16 is appropriate. Recommended specimen dimensions can also be found in ASTM D790-71. For tensile strength determination the specimen is loaded until failure occurs. The maximum tensile stress, Sm, is determined for the three-point loading from the relationship

3PL 2bd2

3PL

four-point loading with a load-span of one-half of support span, 0.17L3

~~b¥

(3-17)

It should be noted that Eqs. (3-16) and (3-17) are valid only for material in which the stress-strain curve is linear to failure. If some nonlinear stress-strain behavior occurs, an error will be introduced into the relationships for Sm. An adjustment in the calculation for Sm must also be made if the deflections become large (ASTM D790-71). Modulus is determined for the three point flexure test from the relationship

3,6.5 Mode I Interlaminar Fracture Test Mode I interlaminar tests are conducted with a double cantilever beam (DCB) specimen, Fig. 3-20. The compliance, C, of the DCB specimen may be obtained from elastic beam theory as C =

2a3 3E/

(3-20)

where a is the crack length and El is the flexural rigidity of the beam. The strain energy release rate, Gl9 is given by: 2w da

P2a2 wEI

(3-18)

where: EB is the modulus of elasticity in bending in N/m 2 , and r1 is the slope of the tangent to the initial straight-line portion of the load-deflection curve in N/m. For

(3-21)

where P is the applied load and w is the width of the specimen. Equations (3-20) and (3-21) give

r3

E * B ~4bd3

(3-19)

(3-16)

where b is the beam width. For four-point loading at quarter points

Td2'

113

Figure 3-20. DCB specimen.

(3-22)

114

3 Polymer Matrix Composites

Critical conditions occur when GY reaches its critical value GIC for P = Pc; i.e., Pc2a2 G ir = wEI

(3-23)

There are several ways to reduce the data from interlaminar fracture tests (Carlsson and Pipes, 1987).

3,7 Damage Tolerance Damage tolerance for aerospace composites is defined as the capability of the composite structure to sustain a level of low-velocity impact with barely visible damage and retain appropriate residual strength. The low-velocity impact has become a major issue due to a concern or damage resulting from drop of tools during maintenance of composite material structures. The main problem is the effect of accidental lateral impact loading on the extent of damage, which is often located in the internal part of the laminate, nonvisible on its external surfaces. The main goal for damage tolerance evaluation and design is to develop empirical and analytical methods aiming to predict the deterioration in structural performance due to such damage. The tests on basic material properties discussed in the preceding section are essential for generating data on design allowables. However, it is also needed to determine the damage tolerance capability of the composite material. Because of the weak interlaminar fracture resistance of polymeric composites, damage tolerance evaluation or tests primarily concern delamination and low-velocity impact resistance of the composite. Graphite/epoxy composites have been reliably used in many secondary aerospace structures. In military aircraft, this class of

materials also has been used in primary structures such as wings and tails. Conventional graphite/epoxy composites which are brittle have the limitation that maximum weight savings cannot be achieved due to their ease of susceptibility to damage. The primary impediment to full utilization of continuous fiber polymeric composites is their inherent tendency to delaminate. Delamination (Fig. 3-21) is the most prevalent type of life-limiting failure in advanced composites. Technologically, it is also one of the most significant problems in advanced composites. There are various sources of delamination. Delamination may develop during manufacturing with improper consolidation of plies. They may result from impact damage or from three-dimensional interlaminar stresses that develop at stress-free edges or discontinuities such as the free surface of a hole. Both interlaminar normal (oz) and shear stresses (TXZ and xyz) as shown in Fig. 3-22 can be detrimental. The interlaminar stresses developing at discontinuities in typical composite structures which may promote delamination are presented in Fig. 3-23. Cut fibers

Matrix cracks

Delamination

Figure 3-21. Delamination and other local damage mechanisms in a composite laminate.

3.7 Damage Tolerance

115

Interface Free-edge

Figure 3-22. Interlaminar normal and shear stresses (circled) in a composite laminate.

Free edge

Notch (hole)

Bond joint

Bolted joint

Figure 3-23. Interlaminar stresses arise from typical discontinuities in composite structures.

Three major failure modes may occur during loading of composite materials namely: fiber fracture, interfiber transverse matrix cracking, and interlaminar fracture or delamination. Fiber fracture rarely occurs in multidirectional laminates under working loads. Presence of a hole may reduce residual strength. Interfiber matrix cracking, which occurs at a low tensile stress transverse to fiber direction, does not significantly reduce composite performance.

Interlaminar stresses can also develop due to compressive loading in laminates. A consequence of compressive loading can be local and global buckling of plies in a laminate. This situation is enhanced due to the propensity of delamination crack growth and separation of plies. An area of considerable concern for conventional graphite/epoxy composites is their response under low velocity impact. The damage may not be visible. However, internal delamination can be quite large

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resulting in the loss of compression strength and structural integrity. The compressive residual strength of the composite structure may be controlled by the size and location of the delamination. Hence, the design of the composite structure should be made resistant to both formation and growth of delaminations. Since it is impossible to control the low energy impacts that can initiate delaminations, composite structural designer must assume that such delaminations exist in the laminate a priori and control their growth through intelligent design. The current damage tolerance philosophy for composite structures quite justifiably leans heavily on impact damage. It is anticipated that composite structures should sustain lifetime design loads in a damaged condition up to the condition where a barely visible impact damage is detectable. Due to these requirements, the design allowable ultimate strain is limited to ~0.4%. However, for further weight savings and for greater damage tolerance performance, design strains of 0.6% are considered highly desirable. Traditional graphite/epoxy composites are brittle and show poor resistance against delamination growth and interlaminar separation resulting from low velocity impact. Although toughened graphite/epoxy composites show good properties, particularly delamination resistance, their performance under hot-wet conditions is of concern. Interleaving of graphite/epoxy composites have shown much promise. In recent years, thermoplastic composites, particularly graphite/PEEK composite, have shown considerable promise in delamination fracture resistance. Thermoplastic matrices give an order of magnitude increase in interlaminar toughness compared to existing epoxy resin composites (GIc of 1500 J/m2 vs. 150 J/m2). Dis-

cussion on toughened polymeric composites including thermoplastic composites can be found in two ASTM publications edited by Johnston (1987) and Newaz (1989), respectively. Fatigue loading is common in aerospace structures. Delamination initiation and growth due to fatigue is also of concern. Delamination growth under fatigue loading in both thermoset and thermoplastic composites has been studied well except that the knowledge base for thermoplastic composites is not as extensive. It is now clear that delamination crack growth rate under fatigue loading is relatively fast in both graphite/epoxy and thermoplastic composites as exhibited by high n values 8-20 in the Paris-type equation. An implication is that fatigue propagation life is short. Thus, an approach to proper damage tolerance design calls for design on the basis of threshold stress-intensity factor; i.e., to avoid any possible growth of the small delamination crack due to fatigue loading. For anisotropic layered material systems, the complexity in actual structures arises from the inherent coupling of Mode I, II and III fracture and possible geometric and material discontinuities. Furthermore, one can envision possible nonlinearities - making the problem of delamination quite complex. The source of nonlinearity may be either due to large deflection or due to ductility of the material as in thermoplastic composites. Thus, the challenge rests on both materials and structural design to avoid this failure mode in aircraft structures. The most severe and frequent event during service and maintenance is the local transverse impact due to foreign objects. The evaluation of their distribution of stress waves due to impact at function of time and space is highly complex due to the

3.7 Damage Tolerance

heterogeneous and anisotropic nature of the composite material and is complicated by loading rate. Hence, the loading rate is an important variable to consider. Another important factor which affects impact loading consequences is the boundary conditions of the loaded part. In real situations such conditions may vary from rigidly fixed to free simply supported and their exact translation to experimental and analytical models are difficult to determine. Hence one way to tackle such variable is by comparing the extreme conditions and establish some empirical factors to adjust the laboratory finding to the real life structural behavior. On the other hand, delamination is considered to be highly critical to the durability of the composite structure under subsequent long term (or cyclic) environmentloading service conditions. Even transverse cracking may be crucial under certain loading conditions as a source for additional delaminations growth initiating from the transverse crack root. Various structural threats vs. local damage are illustrated in Fig. 3-24. Under flexural loading two types of delaminations may develop, namely:

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(a) Edge delaminations at the tensile zone close to the outer surface, where tensile stresses attain their maximum value. (b) Shear delamination, close to the neutral plane where interlaminar shear stresses are predominant under flexure. The chance of predominance of one failure mode above the other depends on span to thickness ratio of the beam or panel considered and on boundary conditions. It is also dependent on loading rate as edge delamination propagates slowly as a function of tensile axial strain, while shear delamination propagates both fast and abruptly when loading energy attains its critical level. Hence shear delamination mode is expected to be predominant under impact loading whereas both modes may develop under quasistatic loading. 3.7.1 Damage Characterization and Detection During fabrication, interlaminar flaws may be formed. During service and maintenance period, damage may be developed. In most cases these flaws or damages may be described geometrically approximately as two-dimensional circular

Resulting local damage Cut fibers Material flaw

i iT

Matrix cracks

i i

Nature of threat

Few MFG error Impact Puncture Hydrodynamic ram

Many

i T Very large number

i

Delamination

Minimal

Characteristic spacing assumed

t

Moderate

i

i i

Extensive

Figure 3-24. Various structural threats cause different amounts of local damage.

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interlaminar separations. In thin sections, or laboratory specimens, such cracks may propagate through the element width and form a "one-dimensional" delamination. The damage (flaw) dimension and its location through the laminate thickness are functions of loading mode and direction, and laminate lay-up. Thus, damage due to processing and due to in-service mechanical loads is a crucial factor which determines the subsequent performance of the laminate under working loads. Hence the importance of reliable and accurate detection methodology for damage (flaw) as well as its geometrical characterization, and correct interpretation cannot be overstated. To date the best method for detection of interlaminar separation pattern is the ultrasonic C-scan technique. With a modern ultrasonic system it is possible to obtain both planar dimensions and level locations of any delamination as long as it is not hidden by another separation located closer to the exposed surface. The full pattern of flaws and damages may be found by C-scan detection applied from the two surfaces together with other NDT methods (radiography etc.). Checking of NDT results may be accomplished by microscopical analysis of sections cut from the damaged laminate (fractography). In the case of narrow laboratory specimens in which one dimensional damages have been formed, the dimension, location and growth of interlaminar damages, under loading may be detected and followed-up by direct on-line video inspection. The chances of existing interlaminar damage (flaw) to propagate under loading depends on its geometry, loading mode and direction, and mainly on the interlaminar fracture toughness (IFT) characteristics of the composite laminate. IFT characterization and its dependence of composite material composition was dealt

widely in numerous publications. Conclusions from these findings point out that by proper design of laminate lay-up and composition (hybridization), a significant improvement of IFT may be achieved, which will in turn, have a positive effect on damage tolerance capacity of the material. Propagation of in-plane cracks in a multidirectional laminate is arrested by the existence of fibers which are misaligned to the crack direction. Interlaminar damage, on the other hand, may propagate almost uninterrupted through the quasi-homogeneous and isotropic interlaminar plane (similar to crack propagation in isotropic homogeneous media). Hence, delamination is a major concern in structural composite materials. Detailed discussions on damage tolerance and interlaminar damage are forwarded by Ishai et al. (1988). 3.7.2 Compressive Loading and Laminate Buckling

The most severe loading condition for a laminate with interlaminar damage (or flaw) is the compressive mode. In such a mode, the damaged laminate is inclined to fail due to sublaminate local buckling mechanism. Such a mechanism was widely studied analytically and investigated experimentally. It was found that the buckled sublaminate provide the driving energy for additional delamination propagation, which in turn enhances premature laminate compressive failure due to the weakening effect and eccentricity caused by the partial separation of the sub-laminate. In standard testing specifications proposed by NASA and some aircraft industries, uniaxial compressive loading after impact procedure is recommended. These loading procedures are carried out on relatively large scale specimens, and can be defined as structural tests.

3.8 References

These tests are expensive especially when considering the number of such large specimens which are required to obtain reliable data. In the future, a cyclic compressive loading test will be probably required in order to investigate the residual fatigue life of impacted specimens. Such a program using panel testing will necessitate substantial funds and hence will be limited to exclusive research centers. A new approach to simplify the post-impact residual characterization procedure is suggested here, based on preliminary tests. In this method, impact and post-impact loading are conducted on small beams, where damage formation and propagation mechanism and its effect on ultimate failure can be detected directly during loading procedure. Residual compressive fatigue test may also be accomplished using "beam specimens". The limitation and advantages of "beam test" vs. "panel test" for damage tolerance evaluation, in view of the different requirements for structural design, material engineering and quality assurance disciplines should be assessed. 3.7.3 Residual Properties

Engineering design considerations must take into account the fact that the composite structure may consist of preformed fabrication flaws and damages due to accidental lateral loads, which will effect the design allowables. Interlaminar flaws are not sensitive to tensile inplane loading. On the other hand under compressive inplane loading such damages (flaws) are likely to propagate and consequently residual compressive strength of damaged laminate is expected to be significantly lower than its virgin reference. Similarly, existing delaminations are expected to propagate under cyclic loading and hence fatigue life is liable to be reduced as compared with that of reference laminate. In most damage tol-

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erance investigations, residual compressive properties of laminated panel after impact is plotted vs. initial impact energy. Damage tolerance methodology for advanced composites is not fully mature primarily due to the complexity in progressive failure modes encountered in service. Except for the case of dominant cracks such as delamination, damage tolerance methodology for interacting failure modes are still continuing to evolve. Current approaches are generally semi-empirical in nature with considerable bias toward experimental data. One current approach is to assess residual properties and relate them to the damage size and propagation modes, which in turn can be related to the state and distribution of stresses within the laminate.

3.8 References Agarwal, B. D., Broutman, L. I (1980), Analysis and Performance of Fiber Composites. New York: John Wiley & Sons. Bellenger, V., Fontaine, E., Fleishmann, A., Saporito, X, Verdu, J. (1984), Polym. Degrad. Stab. 9(4), 195. Carlsson, L. A., Pipes, R. B. (1987), Experimental Characterization of Advanced Composite Materials. Englewood Cliffs, NJ: Prentice-Hall, Inc. Carpenter, J. F. (1977), Assessment of Composite Starting Materials: Physicochemical Quality Control of Prepregs, AIAA/ASME Symp. Aircraft Composites, San Diego, CA. Enns, J. B., Gillham, J. K. (1983), /. Appl. Polym. Sci. 28, 2567. Flynn, J. H., Wall, L. A. (1966), Res. Nat. Bur. Stds. 70A (6), 487. Garton, A. (1984), Polym. Prepr. (Am. Chem. Soc, Div. Polym. Chem.) 25 (2), 163. Gillham, J. K. (1976), Polymer Eng. Sci. 16, 353. Hagnauer, G. L. (1980), Polymer Composites 1, 81. Hagnauer, G. L. (1981), It Res. Dev. 23 (4), 123. Hagnauer, G. L., Dunn, D. A. (1980), in: Materials 1980,12 (12th Natl. SAMPE Technical Conf.), pp. 648-655. Hagnauer, G. L., Dunn, D. A. (1982), Ind. Ed. Chem. Prod. Res. Dev. 21, 68. Hagnauer, G. L., Dunn, D. A. (1984), in: Plastics in a World Economy, ANTEC '84, Society Plastics

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Eng. 42nd Annu. Technical Conf. and Exhibition, pp. 330-333. Ishai, O., Rotem, A., Lifshitz, J. (1988), Damage Tolerance Evaluation of Structural Composite Materials. Haifa: Technion R & D Foundation. Johnston, N. J. (Ed.) (1987), ASTM STP937. Philadelphia, PA: Am. Soc. Test. Mat. Koenig, J. L. (1983), Quality Control and Nondestructive Evaluation Techniques for Composites, Part II: Physicochemical Characterization Techniques. A State-of-the-Art Review, U.S. Army Aviation R & D Command, AVRADCOM TR 83-F-6. May, C. A., Hadad, D. W, Browning, C. E. (1979), Polymer Eng. Sci. 19, 545. Mestan, S. A., Morris, C. E. M. (1984), Macromol. Sci., Rev. Macromol. Chem. Phys. C24(l), 117. MIL-HDBK-17 (1987), Plastics for Aerospace Vehicles: Part 1 - Composite Materials for Aircraft and

Aerospace Applications, Vol. 1 - Guidelines. Watertown, MA: U.S. Army Mater. Lab. Newaz, G. M. (Ed.) (1989), ASTM STP1044. Philadelphia, PA: Am. Soc. Test. Mat. Schapery, R. A. (1968), /. Composite Materials 2 (3), 280. Tsai, S. W, Hahn, H. T. (1980), Introduction to Composite Materials. Lancaster, PA: Techmonic Publishing Co. Whitney, J. M., Daniel, I. M., Pipes, R. B. (1982), Experimental Mechanics of Fiber Reinforced Composite Materials. Brookfield Center, CT: Soc. Experimental Mechanics. Zukas, W X., MacKnight, W. J. Schneider, N. S. (1983), in: ACS Symp. Ser. 227 (Chemorheology of Thermosetting Polymers). Washington, D.C.: Am. Chem. Soc, pp. 223-250.

4 Metal Matrix Composites Krishan K. Chawla Department of Materials and Metallurgical Engineering, New Mexico Institute of Mining and Technology, Socorro, NM, U.S.A.

List of 4.1 4.1.1 4.1.2 4.2 4.2.1 4.2.2 4.2.3 4.2.4 4.3 4.3.1 4.3.2 4.3.3 4.3.3.1 4.3.3.2 4.3.4 4.3.5 4.3.5.1 4.3.5.2 4.3.5.3 4.3.5.4 4.4 4.4.1 4.4.2 4.4.3 4.4.4 4.4.5 4.4.6 4.4.7 4.4.8 4.4.9 4.5 4.6

Symbols and Abbreviations Introduction Types of Metal Matrix Composites Important Metallic Matrix Materials Processing Liquid State Processes Solid State Processes Deposition Techniques In Situ Processes Interfaces in Metal Matrix Composites Importance of the Interface in Composites Crystallographic Nature of the Fiber/Matrix Interfaces Interfacial Bonding in Metal Matrix Composites Mechanical Bonding Chemical Bonding Thermal Stresses in Metal Matrix Composites Microstructure of the Interface Region Alumina/Metal Composites Silicon Carbide/Metal Composites Carbon Fiber/Metal Matrix Composites Interaction Between Interfacial Phenomena Properties Importance of the Metallic Matrix Modulus Strength Matrix Aging Response in Composites Toughness Thermal Properties Fatigue Creep Tribological Properties Applications References

Materials Science and Technology Copyright © WILEY-VCH Verlag GmbH & Co KGaA. All rights reserved.

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4 Metal Matrix Composites

List of Symbols and Abbreviations a a A b d D E E Eo Et

f g G

h IA Jc k K

AK AK th

Klc (1)

I m m n N Nc P P

P P r R R R{ Ro (s) s S S,Sf

t

T

crack length radius of fiber fatigue crack propagation parameter radius of matrix shell diameter of fiber, density diffusivity elastic constants (general) Young's modulus modulus of uncracked material transverse Young's modulus subscript: fiber acceleration due to gravity shear modulus height of composite area of fiber/matrix interface critical current density (of superconductors) Boltzmann constant bulk modulus cyclic stress intensity factor threshold stress intensity plain strain fracture toughness liquid length of composite, length of shaft fatigue crack propagation exponent subscript: matrix creep stress exponent number (of cycles) critical speed (rpm) pressure between interface and matrix property subscript: particle load geometric parameter solidification rate stress ratio inner radius outer radius solid aspect ratio of discontinuous reinforcement stress amplitude precipitates in A l - C u - M g system time temperature

List of Symbols and Abbreviations

AT Tc Tm V w w x Y

temperature interval critical temperature (of superconductors) melting temperature volume fraction subscript: whisker width of composite thickness of the reaction zone a constant, a geometric factor

a Aa s e 0'90" X v a o Aa G0 cra